Au-Sn catalyzed growth of Ge1-xSnx nanowires: growth direction

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Au-Sn catalyzed growth of Ge Sn nanowires: growth direction, crystallinity and Sn incorporation Yong-Lie Sun, Ryo Matsumura, Wipakorn Jevasuwan, and Naoki Fukata Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.9b02395 • Publication Date (Web): 25 Aug 2019 Downloaded from pubs.acs.org on August 25, 2019

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Au-Sn catalyzed growth of Ge1-xSnx nanowires: growth direction, crystallinity and Sn incorporation Yong-Lie Sun*,1,2,, Ryo Matsumura1, Wipakorn Jevasuwan1, Naoki Fukata*,1,2 1International

Center for Materials Nanoarchitectonics, National Institute for Materials

Science, 1-1 Namiki, Tsukuba, 305-0044, Japan 2Institute

of Applied Physics, University of Tsukuba, 1-1-1 Tennodai, Tsukuba, 305-8573,

Japan

ABSTRACT Ge1-xSnx nanowires (NWs) have been a focus of research attention for their potential in realizing next-generation Si-compatible electronic and optoelectronic devices. To control the growth of NWs and increase their Sn content, the growth mechanism needs to be understood. The use of Au-Sn alloy catalysts instead of Au catalysts allows an easier understanding of Ge1-xSnx NW growth, and the effects of Sn at different concentrations in catalysts on growth direction, Sn incorporation and crystallinity of Ge1-xSnx NWs can be clarified. High Sn content in Au-Sn alloy catalysts favors 〈110〉-oriented NW growth and high Sn incorporation in NWs. The higher Sn content in Au-Sn alloy catalysts also improves the crystallinity of NWs. KEYWORDS: Germanium tin, nanowire, vapor-liquid-solid growth, faceting, Au-Sn

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Germanium Tin (Ge1-xSnx) alloys have attracted major attention for their higher carrier mobility, a direct narrow bandgap structure, and great Si-compatibility.1,2 Theoretical3,4 and experimental5,6 studies have shown that indirect-to-direct bandgap transition can occur if Sn incorporation exceeds 6.5 - 11 at.% in Ge1-xSnx alloys. These narrow and direct bandgap properties make them suitable as infrared photodetectors and lightemitters for optical data communication.7,8 On the other hand, incorporation of Sn in Ge increases electron and hole mobilities by lowering the effective mass of carriers,4,9 making Ge1-xSnx an ideal material for high-speed metal-oxide semiconductor field-effect transistors (MOSFETs). However, the low equilibrium of Sn in Ge (< 1 %) needs to be overcome to realize high-speed MOSFETs and optical devices. The structural change from thin films to nanowires (NWs) leads to potential solutions for overcoming the abovementioned difficulties. Ge1-xSnx NWs can be grown at low temperatures to increase Sn content through non-equilibrium incorporation and to minimize Sn segregation, since the equilibrium solubility of Sn in Ge is extremely low (< 1%).10 These requirements can be fulfilled by using a vapor-liquid-solid (VLS) growth mechanism, which provides a growth temperature below the eutectic point11 and over-equilibrium Sn incorporation through the solute trapping process.12,13 NW geometry with a large surface area can also be used to relax the compressive strain in Ge1-xSnx alloys caused by lattice mismatch,


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reducing the Sn content required for indirect-to-direct bandgap transition.3 Ge1-xSnx NWs with a high Sn content of 13 - 19 at.% have been fabricated by microwave-assisted growth14,15. However, NWs grown by this technique show bending and random growth directions. To date, the growth of high-quality Ge1-xSnx NWs with the highest Sn concentrations has been achieved in the range of 9 - 12 at.% using the VLS method.16,17 Growth kinetics such as temperature-dependent Sn incorporation in NWs18 and the dependence of NW growth rate on Sn concentration in catalysts17 have also been investigated. In these reports, mixed Ge-Sn gas is used to grow Ge1-xSnx NWs. The growth is effected by a two-step process: the incorporation of Ge and Sn into Au catalysts and the subsequent incorporation of Sn into GeNWs during vapor-liquid-solid (VLS) growth, making it difficult to understand the growth mechanism17. Using Au-Sn alloy catalysts instead of Au catalyst simplifies the understanding of Ge1-xSnx NW growth. It is important to elucidate the effects of catalysts on Sn incorporation to understand the growth kinetics and thus enable further increased Sn content in Ge1-xSnx NWs. On the other hand, high incorporation of Sn also causes problems, such as causing misfit dislocation, Sn segregation and twin defects in NWs, resulting in low-quality crystals. Thus, changes in crystallinity should also be taken into consideration for Ge1-xSnx NWs with high Sn content. Compositional changes in catalysts significantly affect the growth


