# Exploring the Dynamics of Sn Droplets on Hydrogenated Amorphous Si and Ge Layers for Nanocrystal Design

Xiaoyang Zhang, Xuechun Zhao, Lulu Zheng, Ileana Florea, Fabien Cuvilly, Etienne Talbot, Pere Roca I Cabarrocas, Wanghua Chen

► To cite this version:

Xiaoyang Zhang, Xuechun Zhao, Lulu Zheng, Ileana Florea, Fabien Cuvilly, et al. Exploring the Dynamics of Sn Droplets on Hydrogenated Amorphous Si and Ge Layers for Nanocrystal Design. Crystal Growth & Design, 2025, 25 (13), pp.4932-4939. 10.1021/acs.cgd.5c00423. hal-05132766

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# Exploring the Dynamics of Sn Droplets on Hydrogenated Amorphous Si and Ge Layers for Nanocrystal Design

Xiaoyang Zhang, Xuechun Zhao, Lulu Zheng, Ileana Florea, Fabien Cuvilly, Etienne Talbot, Pere Roca i Cabarrocas, and Wanghua Chen*

https://doi.org/10.1021/acs.cgd.5c00423

<figure><img src="image_2.png" /></figure>

**ABSTRACT:** Studying the interaction between metal droplets and amorphous layers can unlock the potential to nanocrystal design, enhance material properties, and provide fundamental knowledge. With the help of plasma-enhanced chemical vapor deposition (PECVD), this work explores the interplay between Sn metal droplets and amorphous layers (Si and Ge) by investigating two growth modes: standard mode (amorphous layer on metal droplets) and inverted mode (metal droplets on an amorphous layer). Based on these modes, in-plane NWs fabricated from a variety of amorphous layer stacks including single layers, bilayers, and up to 10 layers are compared and discussed. We found that nanostructures are obtained only when Sn droplets can interact directly with an amorphous Ge layer in the inverted mode. A wide variety of nanocrystals including NWs, nanobelts, and nanopillars are obtained when using an amorphous multilayer stack. The results are discussed based on surface tension and interfacial diffusion phenomena.

## 1. INTRODUCTION

The interaction between metal droplets and amorphous layers has exciting potential across various fields.¹⁻⁴ Indeed, amorphous layers can create tunable surfaces with controlled wettability and adhesion. For example, by precisely arranging metal droplets on an amorphous layer, one can engineer materials with exotic optical properties.⁵,⁶ The intimate contact between metal droplets and amorphous layers can create highly active catalytic sites for various chemical reactions.⁷,⁸ Integrating metal droplets and amorphous layers in photo-electrodes can improve their light absorption and charge separation efficiency.⁹ Therefore, understanding the interaction between metal droplets and amorphous layers is a critical issue in numerous material fabrication processes. The interaction between metal droplets and amorphous layers is determined by the wetting behavior of the metal droplets on the amorphous layer, influencing how nanostructures nucleate and grow.¹⁰⁻¹² Different metal-amorphous combinations can lead to diverse nanostructure shapes such as cylindrical, tapered, or branched, tailoring them for specific applications.¹³,¹⁴ Importantly, diffusion of atoms from the amorphous layer into the metal

droplet during annealing is a key step.¹⁵⁻¹⁷ By modulating the interaction through the choice of the metal or amorphous material, one can tune the electrical, optical, and mechanical properties of the nanostructures. For example, switching the metal droplets from In to Pb can change the orientation of Si NWs from out-of-plane to in-plane.¹⁸

In this work, the interaction between Sn metal droplets and the amorphous layers is studied in detail. Two configurations are explored including Sn metal droplets-on-amorphous layer (inverted growth mode) and amorphous layer-on-Sn metal droplets (standard growth mode). We systematically investigate the influence of thin layers on the growth of nanostructures by varying the amorphous layer materials (a-
Si:H or a-Ge:H), their thickness, and the deposition order of both the amorphous layer and metal catalyst droplets. Using scanning electron microscopy (SEM), X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), Kelvin probe force microscopy (KPFM), high-angle annular dark field-scanning transmission electron microscopy (HAADF-STEM), high-resolution TEM (HRTEM), and energy dispersive X-ray spectroscopy (EDS), we comprehensively characterize the resulting nanostructured morphology. These detailed observations are then interpreted within the framework of the surface tension and interfacial diffusion processes.

