Article www.acsnano.org Epitaxially Driven Phase Selectivity of Sn in Hybrid Quantum Nanowires Sabbir A. Khan,* Sara Martí-Sánchez,* Dags Olsteins, Charalampos Lampadaris, Damon James Carrad, Yu Liu, Judith Quiñones, Maria Chiara Spadaro, Thomas Sand Jespersen, Peter Krogstrup,* and Jordi Arbiol* Cite This: ACS Nano 2023, 17, 11794−11804 Read Online ACCESS Metrics & More Article Recommendations * sı Supporting Information ABSTRACT: Hybrid semiconductor−superconductor nano- wires constitute a pervasive platform for studying gate-tunable superconductivity and the emergence of topological behavior. Their low dimensionality and crystal structure flexibility facilitate unique heterostructure growth and efficient material optimization, crucial prerequisites for accurately constructing complex multicomponent quantum materials. Here, we present an extensive study of Sn growth on InSb, InAsSb, and InAs nanowires and demonstrate how the crystal structure of the nanowires drives the formation of either semimetallic α-Sn or superconducting β-Sn. For InAs nanowires, we observe phase- pure superconducting β-Sn shells. However, for InSb and InAsSb nanowires, an initial epitaxial α-Sn phase evolves into a polycrystalline shell of coexisting α and β phases, where the β/ α volume ratio increases with Sn shell thickness. Whether these nanowires exhibit superconductivity or not critically relies on the β-Sn content. Therefore, this work provides key insights into Sn phases on a variety of semiconductors with consequences for the yield of superconducting hybrids suitable for generating topological systems. KEYWORDS: nanowires, topological materials, semiconductor-superconductor hybrid, Sn, quantum computing, interface, epitaxy T in is a common group IV element used in a broad range of industrial applications including electronic circuits, optoelectronic devices, energy storage devices, and coating for commercial products.1−7 Recently, interest was further triggered with the promise of exploring exotic mixed materials present problems with regard to fabrication. While research has been undertaken on crystal phase stability in bulk Sn,6,16,17,19,21,39−46 a comprehensive study of nanoscale Sn phase formation on semiconductor nanowires (NWs) is yet to be performed. Such knowledge is crucial to determining topological phases with Sn−semiconductor hybrids.8−11 selective phase growth and controlling the interface properties However, the type of topology depends on which of the two of the Sn-based hybrid quantum heterostructures. major allotropes (α-Sn or β-Sn) is present. Hybridizing the In this paper, we perform Sn thin-film growth on different semimetallic, cubic α-Sn phase,6,12−18,20 with semiconductors semiconductor NWs and present an in-depth analysis. We tune such as InSb (111) or (100) induces lattice mismatch-related the growth parameters to determine the optimal conditions strain fields. The resulting broken cubic symmetry can lead to and investigate the associated structural changes in the Sn film. topological insulator behavior.8,9,14,21−26 Conversely, metallic When grown on B-polar cubic InSb and InAsSb NWs,11,47,48 β-Sn is a comparatively dense body-centered tetragonal Sn grows predominantly in the lattice-matched α phase. With structure and exhibits superconductivity with a bulk critical increasing film thickness, we observe an increasing density of temperature of 3.7 K.6,12−18,20 Hence, hybrids of β-Sn and one- embedded β-phase grains. Based on the comprehensive crystal dimensional semiconductors with strong spin−orbit interac- tion, such as InSb or InAs, may exhibit topological super- conductivity in the presence of magnetic fields.27−38 Received: March 25, 2023 Unleashing the potential of either of the two Sn phases� Accepted: June 8, 2023 and excluding the potential for unwanted behaviors�requires Published: June 15, 2023 phase-pure crystal growth, which is a challenge in thin films.17,39−41 In particular, the presence of α-Sn significantly reduces the yield of superconducting devices,17 and phase- © 2023 The Authors. Published by American Chemical Society https://doi.org/10.1021/acsnano.3c02733 11794 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article Figure 1. Hybrid InSb−Sn nanowires with multiple Sn phases. (a) Scanning electron microscopy (SEM) image of an MBE-grown InSb nanowire in situ hybridized with a Sn shell. (b) Cross-sectional aberration-corrected HAADF-STEM micrograph of the InSb−Sn interface. (c) Magnified HAADF-STEM image from the InSb segment showing cubic zinc blende crystal structure. (d) Low-magnification longitudinal HAADF-STEM micrograph of the heterostructure where coexistence of both α-Sn and β-Sn is observed. Dashed line indicates the starting of the Sn shell. (e) Magnified image of the interface between cubic-α and tetragonal-β phases. The white dashed line marks a {111}α-Sn plane interfacing with a β-Sn domain. (f) Zoom-in of β-Sn. (g) Fourier filtered map of crystal planes from the squared region marked in panel (d). Blue section matches the β-Sn phase, while red corresponds to InSb and α phases, which could not be unambiguously decoupled due to the equivalent cubic symmetries and their lattice mismatch being close to 0%. (h) Fast Fourier transform employed for (g) structural map showing indistinguishable plane reflections for InSb and α-Sn (red circles), while blue circles correspond to {200} planes of [001]-oriented β-Sn. Scale bars are (a) 100 nm, (b) 5 nm, (c) 1 nm, (d) 20 nm, (e) 10 nm, and (f) 1 nm. analysis of the growth evolution, we suggest a nucleation scanning electron microscopy (SEM) image of an InSb NW mechanism for β-Sn embedded within an α-Sn matrix. hybridized with an approximately 20 nm Sn shell. Here, Sn was Obtaining hybrids with phase-pure β-Sn was possible by grown on the selected facets of the hexagonal NW as indicated using InAs semiconductor NWs. Here, the β-Sn creates crystal in the image. The crystal parameters and the elemental domain matched grains, while differences in crystal symmetry distribution mapping of these NWs are shown in Supporting between InAs and α-Sn and the large associated plane Information, S1 and S2. mismatches strongly suppress α-Sn nucleation. Finally, we For crystal analysis, aberration-corrected high angular correlate structural properties with electrical behavior at low annular dark field (HAADF)−scanning