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directions of NWs.19 NWs with different growth directions show different electronic properties.20 Controlling the growth direction of Ge1-xSnx NWs would be beneficial for further tuning the band structure and adjusting physical properties to suit particular applications. Nanowire-based three-dimensional electronic devices, such as gate-allaround FETs, require the vertical alignment of NWs.21 Ge1-xSnx NWs can also be achieved by using core-shell structures, and their optical and elastic properties have been studied22–24. In this study, we used Au-Sn catalyst alloys instead of Au catalysts to investigate the growth of Ge1-xSnx NWs without the need to take account of Sn incorporation into Au catalysts: this simplifies the process of understanding Ge1-xSnx NW growth. We investigate here the growth direction, Sn incorporation and crystallinity of Ge1-xSnx NWs as a function of Sn concentration over a wide range (0 - 86%). Sn concentration in catalysts was controlled by varying the ratio of deposited Au and Sn. Here, metallic Sn in catalysts was used as the Sn source rather than precursor gas and the Au was used to guide the incorporation of Sn in the NW during the VLS growth process. The influence of catalysts on growth direction will be discussed using Schmidt’s model25 by considering the surface tension change of the catalysts. Sn nanoparticles and Au film were deposited on Si(111) substrate in sequence by


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thermal evaporation. The volume ratio of Au and Sn was controlled by varying the deposition thickness of each metal, monitored by a quartz oscillator with a deposition rate of 0.05 nm per second. These samples were then loaded into an ultra-high-vacuum chemical vapor deposition (UHV-CVD) chamber with a background pressure of 2×10–6 Pa. Ge1-xSnx NWs were grown using germane (GeH4) precursor with Au-Sn alloy catalysts via the VLS process. The details of the UHV-CVD method are reported elsewhere.26–28 After pre-annealing at 280 - 400 ˚C for 10 min, Ge1-xSnx NWs were grown at the same temperature for 20 min using 10 sccm of GeH4 as precursor gas. The total pressure was set at 700 Pa by mixing with nitrogen gas. For GeNW growth, a 2 nm-thick Au film was first deposited on Si(111) substrate and dipped in 1% HF solution for 2 min, then immediately placed in the CVD chamber for VLS growth. Scanning electron microscopy (SEM) images were recorded using a Hitachi S-8000 SEM with acceleration voltage of 5 kV. Transmission electron microscopic (TEM) and energy-dispersive X-ray spectrometry (EDX) analysis were carried out using a JEOL 2100F transmission electron microscope operating at 200 kV. Micro-Raman scattering measurements were performed at room temperature using a 532-nm excitation beam focused to a spot size of ~1 μm with a 100× objective. The power of the excitation laser beam was set at about 0.02 mW to prevent local heating effects.29,30 All the data were


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recorded at a spectral resolution of about 0.3 cm–1. X-ray diffraction (XRD) data were collected using a PANalytical X’Pert Pro MRD system with a parallel Cu Kα beam.

Figure 1. 30˚-tilted SEM images of Ge1-xSnx NWs grown at 320 ˚C using catalysts with Au:Sn ratios of (a) 2:5 (b) 1:5 (c) 1:10 (scale bar: 500 nm). SEM images of Ge1xSnx NWs grown at 360 ˚C using catalysts with Au:Sn ratios of (d) 2:5 (e) 1:5 (f) 1:10 (scale bar: 500 nm). The insets show magnifications of typical NWs. Scale bar: 50 nm. (g) TEM of a NW (Au:Sn = 1:10, 320 ˚C). Scale bar: 10 nm. HRTEM images of a NW grown at 360 ˚C using catalysts with Au:Sn ratios of (h) 2:5 (i) 1:5 (scale bar: 5 nm). SAED patterns in the inset confirm the growth direction and the single-crystalline nature of the NWs.