## 2. EXPERIMENTS

Two configurations of metal droplets and amorphous layers are illustrated in Figure 1. In the standard mode, the amorphous layer is

<figure><img src="image_3.png" /><figcaption>**Figure 1.** Schematic illustrations of two configurations of metal (Sn) droplets and amorphous layers. (a) Standard mode with amorphous layer-on-metal droplets. (b) Inverted mode with metal droplets-on-amorphous layer.</figcaption></figure>

deposited on top of the metal droplets (Figure 1a), while in the inverted mode, the metal droplets are deposited on top of amorphous layers (Figure 1b). Here, the metal droplet is Sn and the amorphous layers correspond to hydrogenated amorphous Si (a-Si:H) or hydrogenated amorphous Ge (a-Ge:H).

First, we started with the widely studied standard growth mode. Building upon our previous studies,¹⁹ SnO₂ NPs were used to produce Sn metal droplets. The detailed SnO₂ NPs preparation conditions can be found in our previous work.²⁰ To generate metallic Sn particles, H₂ plasma was ignited at 235 °C with 10 W RF power, 130 Pa pressure, and 100 sccm H₂ flow for 10 min. Then, a 10 nm a-Si:H layer was deposited on the top of metallic Sn NPs once the sample temperature was cooled down to 120 °C under high vacuum conditions.

In the case of Sn NPs with the inverted mode, the pretreatment process is more complex. The inverted mode involves opening of the vacuum chamber to deposit the metal on the amorphous layers. Moreover, Sn does oxidize easily once it is exposed to air. As a consequence, five distinct process paths based on the inverted mode were attempted, as detailed in Figure 2. Path 1 comprised three primary steps: the sample was first dipped in a 5% HF aqueous solution for 30 s to remove the native oxide layer before being loaded

into the plasma enhance chemical vapor deposition (PECVD) reactor, and various conditions were used to reduce SnO₂ NPs as in the standard mode. Then, the sample underwent annealing in 100 sccm H₂ at 235 °C for 30 min to promote NW growth. However, path 1 failed to produce NWs. Building upon path 1, subsequent paths explored strategies to address its limitations. Path 2: Omitting the HF dipping and inserting a low-power H₂ reduction step (5 W, 5 min) after the standard reduction (both at 235 °C) produced no NW. Path 3: The HF dip and H₂ reduction orders were swapped. However, re-exposure to air after the HF dip reoxidized the surface, impeding direct catalyst/amorphous contact. Path 4: Interestingly, an additional H₂ reduction step activated the GeSn NW growth. Path 5 further enhanced NW activation with an additional postreduction annealing step (235 °C). These results highlight the critical role of the catalyst/amorphous interface and the effectiveness of tailored H₂ reduction and annealing treatments for optimizing NW growth in inverted mode.

The deposition conditions of the amorphous layers (a-Si:H and a-Ge:H) were systematically studied under two growth modes. For a-Si:H, deposition was conducted at 200 °C with a pressure of 50 Pa, RF power of 5 W, H₂ flow rate of 50 sccm, and SiH₄ precursor flow rate of 2 sccm. For a-Ge:H, deposition was performed under identical temperature and pressure conditions (200 °C, 50 Pa), with an RF power of 5 W and GeH₄ precursor flow rate of 50 sccm (1% diluted in H₂). Specifically, for the deposition of hydrogenated amorphous SiGe (a-SiGe:H) layers, the temperature, pressure, and RF power were kept constant with a H₂ flow rate of 50 sccm. Simultaneously, SiH₄ and GeH₄ were introduced with corresponding flow rates being 5 and 50 sccm (1% diluted in H₂), respectively.