transmission electron temperatures. We show a strong correlation between the microscopy (STEM) micrographs were acquired in a cross- presence of superconductivity in ZB (InSb/InAsSb)-Sn with section of the hybrid NWs (Figure 1(b)). Clear contrast the β grain density and that those films expected to have a between the InSb core (dark) and Sn shell (bright) can be vanishingly low β-Sn content exhibit no trace of super- observed. The InSb segment can be identified as a cubic zinc conductivity. Conversely, β-dominant WZ(InAs)−Sn-based blende (ZB) crystal with a lattice constant of 6.479 Å (see hybrids exhibit an induced superconducting gap with energy Δ magnified HAADF-STEM micrograph in Figure 1(c)). In = 440 μeV. Figure 1(b), Sn seems to form a single crystal, but unambiguous phase identification was not possible due to RESULTS AND DISCUSSION either the α−β overlap or elemental migration and damage Multiphase Sn Crystals in Hybrid NWs. Au catalyst- induced during focused ion beam (FIB)-assisted cross-section assisted NWs were grown on InAs (111)B or InAs (100)11 preparation (see Supporting Information S3). Therefore, we substrates using a molecular beam epitaxy (MBE) system. performed structural analysis on longitudinal views of NWs Similar to our earlier works,11,49−51 InSb and ternary InAsSb directly transferred to Cu grids by employing a combination of NWs were grown from InAs stems. Once the NWs were high-resolution TEM (HRTEM) and aberration-corrected grown, Sn deposition was performed in a high-vacuum physical HAADF-STEM imaging (see Figure 1(d)). Here three vapor deposition (PVD) chamber, which is in vacuo connected different HAADF intensities can be observed: the darker to the MBE (see Methods for details). Figure 1(a) shows a contrast originated from the InSb core, and the deposited Sn 11795 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article Figure 2. Sn shell evolution and β-Sn grain formation. (a) HRTEM micrograph of the InSb-Sn NW, where approximately 3 nm Sn was deposited. Due to the ultrathin film and epitaxial relation, the interface is just distinguished by a subtle contrast variation. FFT on a magnified image of the interface confirms the α-Sn phase. (b) HRTEM image of the approximately 5 nm Sn shell where an initial stage twin plane is identified. (c) Large-scale low-magnification HAADF-STEM micrograph of a 20 nm Sn shell where clear contrast between α-Sn and β-Sn can be observed. Zoom-in shows β-Sn grains with well-defined faceting that seem to arise from {111} faceting of the α-Sn shell (white dot lines). The bottom panel shows an atomically resolved HAADF-STEM micrograph of the α−β interface. (d) HAADF-STEM micrograph of 40 nm Sn deposited on InSb. β-Sn grains can be observed on top of the initial α-Sn. (e) Hypothesis for the α, β nucleation of the Sn film on the InSb NW. (i) Initially α-Sn grows epitaxially on the lattice-matched InSb cubic interface. (ii) A few layers away from the interface twin defects appear (indicated by the arrow) mostly oblique to the growth direction. (iii) These twin planes act as preferential nucleation centers for the β-Sn phases when the activation energy barrier is overcome and, hence, a β grain forms. (iv) For thick shells, large β-Sn grains with different orientation can be observed and initial {020}-β faceting dissolves. (f) Schematic of the α ↔ β transition as a function of chemical potential (μ) and temperature. Here, Tg (during growth) and T are the holder temperatures. Scale bars are (a) 10 nm and 2 nm, (b) 10 nm and 2 nm, (c) left: 50 nm, right: 20 nm (both), bottom: 2 nm, (d) top: 5 nm, bottom: 1 nm (both). shows the light contrast with an enclosed area in the shell. epitaxial cubic α-Sn shell. The β grains typically show well- Taking a closer look at Figure 1(d), we can observe that most defined {020} faceting and seem to arise from {111} faceting of the shell matches the α-Sn phase, which is also face-centered on the α shell (indicated with white lines in Figure 1(e)). An cubic with a lattice constant of 6.489 Å. The small lattice atomic-resolution [001]-oriented β phase arrangement and its mismatch (ϵ ∼ 0.15%) between ZB-InSb and α-Sn favors its diffraction pattern are shown in Figure 1(f). Further, in Figure growth on the NW interface, which we will elaborate on in a 1(g) we present a Fourier plane filtered phase map from the later discussion. The brightest small domain can be identified atomically resolved micrograph of the hybrid NW to highlight as a β-Sn grain embedded in the α-Sn matrix. From this and the multiphase structure. As we can see, due to the quasi other analyzed NWs (see Supporting Information S4), we identical atomic arrangement and atomic number, the InSb observed that β grains are mostly not grown directly interfacing core and α-Sn have a perfect epitaxy and are barely with the InSb NW core, but they seem to appear on the distinguishable. On the other hand, due to the different crystal 11796 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article structure and reduced unit cell volume, the β-Sn grain shows α can be observed, which belongs to the {020− β family of distinct contrast and exhibits a sharp interface with the other planes (see more examples in Supporting Information S4). At cubic structures. These differences in the lattice arrangement the bottom of Figure 2(c), we present an atomically resolved can be noticed in the fast Fourier transform (FFT) displayed in micrograph of the α−β interface between cubic and tetragonal Figure 1(h). Sn phases. Finally, an approximately 40 nm deposited Sn shell Optimal Conditions for the Sn Shell on NWs. Thin-film is shown in Figure 2(d). The shell exhibits a continuous but formation of Sn crucially depends on the substrate temperature rough morphology (also see Supporting Information S12). and chemical potential of the incorporation sites, as they Here all the NWs investigated by TEM showed a similar exponentially govern the adatom diffusion length on the NW nanostructure, with multiple β-Sn grains being randomly facets.11,52,53 Low-temperature growth assists in minimizing oriented (most of them out-of-axis), and they are grown from the adatom diffusion length and drives the Sn film to wet on the α phase. Figure 2(d) shows an example of two neighboring the NW surface. Besides, the chemical potential of the β-Sn grains oriented in two different zone