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Au-Sn catalyzed Ge1-xSnx NWs were successfully grown at 320 ˚C and 360 ˚C, as shown in the SEM images in Figure 1. The reason for the low NW density in Figure 1a is that the growth temperature (320 ˚C) is too far below the eutectic temperature (360 ˚C with 60 at.% of Sn).31 Except for Figure 1a, the density of NWs decreases on increasing the added Sn concentration in catalysts at each growth temperature. This can be explained by the larger size of the catalysts (shown in Figure S5c and S5f) and the instability of the catalyst droplets.32 Liquid Sn in the catalyst reduces the surface tension of the Au-Sn-Ge droplets. If there are any perturbations, the droplet will wet the sidewall of the NW and impede its growth. NWs grown at 360 ˚C exhibit a larger diameter (80 nm) than those grown at 320 ˚C (20 nm). This can be attributed to the enlargement of the catalyst droplets due to the faster cracking rate of the germane precursor at the higher temperature and the presence of radial growth resulting from vapor-solid (VS) growth. In addition to this, enhanced surface diffusion of Ge atoms is possible with larger catalyst droplets. The TEM image in Figure 1g shows a straight NW with a spherical seed at the tip, confirming the participation of the Au-Sn catalyst in the VLS growth mechanism. Figure 1h and 1i are high-resolution TEM (HRTEM) images of the 〈 111 〉 -oriented NWs with an approximately 3 nm-thick oxide layer. The corresponding diffraction patterns demonstrate the highly crystalline nature of the NWs. Note that NWs cannot grow at


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temperatures below 320 ˚C or over 360 ˚C: this is likely due to the participation of vaporsolid-solid (VSS) growth at low temperatures and an over-rapid VS growth rate at high temperatures. In addition, the detail information about the Au-Sn catalysts such as thickness of Au-Sn, morphology and the actual Sn concentration is described in the Supporting Information (Figures S5 and S6). The result shows that the actual Sn concentration measured by EDX is consistent with the added amount of Sn. It is essential to study Sn incorporation in Ge1-xSnx NWs quantitatively. TEM-EDX analysis was therefore performed to investigate the composition and distribution of elements in the NWs. The amount of Sn in the catalyst was varied from 0 - 86 at.% by changing the Au:Sn thickness ratio from 1:0 - 1:10, resulting in an increased Sn content in the NWs (in Figure 2a). 1.3 at.%, and 2.2 at.% of Sn were incorporated in NWs by using 61% and 76% of Sn, respectively, in the catalyst. On increasing the added Sn% to 86%, the content of Sn incorporation in the NWs reached 5 at.%. This value far exceeds the equilibrium solubility of Sn in Ge.10 Figure 2b is a linescan of EDX measurements along the length of the NW. The Sn-L peak shows continuous Sn incorporation along the length of the NW. The increase in the intensity of the Ge-K peak towards the bottom of the NW is due to its tapered shape. This tapered structure can be explained by the effect of VS growth and the continuous shrinking of the catalyst at the top of the NWs due to


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the consumption of Sn atoms in the catalyst droplet during growth. The Sn distribution along the NW growth direction (5 at.% of Sn) is shown in Figure 2c. The Sn content at the bottom and the middle region of the NW give a mean Sn content of about 5 at.%, while an extraordinarily high Sn content of 7.3 at.% was estimated in the tip region. Since no catalyst was found either in the TEM images (shown in Figure 5a later) or by elemental mapping using EDX (Figure S2), this high value cannot be attributed to remaining catalyst. A possible explanation for this inconsistency is that the thinner radial growth shell at the tip compared to the bottom results in the higher Sn content, although this could be an overestimation caused by high noise (Figure S1b). EDX mappings in Figure 2d show uniform distribution of Sn inside the NW without any observable segregation. We also performed a linescan along the width of the NW (in Figure S3), which revealed no segregation of Sn at the surface of the NW. No Au M peak can be found from the EDX spectrum (in Figure S1a), indicating negligible Au contamination.