A KPFM system from Asylum Research (Oxford Instruments) equipped with a conductive atomic force microscopy (AFM) tip (AC240TM from Olympus) featuring doped Si cantilevers and a Pt coating was used to characterize the electrical properties of the samples. A Thermo Scientific ESCALAB QXi XPS microprobe was used to acquire XPS spectra. Morphological and structural analyses were performed on a Titan Themis electron microscope from Thermo Fisher Scientific (Waltham, MA, USA) operating at 200 kV. This microscope boasts a Cs aberration probe corrector for enhanced image resolution and a Super X detector, enabling chemical analysis of both light and heavy elements with nanoscale spatial resolution.

## 3. RESULTS AND DISCUSSION

Table 1 summarizes the investigation of 11 samples grown using both standard (amorphous on catalyst) and inverted (catalyst on amorphous) growth modes. SEM revealed varying surface morphologies across the samples. First, the case of single-layer films is discussed. In the following single-layer film cases, the thicknesses of the deposited films are all approximately 10 nm. In-plane NWs using the standard mode on a single a-Si:H layer deposited on Sn catalysts can be

<figure><img src="image_4.png" /><figcaption>**Figure 2.** Flowchart of the various attempts aiming at producing NWs via the inverted growth mode ("X" represents "no NW growth"; "√" represents "NW growth".) The annealing temperature was 235 °C for all cases.</figcaption></figure>
<table><caption>**Table 1.** Schematic Illustration and SEM Images of Different Combinations of Amorphous Layers and Metal Catalyst as well as the Corresponding Growth Results</caption><thead><tr><th>Sample name</th><th>Illustration</th><th>Low magnification</th><th>High magnification</th><th>Result</th></tr></thead><tbody><tr><td>1</td><td>a-Si:H<br>Sn<br>c-Si wafer</td><td></td><td></td><td>NW growth</td></tr><tr><td>2</td><td>Sn<br>a-Si:H</td><td>5 µm</td><td></td><td>Fail</td></tr><tr><td>3</td><td>a-Ge:H<br>Sn</td><td>1 µm</td><td>200 nm</td><td>NW growth</td></tr><tr><td>4</td><td>a-Ge:H</td><td>2 µm</td><td>150 nm</td><td>NW growth</td></tr><tr><td>5</td><td>Sn<br>a-SiGe:H</td><td>2 µm</td><td></td><td>Fail</td></tr><tr><td>6</td><td>a-Si:H<br>a-Ge:H<br>Sn</td><td>2 µm</td><td>300 nm</td><td>NW growth</td></tr><tr><td>7</td><td>Sn<br>a-Si:H<br>a-Ge:H</td><td>2 µm</td><td>500 nm</td><td>Fail</td></tr><tr><td>8</td><td>a-Ge:H<br>a-Si:H<br>Sn</td><td>5 µm</td><td>300 nm</td><td>Fail</td></tr><tr><td>9</td><td>Sn<br>a-Ge:H<br>a-Si:H</td><td>20 µm</td><td>5 µm</td><td>NW, NB, NC growth</td></tr><tr><td>10</td><td>Sn</td><td>2 µm</td><td>1 µm</td><td>NW, NP growth</td></tr><tr><td>11</td><td>Sn</td><td>2 µm</td><td>400 nm</td><td>NW growth</td></tr></tbody></table>

grown, as demonstrated by sample 1. Interestingly, however, inverting the deposition order by placing Sn catalysts on a-Si:H (sample 2) completely prevents NW growth. On the other hand, both growth modes readily produce NWs when an a-Ge:H layer is used. Sample 3 showcases limited NW activation in the standard a-Ge:H-covered catalyst configuration. Only particles between 100 and 200 nm in diameter exhibit activation, while the majority of the observed catalyst particles remain large and micron-sized. In contrast, inverting the catalyst position (sample 4) significantly enhances activation. Activated particles here range from tens of nanometers to 1 µm, and the NW morphology exhibits greater diversity. Notably, an enlarged view reveals zigzag-shaped NWs. Motivated by the apparent role of Ge in promoting NW growth in the inverted mode, we explored an a-SiGe:H alloy as

a layer precursor in sample 5. Unfortunately, this attempt yielded only small nanostructures.