axes, highlighting the incorporation sites depends on the change in Gibbs free polycrystalline nature of the β-Sn shell in contrast to the fully energy that is determined by the interface energy density epitaxial α-Sn. It is also notable that the strong {020}-β between NW facets and deposited Sn.11,52,53 In the initial stage faceting vanishes as the β grains grow. We also observe an of Sn growth, the interface energy depends on the lattice increasing density of β-Sn grains presenting different mismatch between two materials and the area of the interface. orientations as we increase the shell thickness. Therefore, In the latter stage, strain energy also plays a key role in defining from a fully epitaxial α-Sn shell promoted by lattice-matched the interface energy density and evolving grain boundaries. InSb templates, we observe an evolution to form a polycrystal- Hence, selection of NW and a Sn phase with low residual line β-Sn shell in thick films. mismatch will ease the formulation of a continuous Sn shell. In Based on the experimental analysis, in Figure 2(e), we contrast, if the NW and Sn phases have a high residual present a hypothetical model of the structural evolution of Sn mismatch, then a significant thermodynamic force will require phases in the shell grown on a cubic NW template. The to make a continuous film on NW facets. Further, a high flux chemical potential of the α-Sn or β-Sn phase on the NW can rate also assists Sn shell growth by increasing the adatom be associated with the chemical potential of the reference bulk concentration on the NW facets. High adatom concentration phase (μi,bulk) and thin film (μi,ex). The chemical potential in increases the nucleation density on the given surface, which the thin film will again depend on the NW surface (μi,surf), also decreases the average adatom diffusion length. In a interface (μi,int), strain (μi,strain), and grain boundary (μi,gb). nutshell, a combination of the low substrate temperature, Hence, α or β crystal formation is an accumulated factor of all favorable interface energy density, and high flux rate can those parameters: provide an optimized Sn shell on the NW facets. Hence, we i = i ,bulk + i ,ex maintained the Sn flux rate in the range of 3−5 Å/s, and the sample holder temperature was either approximately −100 or = + + + i ,ex i ,surf i ,int i ,strain i ,gb −150 °C during the growth. A detailed study of growth temperature effect on the hybrid Sn shell structure is presented where i = α, β. in the Supporting Information, S5−S9. Since bulk α-Sn is stable below 13.2 °C and the interface β-Sn Nucleation in an α-Sn Matrix. To understand the energy between InSb and α-Sn is very small due to its close thin-film evolution, we performed an extensive structural study lattice matching, in our low-temperature growth condition, it is varying the Sn shell thickness (Figure 2(a−d)). In Figure 2(a), favorable to form α-Sn in the interface. Hence, the epitaxial Sn approximately 3 nm of Sn was deposited on an InSb NW. As film appears initially (Figure 2(e)(i)). In the later stage (Figure α-Sn forms an epitaxy with InSb, it was difficult to differentiate 2(e)(ii)), twin formation in the α-Sn shell initiates mostly in the interface. However, a weak contrast fluctuation from the the {11−1} plane away from the interface. Such twins are interface (indicated with the white dot line) can be observed. oblique to the growth directions and can be found all over the We did not see any tetragonal β-Sn domain in this extremely shell, as shown in the Supporting Information S5−S9. thin shell, which would have been noticeable on the cubic Although epitaxial α-Sn is grown at the interface, it is structure. More details can be found in Supporting Information metastable in temperature exposure and can change the S10. In Figure 2(b), approximately 5 nm of Sn was deposited crystallinity if the chemical potential of α-Sn overcomes the on the InSb NWs. Here, a dominance of α-Sn is still observed, kinetic energy barrier between α and β phases. The schematic and some twinning planes start to be visible, as indicated in the of activation energy dependent α ↔ β transition is shown in magnified section and inset diffraction pattern. In addition, a Figure 2(f) and can be expressed as follows: few seed-stage β-Sn grain domains are also observed in this case (see Supporting Information S11). In some sporadic = cases, we observed β-Sn grains close to the interface with InSb, ji E zy but unambiguous resolution of possible direct contact between expjjj zz j kBT zz{ ,Sn InSb and β-Sn could not be determined due to the 3D k structure of the system. In any case, the threshold thickness for which we start to observe β-Sn nucleation is 5 nm; therefore, Here, ΔΓαβ is the growth rate and δEαβ is the activation energy, we assume some epitaxial α-Sn might be present prior to it which is dependent on μi. Next, tetragonal β grains nucleate on even though it might not be the only β-Sn formation top of the cubic α-Sn shell, overcoming the activation energy mechanism. In Figure 2(c), approximately 20 nm of Sn was barrier. This activation energy barrier may be overcome due to deposited, and full-fledged faceted [001]-oriented β-Sn the change in local conditions during the growth, causing a growing on the {11−1}α-Sn planes can be observed. In the transition from the α to β phase and vice versa. However, a magnified inset images, a clear preferential faceting of β-Sn on likely possibility is that β nucleation occurs while the sample is 11797 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article abruptly exposed to a temperature that is higher than the growth temperature (see Figure 2(f)) (perhaps right after the low-T Sn growth). As we observed, these β-Sn grains typically display well-defined faceting that seems to arise from {111} faceting on the α-Sn (Figure 2(e)(iii)). Hence, it can be inferred that β-Sn overgrows from an α structure by nucleating in twin-induced rough sites instead of a flat surface of InSb NW. Finally, for a thick Sn film (Figure 2(e)(iv)), β-Sn grains gain volume and often multiple grains coalesce into one large grain; simultaneously more β grains start to nucleate in different places of the shell. As a result, different orientations of β-Sn grains can be observed. Once both crystal phases are stabilized at room temperature, no further transition between