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The Sn content can be also estimated by Raman spectroscopy. The Sn composition dependence of the Raman frequency of Ge1-xSnx alloys has been published by numerous research groups, allowing the amount of Sn located at the substitutional sites in the alloy NWs to be estimated. Figure 3a and 3b show the Raman spectra of NWs grown at 320 ˚C and 360 ˚C, respectively. The intense peak around 300.2 cm–1 for bulk Ge is attributed to the Ge optical phonon peak. The Ge optical phonon peak generally shifts to a lower wavenumber on increasing added Sn% in the catalyst. In Figure 3c, the Ge optical phonon

Figure 2. (a) EDX Sn composition in NWs grown at 360 ˚C with different added Sn concentrations in the catalysts. (b) EDX linescan for Ge and Sn along the length of a NW (Au:Sn = 1:10, 360 ˚C). Increased Ge intensity toward the bottom resulting from the tapered shape of the NW. Scale bar: 500 nm. (c) TEM image of the NW and the average Sn content was estimated by EDX within different regions (in red rectangles). Scale bar: 500 nm. (d) TEM image and EDX mapping for Ge and Sn at the bottom of the NW. Scale bar: 30 nm.


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peaks observed for NWs with 1.3, and 2.2 at.% of Sn (estimated by EDX measurements) showed downshifts of –0.32 and –1.02 cm–1 from the value obtained for the bulk Ge. As shown in Figure 1f, Ge wetting layers are present in our samples. To eliminate the effects of the Ge wetting layer, micro-Raman scattering was performed to separate the background signal of the wetting layers from the NW signal. As shown in the inset of Figure S4, an excitation beam 1 μm in diameter was used to probe the region with and without NWs grown with an Au:Sn ratio of 1:10. The intensity of the Ge optical phonon peak observed from the NW region is about double that from the region without NWs. The Raman signal from the NW (colored black in Figure S4) was obtained by subtracting the Raman spectra of the region without NWs from that with NWs, giving a large Raman shift of –3.26 cm–1. The Raman shift is generally affected by the Sn content and strain in Ge1-xSnx alloys. Here, the strain effect is not taken into account, since strain is likely to be released due to the structural properties of one-dimensional nanostructures. The dependence of the Raman shift on the Sn content in Ge1-xSnx alloys (𝜔𝑎𝑙𝑙𝑜𝑦) can been expressed as a linear equation: 𝜔𝑎𝑙𝑙𝑜𝑦 = 𝜔𝐵𝑢𝑙𝑘 +∆𝜔𝑎𝑙𝑙𝑜𝑦𝑥, with 𝜔𝐵𝑢𝑙𝑘 being the Raman shift of bulk Ge and 𝑥 being the Sn concentration. We therefore fitted the Raman shift (𝜔𝑎𝑙𝑙𝑜𝑦) as a function of Sn concentration ( 𝑥) with this expression and obtained ∆𝜔𝑎𝑙𝑙𝑜𝑦 = ―(60.0 ± 7.5) cm–1. This value is in good agreement with the value of


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―(64.3 ± 0.1) cm–1 for Ge1-xSnx NW16 and ―(68 ± 5) cm–1 for unstrained Ge1-xSnx alloy film.33 This consistency gives further confirmation of the substitutional incorporation of Sn in our samples. The decreased full width at half maximum (FWHM) for the NW with low Sn content in Figure 3d indicates the improvement of crystallinity in the NWs, while the broadening of the FWHM for the NW with high Sn content is due to high Sn incorporation. XRD measurements were performed to further confirm Sn incorporation and crystallinity in the Ge1-xSnx NWs shown in Figure 3e - 3h. The shift in Ge (111) reflections and corresponding calculated lattice constant indicates that the unit cell grows

Figure 3. Raman spectra of Ge1-xSnx NWs grown at (a) 320 ˚C and (b) 360 ˚C, and XRD pattern of Ge1-xSnx NWs grown at (e) 320 ˚C and (f) 360 ˚C with different Au:Sn ratio in the catalysts. The black dashed line shows the tendency of the peak shift. (c) Raman shift of Ge optical phonon peak and (d) corresponding FWHM. (g) Lattice constant calculated from the Ge (111) peak and (h) corresponding FWHM. The green dashed line shows the bulk value. The error bar indicates the standard deviation of the Raman measurements and fitting of XRD spectra. The actual Sn content in the NWs (estimated by EDX) is indicated in (c) and (g).