In the following part, we implemented a bilayer approach using a-Si:H/a-Ge:H combinations. In the double-layer film structure, the thickness of the a-Si:H film is approximately 4.9 nm, and the thickness of the a-Ge:H film is approximately 5.3 nm, making the total thickness of these films close to that of a single-layer film. Four sample configurations were devised across both growth modes (samples 6–9 in Table 1). Analysis of the results reveals a trend that NW growth occurs only when the a-Ge:H layer directly contacts the catalyst. This aligns with observations from the single-layer experiments, suggesting Ge’s superior incorporation into the Sn catalyst compared to Si. Depositing the catalyst first in the bilayer configuration (sample 6) yields outcomes like those in sample 3. While the catalyst size distribution shrinks, the activation rate remains low. Conversely, the inverted bilayer arrangement (sample 9) demonstrates a wealth of nanostructures, including not only NWs but also nanobelts (NBs) and nanocrystals (NCs).

Two configurations were investigated, including 6-layer (sample 10) and 10-layer (sample 11) structures (Table 1). For the 6-layer amorphous structure, the thickness of each a-Si:H layer and a-Ge:H layer is approximately 2.1 nm. For the 10-layer amorphous structure, the thickness of each a-Si:H layer and a-Ge:H is approximately 1 nm to maintain consistency in total thickness with that of the aforementioned single-layer and double-layer film structures. Sample 10 exhibits NWs with intriguing characteristics. Their diameters show a degree of periodic variation, and a unique nanostructure we term “nanopillars (NPs)” as shown in the inset. Sample 11 presents NWs with a similar overall shape but more pronounced periodic diameter changes and larger diameters. These observations suggest that while increasing the number of layers in the multilayer stack may not significantly alter the basic NW shape, it could influence dimensions (diameter, length) and potentially the activation rate.

Figure 3 shows the XPS spectra for Si 2p, Ge 3d, and Sn 3d of bilayer and multilayer samples including a standard mode sample (sample 8 without NW growth) and inverted mode samples (samples 9, 10, and 11 with NW growth). Five elements were recorded for the corresponding elements of C 1s, O 1s, Si 2p, Ge 3d, and Sn 3d. Current traditional correction methods relying on adventitious carbon have been demonstrated to be unreliable. Meanwhile, the energy level of adventitious carbon being in alignment with the vacuum level of the sample has been established as scientifically valid and dependable when the C 1s peak is used.²¹ Precisely, a direct relationship exists between the C 1s peak position ($E_B$) and the sample work function ($\Phi_{SA}$) that is quantified by the equation $E_B = 489.58 - \Phi_{SA}$.²² In this study, we use the work function of SiGe (5.02 eV)²³ to calibrate the binding energy of C 1s (284.56 eV). Figure 3a shows that the Si 2p3/2 and Si 2p1/2 in sample 8 correspond to 98.7 and 99.3 eV, respectively, as well as two oxidation states, $Si^{3+}$ (101.5 eV) and $SiO_2$ states (103.2 eV), which is in agreement with the previous literature.²⁴ In samples 9, 10, and 11, the intensity of the $SiO_2$ peak was significantly enhanced. Figure 3b shows the Ge 3d XPS spectra, indicating that a small part of sample 8 is oxidized into $GeO_2$. However, in samples 9 to 11 where nanostructures were grown, Ge exists mainly in its elemental state.²⁵ In addition, the XPS spectra of Sn 3d in Figure 3c show that in the standard growth mode (sample 8), the possible
<figure><img src="image_5.png" /><figcaption>**Figure 3.** High-resolution XPS spectra of (a) Si 2p, (b) Ge 3d, and (c) Sn 3d for bilayer and multilayer samples including the standard mode sample (sample 8) and inverted mode samples (samples 9, 10, and 11).</figcaption></figure>