α-Sn and β-Sn happens, as the activation barrier for the transition is very high at this point. Hence, both α and β crystal phases coexist in the Sn shell on cubic NWs. In Supporting Information S13 we performed thermal annealing on the 20 nm shell NWs and observed no changes in the Sn heterostructure even when turning the conditions to completely favor the β-Sn phase, proving the structural stability of the achieved configuration. Increasing Interfacial Lattice Mismatch. Based on the above discussion, it can be presumed that the lattice parameter of the NW core is key to control the Sn-shell phase formation in the hybrid. Therefore, other NW compositions were employed to decrease the lattice match of α-Sn and, hence, hinder its formation. With this purpose, we grew NWs with 80% Sb and 20% As composition, which implies a lower lattice constant than in pristine InSb NWs given Vegard’s law,54 while preserving the ZB cubic phase. This configuration implies an increase in lattice mismatch between cubic α-Sn and the NWs, rising from 0.15% with pure InSb to 1.5% for InAs0.2Sb0.8. Figure 3(a) presents InAsSb NWs with a 20 nm Sn shell grown on selected facets of the NWs, where the growth rate was approximately 3 Å/s and the sample holder temperature was approximately −150 °C. Magnified sections of these NWs are shown in panel (b) (bottom part) and in panel (c) (top part). Two major observations from these hybrids are, first, the Sn shell exhibits a continuous but inhomogeneous morphology on InAsSb NWs compared to InSb NWs; second, the hybrid InAsSb−Sn NWs show strong bending along the deposition Figure 3. Morphology and structural analysis of InAsSb/Sn NWs. direction. Figure 3(d) presents a low-magnified HAADF- (a) Tilted SEM image of the InAsSb/Sn NWs. Here 20 nm Sn is grown on the three facets of InAsSb NWs. The white arrow shows STEM image and an electron energy loss spectroscopy (EELS) the deposition direction of the Sn. (b) Magnified bottom section of elemental map of In/Sn in an NW, which confirms the the NW from panel (a). Continuous but rough morphology of the continuity of the shell. Taking a close look into the crystal Sn shell can be observed. (c) Magnified top section of the NW arrangement of the shell, Figure 3(e) shows an example of the from panel (a). (d) HAADF-STEM image of InAsSb (left) and most commonly observed crystal phase of Sn on the InAsSb EELS-based In/Sn compositional map, which reveals continuous NWs, belonging to the α-Sn phase and forming a similar Sn coverage. (e) HRTEM image and its FFT of α-Sn on InAsSb. structural configuration to that observed in InSb NWs. As The interface is barely distinguishable due to the same crystal previously discussed, the α-Sn shell contains several planar symmetry (see inset model) between cubic α-Sn and ZB InAsSb, defects including twinning and stacking faults, which form at but twinning is present in Sn. (f) HRTEM image and its FFT of observed β-Sn grain on α-Sn on InAsSb. Scale bars are (a) 1 μm, ∼71° with the growth direction plane. The model in the inset (b−d) 100 nm, and (e, f) 20 nm. of Figure 3(e) shows the atomic arrangement between α-Sn and InAs0.2Sb0.8. In addition, we could observe some β-Sn grains embedded in the α-Sn, also exhibiting the same crystal energy results in compressing strain fields that relax toward the configuration with well-defined faceting in the {200} planes, outer Sn shell and thus strongly bend the NW (Figure 3(a)). A but they are discretely distributed through the NW shell, which schematic of compressive strain fields originating the bending displays a prevalence of the α phase. More micrographs on is displayed in Supporting Information S14. As a consequence, different areas are presented in the Supporting Information this configuration still favors α-Sn nucleation over β-Sn. In (S14 and S15). addition, this bending can be responsible for the rough Hence, despite increasing the interfacial lattice mismatch of morphology observed in the shell. Although the Sn crystal the NW with the cubic Sn phase, we still obtain a preference structure on InAsSb is identical to the InSb one, careful for α-Sn nucleation. We presume that the increase in interface attention should be put on the resulting electronic behavior of 11798 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article the InAsSb−Sn system because of the strain accumulation at lattice-matched configuration for either of the two Sn phases, the interface, as it is capable of locally modifying its leading to a high interface energy between these materials (see properties.55 Supporting Information S1). 11,53 As a result, a high Selective β-Sn Nucleation in a Poor Lattice Matched thermodynamic driving force is needed for Sn to wet on the WZ Core. Finally, we investigate the deposition of Sn on the InAs NW facets during the growth. Hence, in our growth WZ InAs NWs. Figure 4(a) shows InAs NWs with a ∼20 nm conditions with a limited low-temperature range, Sn remains as dewetted islands on the InAs NW. Therefore, to achieve a Sn shell grown at the rate of 3 Å/s and a holder temperature of continuous film with the same thickness, the temperature approximately −150 °C. Sn forms discrete grains with uniform should be further lowered so that the Sn adatoms can be thickness on InAs NW facets. The NW crystallizing in the immobilized during the growth. Another way to achieve a hexagonal phase implies that there is no particularly favorable continuous film is by increasing the Sn thickness. Figure 4(b) shows an example of ∼45 nm Sn deposited on an InAs NW keeping the same growth conditions. Doubling the Sn thickness enhanced the length of the islands. However, the shell remains discrete and more material should be deposited to achieve a continuous film. As presented in the schematics on the left, upon reaching a critical thickness, the neighboring islands coalesce and evolve to form large islands, which can merge again with the adjacent ones. Figure 4 (c) shows an EELS-based compositional map of Sn and As of an NW together with a HAADF-STEM micrograph of the area of study. We can observe that despite the grains tending to be relatively flat and homogeneous in terms of thickness, the grain size distribution is highly inhomogeneous. When moving to analyze the crystal phase of Sn, we found most of the grains, including the larger ones, are out-of-axis with respect to