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with increased added Sn%. The enlargement of the unit cell can be explained by the incorporation of Sn atoms into the Ge lattice according to Vegard’s law. The large difference in lattice constant between the sample (Au:Sn = 1:10, 360 ˚C) compared to the bulk Ge, with Sn = 4.1 at.%, puts it in good agreement with the actual Sn content of 5 at.% as determined by EDX. The FWHM of the Ge (111) peak in Figure 3h shows a similar trend to the result obtained by Raman spectroscopy, which further confirms the improvement of the crystallinity at low Sn content and increased Sn% in catalyst giving high Sn incorporation. A possible explanation for this enhancement of crystallinity is that Ge supersaturation can significantly suppress the homogeneous nucleation rate of the Au.34 Ge supersaturation can be increased by adding Sn to the Au catalyst to reduce the equilibrium concentration of Ge (Equation 2, shown later), resulting in low Au contamination of NWs with high crystallinity. Moreover, an overall decrease of FWHM of the Ge(111) peak for the 360 ˚C-grown NW compared to 320 ˚C could also be explained by the lower nucleation rate of Au at high temperatures.34 Non-equilibrium incorporation of Sn during the Ge NW growth is considered through the solute trapping mechanism carried out under step-flow kinetics.12 A step-flow cycle includes an accumulation of chemical potential called the incubation process, which is followed by a rapid layer addition called step flow. The impurities in the catalyst droplet


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can then be trapped by the sudden growth of the host layers. The amount of impurity incorporation can be increased by a faster trapping rate, which is related to a higher step velocity. System size is one of the factors that can determine the step velocity;13 thus, NWs with a larger diameter give a higher step velocity, resulting in the increased Sn% in NWs. Another item to mention is that performing a quick step flow requires a high Ge supersaturation level in catalysts.13 Both a higher growth temperature with a faster precursor cracking rate and a larger diameter35 due to the Gibbs-Thomson effect will increase Ge supersaturation, resulting in a higher Sn content in the NW. On the other side, experimental results also show that a higher Sn concentration in catalysts will increase the growth rate of Ge1-xSnx NWs,17 which indicate a higher step velocity and the increasing of Sn incorporation in NWs. To make the integration of Ge1-xSnx NWs on a silicon platform more practicable, the key is to control the growth direction of Ge1-xSnx NWs. Figure 1a - 1f clearly show that the NW growth direction depends on the Sn concentration in the catalyst. Since the GeNWs were mainly grown in the 〈111〉 and 〈110〉 directions, but rarely in the 〈112〉 direction,36,37 3D models of 〈111〉 and 〈110〉-oriented NWs grown on (111) substrate were built in Figure 4a and 4b to simulate the same view of 30˚ tilted SEM images for comparison. The relative proportion of NWs with different growth directions grown at


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320 ˚C and 360 ˚C are shown in Figure 4c and 5d. The relative proportion of 〈110〉 direction significantly increased on increasing the Sn% in the catalyst, giving a crossover Sn% of approximately 75, and 80 in the 〈111〉 direction at growth temperatures of 320 ˚C, and 360 ˚C, respectively. This phenomenon can be explained as follows. Diameter dependence of Si NW growth direction has been demonstrated25 while

Figure 4. Schematic 30˚ tilted view of (a) 〈110〉-oriented and (b) 〈111〉-oriented NWs grown on Si(111) substrate. The inset shows the viewing direction. Relative proportion of the different growth directions of the NWs for the growth temperature of (c) 320 ˚C and (d) 360 ˚C. (e) Contour plot of the separation between the free energy per unit circumference of the 〈111〉 and 〈110〉-oriented NWs as a function of NW radius and Sn concentration in a catalyst droplet. The black solid line marks zero energy separation.


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remaining unknown for Ge1-xSnx NWs. Here, we explain how the direction of growth of Ge1-xSnx NWs depends on the diameter of the Ge1-xSnx NWs and the Sn concentration in the catalyst, using a theoretical model established by Schmidt et al.25 The dependence on Sn concentration was integrated into the Schmidt model as shown in Equation 1, 𝑓〈𝑢𝑣𝑤〉 = ∆𝑧𝜎〈𝑢𝑣𝑤〉 + 𝑎〈𝑢𝑣𝑤〉𝜎〈𝑢𝑣𝑤〉 𝑟 𝑠 𝑙𝑠