<figure><img src="image_6.png" /><figcaption>**Figure 4.** SEM image of different nanostructures for (a) sample 9 and (b) sample 10, respectively. (c) Plots of trench width distribution for four types of nanostructures.</figcaption></figure>

forms of Sn elements are elemental Sn, SnO, and SnO₂. In samples 9–11, the Sn element is mainly in the form of SnO and SnO₂.²⁶

Figure 4a presents a detailed SEM image of sample 9 with annotations. White double arrows highlight the channel width, while red arrows point to NBs and NCs. Notably, the NCs demonstrate a remarkably ordered arrangement within the trench, maintaining a consistent spacing. Interestingly, regions of coexistence with NCs become apparent along the path of NW growth. Figure 4b showcases another intriguing observation: the presence of both NWs and nanopillars on sample 10. Here, a clear distinction between these two nanostructures is evident. Notably, the catalyst particles at the NP ends are significantly larger, reaching approximately 1 µm. This observation provides compelling evidence that distinct nanostructures can arise on the same sample depending on the size of the catalyst particles. To investigate the potential link between the trench width and nanostructure type, we employed ImageJ for precise measurements. Figure 4c reveals a striking trend: the average trench width gradually decreases across the four identified nanostructures. The inset table further details this observation, showcasing the range and mean values for each category from NCs: 1.42–8.7 µm (mean: 3.66 µm), NBs: 0.67–1.76 µm (mean: 1.22 µm), and NPs: 0.47–

0.87 µm (mean: 0.64 µm) to NWs: 0.18–0.45 µm (mean: 0.36 µm).

Figure 5a presents an SEM image of an island-chain Si/Ge heterostructure NW grown from a six-layer stack (sample 10). Strikingly, the diameter of the nanostructure exhibits significant variation throughout its length, with some sections showing some degree of periodic fluctuation. To gain further insight into these longitudinal variations, we employed AFM for topographic characterization of a selected NW segment. Figure 5b shows a convex morphology of its surface. In the same region, KPFM was used to probe the contact potential difference (CPD). As shown in Figure 5c, two brighter areas appear on the nanostructure, indicating a higher CPD compared to the surrounding regions. This localized variation in CPD suggests a potential change in the NW composition along the segment, leading to the inference of two distinct compositional domains within the measured area. The observed CPD variations on the NW (Figure 5c) suggest intriguing possibilities. For example, accumulation of holes (electrons) in Ge-rich islands (Si-rich chains) could lead to higher (lower) electrostatic potential, indicating the formation of Si/Ge heterojunctions. However, changes in surface height can also influence CPD. To conclusively demonstrate the presence of heterojunctions, we need to isolate the effect of height variations. Similarly, differentiating NBs from NWs
<figure><img src="image_7.png" /><figcaption>**Figure 5.** (a) SEM image of island-chain Si/Ge hetero-nanostructure produced from the six-layer layer approach (sample 10). (b) Morphology of the nanostructure obtained by AFM analysis. (c) Contact potential difference of the nanostructure by KPFM scan. (d) SEM image of NB from the two-layer layer case (sample 9). (e) Morphology of NB obtained by AFM analysis. (f) Contact potential difference of NW by KPFM scan.</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>**Figure 6.** (a) Variation of the diameter of the NW and NB in the section indicated by the white arrow in Figure 5a,d as a function of length. Distribution of height (blue) and potential (red) changes in (b) NW and (c) NB extracted along the corresponding scan lines in Figure 5 b,e and c,f, respectively.</figcaption></figure>