the NW (Supporting Information S16), while we could find some small sections being in-axis with the InAs (see Figure 4(d−f)). We can identify all these grains as β-Sn (Figure 4(f), and they are in direct contact with the InAs NW (Figure 4(e)), with no presence of α-Sn prior to their nucleation. More micrographs with β phase identification are available in the Supporting Information (S16−S20). Based on the analyzed structures, we presume that the presence of a hexagonal crystal as growth template inhibits the formation of the cubic α-Sn phase. In the previous cases, the same cubic crystal nature of InSb and InAsSb acted as a template for cubic α-Sn growth and made it the most stable phase by hindering the nucleation of tetragonal β-Sn. However, when employing hexagonal InAs as a core material, the crystal symmetry equivalence is broken. The most similar configuration between cubic and hexagonal would be InAs(0002)/α-Sn(111),56,57 but this still implies a 7% plane mismatch along the NW growth axis direction. Therefore, lower mismatch β-Sn growth gets favored instead of α-Sn, and it is the only phase we observed in all the NWs we examined. While in previous cases β-Sn grains nucleated on the α-Sn matrix showed some preferential out-of-plane orientation, β-Sn grains on InAs show a high degree of randomness regarding their orientation with respect to the InAs lattice, especially for Figure 4. Sn shell growth on wurtzite InAs NWs. (a) Tilted SEM smaller grains. However, for the long-range grains, we observe image of the InAs NW with a Sn shell. Sn forms discrete grains a dominant Sn orientation in the in-plane direction, although it with uniform thickness (∼22 nm), but nonuniform length (>100 is not in axis with the InAs crystal. In Supporting Information nm). Arrow shows the deposition direction. (b) Doubling the Sn S16−S18 multiple HRTEM micrographs of the InAs-Sn deposition time, where the Sn shell thickness is ∼45 nm. With heterostructure with larger Sn grains are presented. A 2.9 Å longer growth time nearby, Sn grains grow larger and eventually plane reflection parallel to InAs(0002) is visible in all of them collapse on each other as shown in the right-side schematic. (c) and creates periodic moiré patterns in areas overlapping with HAADF-STEM micrograph of an InAs-Sn NW and EELS InAs. This reflection belongs to the {020}β-Sn family of compositional mapping of As and Sn on the squared region. (d) planes. Given the differences in lattice symmetry between β-Sn HRTEM micrograph showing two Sn grains in direct contact with the InAs NW. (e) Fourier filtered structural map of (d), showing and WZ InAs, there is no low-index 3D configuration with low the spatial distribution of two differently oriented β-Sn grains on associated mismatch in all matching planes, and this explains InAs. (f) Zoom-in of the squared regions in (d) with their the appearance of multiple crystal orientations in the corresponding indexed FFTs and phase identification as β-Sn. heterostructures. In the case of the observed InAs(0002)/β- Scale bars are (a−c) 100 nm, (d) 10 nm, and (f) 1 nm (both). Sn{200}, they provide a mismatch of εInAs(0002)/β-Sn{200} 11799 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article = 16%. However, the domain formation consisting of 5×(002)InAs with 6×{200} β-Sn planes reduces the mismatch to ε5×InAs(0002)/6×β-Sn{200} = 0.017%. In fact, the moiré periodicity found in the overlapping areas matches with this plane combination (S16−S18). We assume that this low plane mismatch domain is therefore more energetically favorable than the other observed orientations and can be maintained in a longer spatial range, thus being the orientation observed in larger Sn grains. More details and schematics of the plane matching in the formed domains can be found in S16. Last but not least, in order to cross-check that the observed preclusion of α-Sn growth on InAs NWs is caused by its structural symmetry rather than its composition, we radially grew an approximately 2 nm shell of InAs covering the InSb NW before depositing Sn (see Supporting Information S21). In these conditions, InAs achieves ZB phase and remains completely strained on top of the InSb, adopting its crystal symmetry and lattice constant. We observed that regardless of the presence of InAs covering InSb, Sn crystallizes mainly in the α phase, with some β grains being embedded on it (see S21). This fact is consistent with the NW crystal structure being one of the key parameters driving Sn phases on hybrid NWs. A schematic of Sn phase dynamics as a function of lattice parameters is presented in Supporting Information S22. Low-Temperature Electrical Measurements. To inves- tigate the superconducting properties of the hybrids, we fabricated and measured electrical devices at low temperatures. Individual hybrid NWs were transferred from the growth substrate to a substrate of highly doped silicon capped with 200 nm of SiO2. Nanowires were located with respect to predefined alignment grids, and Au/Ti contacts were defined using standard e-beam lithography and metal evaporation. See Methods section for details of the fabrication. Figure 5(a) shows a SEM micrograph of a typical device with four contacts on the same NW. A 10 nA ac current was sourced between the outer contacts, and the resulting voltage drop between the inner pair was measured using standard lock-in techniques, yielding the resistance between the inner pairs. Measurements Figure 5. Low-temperature measurements of NW−Sn hybrid were performed in a dilution cryostat with a base temperature devices. (a) SEM micrograph of a typical device. (b) Temperature of 15 mK. Figure 5 shows the temperature dependence of the dependence of the four-terminal differential resistance for nine resistance normalized to the value at the highest temperature different devices from three different hybrid types. See the text and for nine different NWs from six different growths of three Supporting Information S23 for details. For each device the different hybrid types. D2−D4 and D7−D9 are InSb-Sn resistance is normalized to the value at the highest temperature and curves are offset for clarity. (c, d) V−I curves for device D9 devices, D5 and D6 are InAsSb-Sn devices, and D1 is an measured as a function of temperature and perpendicular magnetic InSb(InAs shell)−Sn device. See the table in Supporting