(1)

where 𝑓⟨𝑢𝑣𝑤⟩ is free energy per unit circumference of a 〈uvw〉 oriented nanowire, ∆𝑧 is the interfacial thickness of the Ge1-xSnx NWs side, 𝜎⟨𝑢𝑣𝑤⟩ is the surface energy of a 𝑠 〈uvw〉-oriented NW, 𝑎⟨𝑢𝑣𝑤⟩ is the geometrical parameter of the cross section of a 〈uvw〉oriented NW, 𝜎⟨𝑢𝑣𝑤⟩ is the interfacial tension of the liquid-solid interface, and 𝑟 is the 𝑙𝑠 radius of the NW. An intermixed region of Si and Au of an Au thin film on a Si substrate has been calculated by molecular dynamics simulation.38 Considering this simulation result, Schmidt et al. set ∆𝑧 at 1 nm for the growth of SiNWs and explained the relationship between the diameter and growth direction of SiNWs. Biswas et al. performed highresolution EELS mapping of Ge1-xSnx NWs and estimated the intermixed region to be about 1 nm.16 Based on all these results, we can also set ∆𝑧 at 1 nm for the growth of Ge1-xSnx NWs. Here, we only consider 〈111〉 and 〈110〉-oriented NWs. The preferential growth direction of growth can be estimated from the difference between 𝑓⟨111⟩ and


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𝑓⟨110⟩. The calculation for estimating the surface energy of 𝑓⟨111⟩ and 𝑓⟨110⟩ alloys is described in detail in the Supporting Information. A contour plot of energy separation between 𝑓⟨111⟩ and 𝑓⟨110⟩ is shown in Figure 4e. Positive and negative values respectively indicate 〈 110 〉 and 〈 111 〉 to be preferential. The solid black line located at the energy separation of zero shows that radius 𝑟0 increases as the Sn concentration in catalyst increases, which means that, with higher Sn%, NWs tend to grow in the 〈110〉 rather than the 〈111〉 direction. The reason is that higher Sn% reduces liquid-solid interfacial energy, making the surface energy of the NW become more dominant than the interfacial energy, and vice versa. For NWs grown at 320 ˚C, a transition between the 〈110〉 and 〈111〉 orientation takes place at Sn% of 75 in Figure 4c, which corresponds to a radius 𝑟0 ≈ 10 nm at the energy separation of zero in Figure 4e. This value is in good agreement with the average diameter measured from the bottom of NWs (20 nm). However, at a growth temperature of 360 ˚C, the transition Sn% of 80 gives a radius of 𝑟0 ≈ 10 nm, which does not match the average diameter of 60 nm after subtracting the shell thickness measured from Figure 5b. One of the most likely reasons for this discrepancy is the decoration of Au-Sn alloy to the NW sidewall. The Sn-rich catalyst droplets are unstable due to their low surface tension, which permits them to wet and decorate to the NW sidewall.32 This decoration will


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considerably alter the surface free energy of the system and lead to 〈110 〉-oriented growth.39 Control experiments were carried out to prove that the total thickness of Au-Sn had almost no influence on the growth direction of Ge1-xSnx NWs as shown in Figure S7. The use of bimetallic alloy seeds for growing NWs can affect the crystalline features of NWs,40 while NWs grown by single-metal catalysts generally show single-crystal structures.26–28 HRTEM images with corresponding selected area electron diffraction (SAED) patterns were acquired to investigate the crystallinity of NWs grown at 360 ˚C with the highest Sn% in the catalyst. The NW appears to be a single crystal and, judging by the HRTEM image (Figure 5a and 5b), is covered with a 2-nm oxide layer; however, some extra spots are observed in the SAED pattern, indicating hidden defects in the NW. Figure 5c shows the HRTEM image for three different regions. Fast Fourier transform (FFT) analysis was performed on regions A and B in the HRTEM image in Figure 5c, showing them to correspond to different crystallographic grains. In the middle region, a periodicity of three {111} planes appears gradually and reaches the highest contrast at the center, then changes to another type of grain. The corresponding SAED pattern in Figure 5d shows a pseudo-hexagonal symmetry with some additional spots. These odd phenomena can be explained by the superposition of two twinned grains.41,42 As shown in Figure 5f, grains A and B are partially superposed on each other along the incident


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electron beam, naturally generating a 3𝑑111 period that closely agrees with the TEM observation. A possible configuration of grains A and B in the NW is illustrated in Figure 5e. The relative proportion of grain A decreases from left to right and reaches 0.5 at the center, resulting in the contrast change that is observed from the TEM image. The corresponding SAED pattern also can be explained by superposing FFT images of region A on region B as shown in Figure 5d, and the presence of many additional spots is caused