requires careful analysis of both morphology and structure. Figure 5d presents the SEM image of a nanostructure grown from a 4.9 nm a-Si:H/5.3 nm a-Ge:H bilayer layer (sample 9). AFM characterization of the selected area (Figure 5e) reveals its distinct NB shape. The surrounding dark trench indicates the layer consumed by the Sn droplets during growth. Further analysis reveals the nanostructure's distinctive characteristics. Although the width of the NB approaches 100 nm, its maximum height barely reaches 12 nm, resulting in a remarkable width-to-height ratio of approximately 10:1, consistent with a belt-like morphology. Interestingly, the NB's height closely matches the total thickness of the deposited a-Si:H/a-Ge:H bilayer layer (10.2 nm). To investigate the potential presence of Si/Ge heterojunctions within the NB, KPFM measurements were conducted, as shown in Figure 5f. However, in contrast to the six-layer configuration presented in Figure 5c, the CPD profile along the

NB in Figure 5f (two-layer configuration) remains relatively uniform. This observation suggests that no apparent Si/Ge heterostructure has formed within NB grown from the bilayer sample.

To accurately quantify the diameter variation of NWs and NBs as a function of length, we performed image analysis using ImageJ. The segments indicated by white arrows in Figure 5 a,d were measured, and the results are presented in Figure 6a. The black curve represents the measured NW diameter data, while the straight line indicates an average diameter of 110 nm. Notably, the diameter exhibits a distinct simple periodic variation, oscillating within a 60 nm range above and below the average value, suggesting strong underlying regularity. The spacings between adjacent peaks in the diameter profile were determined to be 210, 325, 391, 268, 283, 219, 172, 206, 296, and 99 nm, respectively. Analyzing the red curve in Figure 6a, we observe a similar harmonic variation for the nanobelt
diameter. This oscillation occurs within a range of 50 nm above and below the average diameter of 200 nm. Notably, the spacings between adjacent peaks in the NB profile differ from those of the NWs, measured as 357, 663, 204, 102, 306, 561, 53, and 608 nm, respectively. This translates to a maximum lateral diameter change of approximately 357 nm for the NBs. Comparing the NW and NB data in Figure 6a, we can draw several key conclusions. NBs exhibit a smaller periodic variation in diameter compared to NWs. The larger amplitude of NW diameter changes makes them more readily observable in the figure. The average diameter of NBs is roughly twice that of NWs, which aligns with the trench width measurements.

Figure 6b,c depicts the height (blue) and CPD (contact potential difference, red) variations along scan lines across NW and NB, respectively, extracted from Figure 5b,e and c,f. Clear peaks in the NW CPD profile (Figure 6b, red curve) exhibit prominent differences of up to 17 mV, aligning with the bright areas in Figure 5c. Notably, the height profile (blue curve) mirrors this trend with peak positions matching those of the CPD curve. This correspondence suggests a strong correlation between the height and CPD variations in NWs. In contrast, the NB data in Figure 6c reveal a distinct pattern. CPD peaks and valleys do not correspond to those in the height profile, implicating decoupling of height and CPD variations in NBs. Thus, NWs exhibit a clear correlation between height and CPD, suggesting a potential height-induced CPD modulation. NBs demonstrate no such correlation, indicating that CPD changes arise from factors independent of the height variations.