field, respectively. The sweep direction was from negative to Information S23 for the growth parameter details. Devices D3 positive current. (e) SEM micrograph of a two-terminal InAs−Sn and D1 show no changes in the resistance upon cooling, which N−S device. (f) Differential conductance as a function of voltage we attribute to the absence of β-Sn grains in the gap between bias measured for the InAs−Sn device in different temperatures. the inner voltage probes for these devices. All other devices, Scale bars are (a) 1 μm and (e) 200 nm. however, show a clear drop in resistance around 3.7−3.9 K close to the expected transition temperature of β-Sn. While D9 including proximity coupling neighboring grains through the shows an abrupt drop to an R = 0 Ω at ∼3.8 K, the remaining normal matrix or semiconductor wire. A similar trend was devices exhibit a gradual transition with multiple distinguish- shown in systems of designed Josephson arrays of super- able steps upon cooling. This behavior is consistent with a Sn conducting regions in a normal matrix.59 These results are shell composed of discrete β-Sn grains electrically coupled consistent with the structural analysis presented in Figures 1−3 through a non-superconducting α-Sn matrix and the semi- for ZB InSb and InAsSb NWs based hybrids. Figure 5(c,d) conductor NW. At TβC, isolated β-Sn grains turn super- show the V−I characteristics of D9 as a function of conducting, and while the resistance drops, the resistance temperature and perpendicular field, respectively. The remains finite. Other grains have transition temperatures measurements were performed sweeping the current from suppressed either due to finite size effects58 or because of an negative to positive, and the transition from the resistive to the inverse proximity effect from the surrounding matrix. Upon superconducting state upon sweeping I from a large value lowering the temperature, a gradual transition thus occurs until toward zero occurs at a smaller absolute current than the a global superconducting phase is established possibly transition from the superconducting to the dissipative state 11800 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article observed when increasing I from zero. We attribute this performance of NW−Sn-based topological and superconduct- asymmetry to Joule heating in the normal state. The critical ing quantum devices. current can be observed up to 3.5 K (for B⊥ = 0) and B⊥ = 650 mT (for T = 15 mK), consistent with panel (b). METHODS Finally, to measure the size of the superconducting gap and Hybrid Crystal Growth. Molecular Beam Epitaxy. A Vecco Gen confirm that superconductivity extends from the Sn shell into II molecular beam epitaxy device connected to a high-vacuum the semiconductor NW, we further fabricated a two-terminal electron gun assisted physical vapor deposition system was used to device on the InAs−Sn stem, as shown in Figure 5(e). Here a grow NW−Sn hybrids. III−V NWs were grown in the MBE chamber; large isolated grain could be seen with one contact positioned later the samples were transferred into the PVD chamber through the on the grain and the other on the adjacent Sn-free high-vacuum tunnel to grow Sn on the selected facets. All of the NW semiconductor segment. As explained in Figure 4, the discrete batches were grown on an InAs (111)B single-side polished substrate. Sn grain here is expected to be only the superconducting β-Sn Au was used as growth catalyst and deposited in situ in the MBE phase without the α mix. The electron density in the chamber before NW growth. In the MBE, material cells (III−V and Au) were preheated and stabilized before the growth. Optimized cell semiconductor can be tuned using the potential on a nearby parameters for NW growth are In, 865/815 °C, As, 400/345 °C, Sb, electrostatic side-gate, and close to pinch-off a quantum dot 680/500 °C, and Au, 1100 °C. While loading the InAs (111)B (QD) forms in the semiconductor (see Supporting Informa- substrate, standard 200 °C for 2 h of baking at a load-lock chamber tion S24 for further characterization). In Coulomb blockade followed by 250 °C for 1 h degassing at the buffer tunnel was carried the QD acts as a tunnel barrier allowing tunnel spectroscopy of out, confirming a clean InAs surface for the growth. Subsequently, the the density of states in the semiconductor. Figure 5(f) shows samples were transferred to the MBE growth chamber, where the measured differential conductance dI/dVSD vs VSD for annealing was performed for 6 min at a 610 °C substrate temperature different temperatures. Here VSD is the applied voltage bias. under high As4 overpressure. Next, the substrate was cooled down to Superconducting coherence peaks are observed at VSD = ±Δ*/ 590 °C and Au was deposited for 1 s. After Au deposition, the substrates were further lowered to growth temperature, which is 447 e, corresponding to an induced gap of Δ* = 440 μeV, close to °C, and the In flux was introduced in the growth chamber to nucleate the value expected for a gap induced by β-Sn. the Au catalysts. For InAs NWs a 35 min growth was performed. For InSb and InAsSb NWs, Sb flux was introduced after the InAs stem CONCLUSION section. Usually the stem was grown for 6 min. For InSb, the As flux was turned off after the stem, whereas for InAsSb, As flux was In summary, we have presented hybrid NWs using various III- compensated according to the desired Sb:As ratio of the NW. The V semiconductors and Sn shells and performed an in-depth substrate was cooled to 150 °C after the NW growth and transferred analysis of Sn crystal phase formation in the heterostructure. to the PVD chamber for hybridizing with Sn. InSb, InAsSb, and InAs NWs were grown using MBE and later Electron-Beam Physical Vapor Deposition. For low-temperature in situ hybridized with Sn by determining optimal conditions growth of the Sn thin film, substrate cooling was carried out with a and attaining the utmost interface quality. We have provided a liquid nitrogen supply. The lowest holder temperature reach of our comprehensive structural analysis at the atomic scale that system is approximately ∼−150 °C, which is recorded with the thermocouple connected at the back of the substrate holder. Hence, reveals the coexistence of the two different Sn crystal phases we point out that the real temperature on the sample surface may (α-Sn and β-Sn) when