Figure 5. (a) HRTEM image of the tip of a Ge1-xSnx NW (Au:Sn = 1:10, 360 ˚C). (b) HRTEM image of the bottom of the NW with the corresponding SAED pattern in the inset. Radial growth 10 nm in thickness can be observed. (c) HRTEM image taken from the middle of the NW. The insets show the FFT patterns corresponding to region A (red rectangles) and region B (blue rectangles). (d) Corresponding SAED pattern showing pseudo-hexagonal symmetry. (e) Cross-sectional schematic illustration showing a possible configuration of grains A and B in the NW and an electron beam-generated TEM image. (f) Atomic model of the {111}twinned grain A and B with a partial superposition in the middle. The view is perpendicular to the electron beam. All scale bars: 5 nm.


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by the double diffraction effect. Figure 5a shows that these two twinned grains also appear at the tip of the NW, indicating a longitudinal twin plane along the NW growth axis. Twin defects can be induced in GeNWs if the twin energy of the catalyst is smaller than ∆𝜇/𝑆.42 Here, ∆𝜇 is a thermodynamic quantity called supersaturation, which is defined as the chemical potential difference between the Ge in the catalyst droplet and that in the NW. 𝑆 is the inverse of the nucleus density on a {111} plane. According to Biswas et al.,40 supersaturation can be expressed as

()

∆𝜇 = 𝑘𝑇 𝑙𝑛

𝐶 𝐶𝑒

(2)

where 𝑘 is the Boltzmann constant, 𝑇 is the temperature, 𝐶 is the concentration of the growth species (here it is Ge) in a liquid or solid and 𝐶𝑒 is the equilibrium concentration of the growth species. High Sn% in the catalyst gives an extremely low 𝐶𝑒 due to the low solubility of Ge in Au-Sn liquid on the Sn-rich side,43 leading to high supersaturation, ultimately enhancing twin formation. On the other hand, the lower stacking fault energy of Au-Sn alloy compared to Au44 could also be the reason, since twin density usually increases on decreasing the stacking fault energy.45 Incidentally, these twin defects observed in our samples are unlikely to affect the Sn incorporation, since the impurities in the catalyst are first incorporated in the NW with a uniform distribution and then segregate to the defects.46


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In summary, the growth of Ge1-xSnx NWs closely depends on the Sn content in Au-Sn catalysts. Using Au-Sn alloy catalysts simplified the understanding of the growth of Ge1xSnx

NWs. The effects of Sn concentration in the catalysts on the growth direction, the

Sn incorporation into GeNWs, and the crystallinity of Ge1-xSnx NWs were elucidated. By applying a model proposed by Schmidt et al., we were able to successfully explain the experimental results and illustrate that the main factor is the surface energy. NWs grown with high Sn% in catalysts favor the 〈110〉 growth direction whereas NWs with low Sn% favor the 〈111〉 direction. Ge1-xSnx NWs with Sn content up to 5% were successfully grown by VLS growth using Au-Sn catalysts. High-Sn% catalysts enhance the incorporation of Sn and the crystallinity of NWs. Using Au-Sn alloy catalysts also induced twin defects in NWs with high Sn%. High Sn% in the catalyst gives an extremely low 𝐶𝑒 due to the low solubility of Ge in Au-Sn liquid on the Sn-rich side, leading to high supersaturation, ultimately enhancing twin formation. Interesting electrical, optical and thermoelectric properties can be expected by manipulating and controlling twinning in Ge1-xSnx NWs. Gaining a fuller understanding the role of twinning in VLS growth also has the potential to provide effective tools for nanostructure and band structure engineering.


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ASSOCIATED CONTENT Supporting Information. Calculation of free energy of 〈111〉 and 〈110〉-oriented nanowires, additional EDX and Raman data of Ge1-xSnx NWs, morphology and Sn concentration of Au-Sn catalysts, and control experiment data of the growth direction of Ge1-xSnx NWs.

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected]. *E-mail: [email protected].

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS: This work was supported by JSPS Kakenhi (No. 26246021, 17H07351), The Murata Science Foundation, and the World Premier International Research Center Initiative (WPI Initiative), MEXT, Japan.


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Table of Contents graph:


Au-Sn catalyzed growth of Ge1-xSnx nanowires: growth direction

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