To further investigate the morphological and compositional characteristics of these nanostructures, we employed STEM analyses. Based on the insights from Table 1 and Figure 7a, we

<figure><img src="image_9.png" /><figcaption>**Figure 7.** (a) SEM image of an NB with the red square indicating the location of the lamella extracted by FIB from sample 9 in Table 1. The inset is an HAADF-STEM image of the Ge-rich part of the NB indicated by the white arrow in the red box in panel a and the STEM image. (b) The corresponding EDS mapping of Ge, O, Si, and Sn elements.</figcaption></figure>

selected a NB exhibiting distinct Ge-rich segments (sample 9) for in-depth cross-sectional examination. The lamella was extracted by focused ion beam (FIB). The HAADF-STEM image in the inset of Figure 7b further reveals the morphology of the Ge-rich region marked by a white arrow within the red box in Figure 7a. Consistent with the findings from the AFM characterization in Figure 5e, the cross-section displays a belt-like profile. Figure 7b presents the corresponding EDS mapping of Ge, O, Si, and Sn atoms within the inset in Figure 7a. Notably, Ge, O, and Sn atoms exhibit a homogeneous distribution within the NB, whereas Si atoms are practically absent. This observation can be attributed to the significantly

higher solid solubility of Ge in Sn compared to Si.²⁷ Consequently, the incorporation rate of Ge atoms into Sn droplets at the growth interface surpasses that of Si. This phenomenon effectively allows the Sn droplet to “float” over the a-Si:H layer, preventing the incorporation of Si into the resulting NWs.

HRTEM observations were done on nanowires corresponding to sample 11 for accessing information on the sample microstructure and crystalline quality. For this, a preliminary standard lift-out process was used for preparing TEM lamella using the focused ion beam (FIB) procedure. Figure 8a

<figure><img src="image_10.png" /><figcaption>**Figure 8.** HRTEM characterization of two NWs from sample 11. (a) Top-view SEM image of NWs where the TEM lamella was prepared. (b) Cross-section TEM image including several NW segments. (c) TEM image of NW 1 with the corresponding FFT (selected area marked by green dashed square) in panel d. (e) Enlarged view of the atomic plane of NW 1. (f) TEM image of NW 2 with the corresponding FFT (selected area marked by blue dashed square) in panel g. (h) Enlarged view of the atomic plane of NW 2.</figcaption></figure>

presents the SEM image where the TEM lamella (Figure 8b) was prepared. TEM of one of the NW segments (NW 1) is presented in Figure 8c. The fast Fourier transformation (FFT) of a selected area marked by a green dashed square is shown in Figure 8d, indicating the crystalline structure of the studied NWs. The enlarged view of Figure 8c is presented in Figure 8e. TEM of another NW segment (NW 2) is shown in Figure 8f with the corresponding FFT (selected area marked by blue dashed square) and enlarged image presented in Figure 8g and h, respectively.

As far as Sn is considered, as seen in Table 1, utilizing Sn as a catalyst and a single a-Si:H layer for SiSn NW growth leads to divergent outcomes between the two modes. While the standard mode allows NW formation (sample 1), the inverted mode does not (sample 2). Conversely, introducing a single a-Si:H layer yields the growth of Ge NWs in both standard and inverted scenarios (samples 3 and 4). This disparity can be attributed to the different wetting behaviors between Sn droplets and the underlying amorphous layers. In the case of single-layer films, the interaction between the Sn droplet and the single-layer a-Si:H depends on the contact area and diffusion between the two. Under the standard mode, due to
the large contact area and the effect of surface tension, it is
difficult for Sn atoms to diffuse into the a-Si:H layer, while
silicon atoms in the a-Si:H layer can easily diffuse into the Sn
droplet, increasing the effective diffusion of Si atoms. However,
under the inverted mode, there is no difference in the diffusion
interface between Sn and a-Si:H. Therefore, Sn atoms in the
Sn droplet and Si atoms in the a-Si:H layer can diffuse
bidirectionally without obstacles, reducing the effective
diffusion of Si atoms. Moreover, under the inverted mode,
the contact area between the Sn droplet and the a-Si:H is
smaller, which leads to the failure of NW growth under the
inverted mode.