deposited on cubic crystal NWs like slightly deviate from the reading. Once the desired growth InSb or InAsSb. Based on the experimental assessment in temperature was reached, we waited further to completely stabilize different stages of the Sn film, we proposed a hypothesis of α,β- the temperature before preparing the chamber for deposition. Sn phase nucleation in the hybrid. The NWs with a cubic Subsequently, the substrate was aligned with respect to the Sn lattice provide a template to epitaxially grow α-Sn due to their source, and the deposition angle was corrected to achieve selected quasi-identical crystal structure (space group, atomic positions, facets (two or three) growth. Details about NW facet selection growth and lattice constant). Later, β-Sn nucleates preferentially from and angle corrections are discussed in ref 11. Next, the source Sn materials were heated up with energetic electron beams, and the the defective sites of the film by overcoming an activation evaporated Sn flux rate was monitored through a quartz-crystal energy, and eventually isolated grains are observed to be microbalance (QCM). Once our desired flux rate was achieved, Sn embedded in the dominant α-Sn phase. The density of β-Sn was deposited on the NW facets. The deposition rate was grains has been shown to increase with the shell thickness. On experimented from 0.5 to 5 Å/s to find the optimized morphology. the contrary, selective growth of β-Sn was achieved on InAs After deposition, the hybrid NW sample was kept on the cold holder NWs, which is assumed to be a result of the poor lattice match for a while and then unloaded through the load-lock. During the between the WZ hexagonal crystal and the cubic α phase. In unloading process, the NW was initially exposed to oxygen for agreement with the structural analysis, low-temperature approximately 15 min and later vented with nitrogen before introducing to room temperature. electrical measurements on InSb−Sn and InAsSb−Sn devices Morphological and Structural Characterization. Scanning manifest the influence of discrete β-Sn grains on the Electron Microscopy. A JEOL JSM-7800F SEM was used to surrounding α-Sn matrix. While one device demonstrated a investigate the morphology of the hybrid NW. Standard operating sharp SU transition, others show a multistep transition, which conditions for these samples are a 15 kV accelerating voltage and a is possibly due to the inverse proximity effect from a 10−15 mm working distance with a secondary electron detector. conducting−nonconducting mixed structure. Additionally, by Depending on the imaging requirements and investigating materials, measuring a phase-pure β-Sn grain on an InAs NW, we we also used a 30° tilted viewing angle. Further, for a detailed surface demonstrate the superconducting size and hybridization of morphology investigation of the hybrid facets we used a 2−5 kV accelerating voltage with a 5 mm working distance, which helped us superconductivity to the semiconducting base. In brief, along gain information about the roughness and residual metals. with the optimal growth conditions, choosing the right lattice (Scanning) Transmission Electron Microscopy. For structural template is the key to realizing the desired Sn phases. characterization with (S)TEM, NWs were directly transferred to Therefore, this work presents a crucial development to achieve copper grids by a simple dry transfer technique with a clean room a selective Sn phase in the hybrid and, thus, to determine the wipe. HRTEM micrographs were acquired in a TECNAI F20 11801 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article microscope operated at 200 kV. Atomically resolved HAADF data AUTHOR INFORMATION were acquired in a probe-corrected TITAN microscope operated at 300 kV, and a Wiener filter was applied to reduce noise in the Corresponding Authors displayed images. EELS mapping was carried out in a GATAN Sabbir A. Khan − Center for Quantum Devices, Niels Bohr QUANTUM spectrometer coupled to a TECNAI F20 microscope Institute, University of Copenhagen, 2100 Copenhagen, operated at 200 kV with an approximately 2 eV energy resolution and Denmark; Danish Fundamental Metrology, 2970 Ho̷ rsholm, 1 eV energy dispersion, collecting the range of 350−2398 eV of Denmark; Email: skh@dfm.dk energy loss. Principal component analysis (PCA) was applied to the Sara Martí-Sánchez − Catalan Institute of Nanoscience and spectrum images to enhance the S/N ratio. In M4,5, Sn M4,5, Sb M4,5, Nanotechnology (ICN2), 08193 Barcelona, Catalonia, and As L2,3 were the edges employed for areal density mapping. 2D Spain; Email: sara.marti@icn2.cat and 3D atomic models were created using Rhodius software from the Peter Krogstrup − NNF Quantum Computing Programme, University of Cadiz.60 For the presented cross-sections, NWs were Niels Bohr Institute, University of Copenhagen, 2100 transferred to a SiOx substrate using a micromanipulator, where a 100 Copenhagen, Denmark; Email: krogstrup@nbi.dk nm tungsten needle was used to selectively pick the NWs and place Jordi Arbiol − Catalan Institute of Nanoscience and them parallel to each other. Next, an atomic layer deposition system Nanotechnology (ICN2), 08193 Barcelona, Catalonia, was used to deposit 30 nm of AlOx as a protective layer. A focused ion Spain; ICREA, 08010 Barcelona, Catalonia, Spain; beam (Helios 600 dual beam microscope) was used to prepare cross- sectional NW sections. orcid.org/0000-0002-0695-1726; Email: arbiol@ Device Fabrication. NW−Sn devices were fabricated on a highly icrea.cat p-doped Si/SiO2 substrate with a predefined alignment grid. Authors Individual NWs were transferred from the growth wafer to the device substrate with a mechanical micromanipulator carrying a 100 nm Dags Olsteins − Center for Quantum Devices, Niels Bohr tungsten needle and imaged with SEM to define their exact location. Institute, University of Copenhagen, 2100 Copenhagen, Contacts were patterned using electron-beam lithography and e-beam Denmark; orcid.org/0000-0002-2946-7531 evaporation of Ti/Au (5 nm/300 nm). A6 resist was used for spin Charalampos Lampadaris − Center for Quantum Devices, coating for 45 s with 4000 rpm and pumped in a vacuum chamber for Niels Bohr Institute, University of Copenhagen, 2100 3 h. Patterns were developed in 1:3 MIBK/IPA for 45 s, passivated in Copenhagen, Denmark IPA for 30 s, and washed with