As for the interaction between Sn and a-Ge:H, since the
solubility of Ge in Sn is relatively large, the interaction mainly
depends on the wetting effect. Our previous work¹⁹ has
conclusively demonstrated that the wetting interaction
between Sn and a-Ge:H is significantly stronger compared to
that between Sn and a-Si:H. In the standard growth mode,
before the droplet reaches saturation, due to the strong wetting
effect, the Sn droplet loses its droplet shape and cannot
precipitate NWs. However, under the inverted mode, only the
bottom of the Sn droplet is in contact with the a-Ge:H.
Affected by the initial wetting effect, the Sn droplet expands
slightly at the bottom, but during the growth process, the
wetting is mainly downward and does not expand infinitely in
the horizontal direction. The Sn droplet can still maintain its
droplet shape, and NW growth can be achieved. Using Sn
dioxide particles as the catalyst precursor can weaken the
influence of the wetting effect, allowing NW growth under the
standard mode, but the activation rate is very low. In contrast,
the activation rate under the inverted mode is relatively higher.

4. CONCLUSIONS

In summary, we have comprehensively investigated the
influence of the material properties and growth configurations
on the formation and morphology of NWs. Our findings shed
light on various crucial factors governing NW growth and offer
valuable insights into tailoring nanostructure properties for
specific applications. The Sn metal droplet and amorphous
layer interactions lead to distinct NW growth behaviors. While
a-Ge:H enables NW formation in both standard and inverted
modes, a-Si:H only supports standard growth, highlighting the
critical role of the material properties. Multilayer config-
urations influence NW growth patterns; in particular, in the
inverted bilayer configuration, diverse nanostructures (NWs,
NBs, NCs, and nanopillars) form with a strong correlation
between morphology and catalyst particle size. Higher layer
numbers impact the NW diameter, length, and activation rate.

Overall, this work demonstrates the significant role of
material types (a-Si:H and/or a-Ge:H), growth configurations,
and annealing conditions in tailoring NW structure and
density. These insights pave the way for precise control over
nanostructure properties for specific applications including
neuron attachment, biocompatible interfaces, and tailored
optoelectronic devices.

■ AUTHOR INFORMATION

Corresponding Author

Wanghua Chen — School of Physical Science and Technology,
Ningbo University, Ningbo 315211, China; ◎ orcid.org/
0000-0001-6167-3178; Email: chenwanghua@nbu.edu.cn

Authors

Xiaoyang Zhang — School of Physical Science and Technology,
Ningbo University, Ningbo 315211, China

Xuechun Zhao – School of Physical Science and Technology,
Ningbo University, Ningbo 315211, China

Lulu Zheng – School of Physical Science and Technology,
Ningbo University, Ningbo 315211, China

Ileana Florea – LPICM, CNRS, Ecole Polytechnique, Institut
Polytechnique de Paris, Palaiseau 91128, France; Université
Cote d'Azur, CNRS, CRHEA, Sophia Antipolis Cedex
06905, France

Fabien Cuvilly — GPM, Université Rouen Normandie, INSA
de Rouen, CNRS, Saint Etienne du Rouvray 76801, France

Etienne Talbot – GPM, Université Rouen Normandie, INSA
de Rouen, CNRS, Saint Etienne du Rouvray 76801, France;
◎ orcid.org/0000-0001-6389-1670

Pere Roca i Cabarrocas – LPICM, CNRS, Ecole
Polytechnique, Institut Polytechnique de Paris, Palaiseau
91128, France; ◎ orcid.org/0000-0003-2241-2762

Complete contact information is available at:
https://pubs.acs.org/10.1021/acs.cgd.5c00423

■ ACKNOWLEDGMENTS

W.C. is supported by the Zhejiang Provincial Natural Science
Foundation of China (Grant LY24F040001). TEM and EDS
characterizations were performed in the framework of the
French Government program of investment for the future
(Programme d'Investissement d'Avenir—ANR-IEED-002-01
and TEMPOS Equipex—ANR-10-EQPX-50, pole NanoTEM).

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