oxygen plasma for 30 s. Prior metal Damon James Carrad − Center for Quantum Devices, Niels evaporation RF Ar ion milling was employed with 7 W for 5 min at an Bohr Institute, University of Copenhagen, 2100 Copenhagen, 18 mTorr pressure. The deposition rate of Ti and Au was 0.1 and 0.3 Denmark; Department of Energy Conversion and Storage, nm/s, respectively. Acetone was used for lift-off. For the devices of Technical University of Denmark, Lyngby 2800, Denmark Figure 5(b), the distance of the inner probes was defined by the Yu Liu − Center for Quantum Devices, Niels Bohr Institute, design software at 250−500 nm. University of Copenhagen, 2100 Copenhagen, Denmark; orcid.org/0000-0002-1869-5811 ASSOCIATED CONTENT Judith Quiñones − Catalan Institute of Nanoscience and Nanotechnology (ICN2), 08193 Barcelona, Catalonia, Spain * sı Supporting Information Maria Chiara Spadaro − Catalan Institute of Nanoscience and The Supporting Information is available free of charge at Nanotechnology (ICN2), 08193 Barcelona, Catalonia, https://pubs.acs.org/doi/10.1021/acsnano.3c02733. Spain; orcid.org/0000-0002-6540-0377 S1. Unit cells and crystal structure parameters of InSb, Thomas Sand Jespersen − Center for Quantum Devices, Niels InAsSb, InAs, α-Sn, and β-Sn; S2. InSb−Sn morphology Bohr Institute, University of Copenhagen, 2100 Copenhagen, and EELS map; S3. FIB- or ALD-induced damage Denmark; Department of Energy Conversion and Storage, during cross-section preparation; S4. InSb−Sn: β-Sn Technical University of Denmark, Lyngby 2800, Denmark; grain formation within α-Sn shell; S5. Temperature orcid.org/0000-0002-7879-976X effect on the hybrid structure; S6. InSb−Sn grown at Complete contact information is available at: room temperature; S7. InSb−Sn grown at sample holder https://pubs.acs.org/10.1021/acsnano.3c02733 temperature −50 °C; S8. InSb−Sn grown at sample holder temperature −100 °C; S9. InSb−Sn grown at Author Contributions sample holder temperature −150 °C; S10. Additional S.A.K. and S.M.-S. contributed equally to this work. S.A.K HRTEM micrographs of 3 nm Sn on InSb NW; S11. conceived the idea and experimental plans and performed Defect formation and initial β-Sn nucleation in a 5 nm substrate fabrication, materials growth, and morphological Sn shell on an InSb NW; S12. Additional micrographs of characterization. S.M.S. performed the (S)TEM-related experi- a 40 nm Sn shell on an InSb NW; S13. Thermal stability ments, analyzed TEM and EELS data, and prepared the atomic of the α-β-Sn/InSb heterostructure; S14. Additional models. S.A.K., S.M.S., and P.K. proposed the growth SEM images of InAsSb−Sn morphology; S15. Addi- tional structural analysis of InAsSb/Sn; S16. Structural mechanisms based on the experimental analysis. D.O., C.L., analysis of InAs/Sn NWs; S17. β-Sn phase analysis on an and D.J.C. performed the device fabrication, electrical InAs NW; S18. Grain boundary within β-Sn on InAs measurements, and related data analysis. Y.L. assisted with facets; S19. Additional β-Sn analysis on am InAs NW; the materials growth experiments. J.Q. and M.C.S. assisted S20. Additional β-Sn orientations on an InAs NW; S21. with the TEM experiments. S.A.K. and S.M.S. prepared the InSb−InAs(2 nm)−Sn; S22; Summary of 3D models of manuscript with the contribution from all other authors. T.S.J., the α-β/InX configuration; S23. Growth parameters of P.K., and J.A. supervised the project. the fabricated devices; S24. Additional device measure- Notes ment (PDF) The authors declare no competing financial interest. 11802 https://doi.org/10.1021/acsnano.3c02733 ACS Nano 2023, 17, 11794−11804 ACS Nano www.acsnano.org Article ACKNOWLEDGMENTS et al. Parity-preserving and magnetic field−resilient superconductivity in InSb nanowires with Sn shells. Science 2021, 372, 508−511. The project was supported by Danish Agency for Higher (11) Khan, S. A.; Lampadaris, C.; Cui, A.; Stampfer, L.; Liu, Y.; Education and Science, European Union Horizon 2020 Pauka, S. J.; Cachaza, M. E.; Fiordaliso, E. M.; Kang, J.-H.; research and innovation program under the Marie Skłodow- Korneychuk, S.; et al. Highly transparent gatable superconducting ska-Curie Grant No. 722176 (INDEED), Microsoft Station Q, shadow junctions. ACS Nano 2020, 14, 14605−14615. and the European Research Council (ERC) under Grant No. (12) Busch, G.; Kebn, R. Semiconducting Properties of Gray Tin; Solid 716655 (HEMs-DAM). Materials growth of this project is State Physics; Academic Press: New York, 1960; Vol. 11; pp 1−40. performed at NBI Molecular Beam Epitaxy System, and the (13) Ewald, A.; Tufte, O. Gray tin single crystals. J. Appl. Phys. 1958, authors acknowledge C. B. Sørensen for all the maintenance 29, 1007−1009. and keeping the system up and running. ICN2 acknowledges (14) Groves, S.; Paul, W. Band structure of gray tin. Phys. Rev. Lett. funding from Generalitat de Catalunya 2021SGR00457. This 1963, 11, 194. study was supported by MCIN with funding from European (15) Farrow, R.; Robertson, D.; Williams, G.; Cullis, A.; Jones, G.; Union NextGenerationEU (PRTR-C17.I1) and Generalitat de Young, I.; Dennis, P. The growth of metastable, heteroepitaxial films of α-Sn by metal beam epitaxy. J. Cryst. Growth 1981, 54, 507−518. Catalunya and by “ERDF A way of making Europe”, by the (16) Legrain, F.; Manzhos, S. Understanding the difference in European Union. ICN2 is supported by the Severo Ochoa cohesive energies between alpha and beta tin in DFT calculations. AIP program from Spanish MCIN / AEI (Grant No.: CEX2021- Advances 2016, 6, No. 045116. 001214-S) and is funded by the CERCA Programme/ (17) Didschuns, I.; Fleischer, K.; Schilbe, P.; Esser, N.; Richter, W.; Generalitat de Catalunya. The authors acknowledge the use Lüders, K. Superconductivity in Sn films on InSb (110) taking of instrumentation as well as the technical advice provided by account of the film morphology and structure. Physica C: Super- the National Facility ELECMI ICTS, node “Laboratorio de conductivity 2002, 377, 89−95. Microscopías Avanzadas” at University of Zaragoza. M.C.S. has (18) Molodets, A. M.; Nabatov, S. S. Thermodynamic potentials, received funding from the postdoctoral fellowship Juan de la diagram of state, and phase transitions of tin on shock compression. Cierva Incorporation from MICINN (JCI-2019) and the High Temperature 2000, 38, 715−721. 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