Nanoparticle-Templated Thickness Controlled Growth, Thermal

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Nanoparticle-Templated Thickness Controlled Growth, Thermal Stability and Decomposition of Ultrathin Tin Sulfide Plates Eli Sutter, Jia Wang, and Peter Sutter Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 05 Mar 2019 Downloaded from http://pubs.acs.org on March 5, 2019

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Chemistry of Materials

Nanoparticle-Templated Thickness Controlled Growth, Thermal Stability and Decomposition of Ultrathin Tin Sulfide Plates Eli Sutter,1,* Jia Wang, and Peter Sutter2 1Department

of Mechanical and Materials Engineering, University of Nebraska-Lincoln, Lincoln, NE 68588 (USA) 2Department of Electrical and Computer Engineering, University of Nebraska-Lincoln, Lincoln, NE 68588 (USA) ABSTRACT

Tin monosulfide (SnS) is a layered two-dimensional semiconductor attractive for instance for photovoltaics, optoelectronics, and valleytronics. However, these applications require synthesis of SnS flakes with controlled thickness and shape. Here we demonstrate the possibility of using Au nanoparticles to seed the nucleation, determine the position and control the thickness of the growing SnS plates. Further, we use in-situ transmission electron microscopy on individual ultrathin SnS flakes during annealing at high temperatures to establish their thermal stability and decomposition pathways, important due to the large surface areas of ultrathin flakes. We find that ultrathin SnS plates decompose rapidly at temperatures above 400ºC. The decomposition invariably starts at the edges of the flakes and is highly anisotropic. Real-time observations show that while the {010} and the {110} facets are relatively stable and retract slowly, rapid removal of material along the directions leads to the fast decay of {100} facets, which are thus replaced by extended {110} facets. The comparison of shapes developed during growth and sublimation establishes that the typical crystal habits observed for SnS and related group IV monochalcogenides are indeed kinetic growth shapes rather than equilibrium (Wulff) shapes determined by different facet surface free energies. These findings can further guide the design of growth experiments for achieving particular shapes of flakes of SnS and related layered crystals.

*Corresponding author: [email protected]

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1. INTRODUCTION Tin sulfide (SnS) is a member of the relatively unexplored layered twodimensional (2D) group VI monochalcogenide semiconductors (MX, where M: Sn, Ge; X: S, Se) that are isostructural to black phosphorus,1-2 stable under ambient conditions, and are particularly attractive for next-generation low-dimensional opto-electronics3-5 benefiting from carrier confinement and other unique characteristics, such as anisotropic crystal structure and properties,6 large exciton binding energies7-8 and strong light-matter interactions. SnS has attracted interest for light harvesting due to its suitable bandgap9 and large absorption coefficient (> 104 cm-1), low-cost, non-toxicity as well as anisotropic electronic,6 optical,10 thermal and mechanical properties.11-12 To access the extraordinary properties predicted for mono- and few layer SnS for applications,13-14 the controlled synthesis of SnS with particular thickness and shapes needs to be developed. The growth of SnS plates by physical vapor deposition or vapor transport uses stoichiometric SnS powder10, 15-16 or a combination of SnO/SnO2 and S17-18 as source materials and yields different types of SnS flakes depending on the substrate used, i.e., basal-plane oriented layered SnS crystals on flat and relatively inert supports such as graphite, graphene19 or mica,10, 15-16, 20 and vertical/standing SnS flakes on more reactive supports such as Si or SiO2.19,

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Standing SnS flakes integrated in devices demonstrate very high response

speeds for fast photodetection, several orders of magnitude higher than horizontal geometries.21-22 Devices such as IR photodetectors incorporating SnS sheets show excellent external quantum efficiency, which can be further enhanced by decoration with Au particles.10

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SnS growth is usually not very uniform. The SnS flakes nucleate either at step edges of van der Waals substrates or randomly on flat substrate terraces, are polydisperse, have a variety of thicknesses (typically multilayers exceeding 10-15 nm), and even on flat substrates such as mica are not all horizontal.10 Interestingly, the SnS flakes invariably exhibit faceted shapes bounded by {110} side facets rather than the armchair and zigzag edges,10,

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and the same types of faceted crystal shapes are observed for the

isostructural and chemically similar SnSe.8 Crystal facets with a lower surface energy are expected to be part of the equilibrium (Wulff) shape that minimizes the total surface energy of a crystal. For SnS the lowest energy facets are predicted to be the (100) and (110).23 Limited experiments have used Si substrates covered with Au films or nanoparticles and subsequently exposed to Sn-S precursors with the aim of growing SnS nanowires,24-26 similar to the growth of GeS nanowires27 and plates28 demonstrated recently. Even though such nanowires grow as disordered forests with generally low uniformity and poor size control, their formation suggests that Au nanoparticles could be used to seed the growth of SnS nanostructures. In the present work we address two important aspects of the growth of few-layer SnS: (i) Seeded nucleation with thickness control; and (ii) the origin of the strongly faceted shapes of SnS flakes. We investigate the growth of SnS plates catalyzed by Au nanoparticles on Si substrates and demonstrate that the Au nanoparticles play the role of nucleation sites and support the formation of predominantly standing SnS plates. Importantly, we find that the thickness of the flakes is governed by the diameter of the Au nanoparticles, i.e., both the positioning and thickness of the SnS flakes can be controlled deterministically by the placement and size of suitable Au seed nanoparticles.

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Further, we use real-time observations by transmission electron microscopy (TEM) to establish the thermal stability and decomposition pathways of SnS flakes during annealing. Fundamentally, as the inverse of crystal growth, thermal sublimation can shed light on atomistic synthesis processes, e.g., help distinguish between equilibrium (Wulff) shapes29 and kinetic growth shapes,30 which are determined by anisotropic speeds of growth or sublimation instead of the specific surface free energy of different facets so that faster growing (or shrinking) orientations will “grow out” and eventually cease to exist, leaving the crystal shape to be bounded by its more slowly growing (shrinking) orientations.31 Annealing of SnS films has been discussed in the context of sintering of absorber layers in solar cell devices, where it results in improved crystallinity, larger grain size, and enhanced optical absorption.32-34 However, annealing experiments have not been carried out for individual SnS flakes, in contrast to exfoliated few layer black phosphorus for which the decomposition by sublimation has been analyzed previously.35-38 Here we follow the time-evolution of individual SnS flakes using in-situ transmission electron microscopy (TEM) during annealing at temperatures up to 400ºC and establish their decomposition pathways. Our results show that the decomposition is highly anisotropic and results in the formation of extended {110} facets. This family of facets coincides with the bounding facets typically observed for SnS growth and thus rationalizes these ubiquitous crystal habits as kinetic growth shapes of anisotropic group IV monochalcogenides. These findings can further guide the design of growth experiments for achieving particular shapes of flakes of SnS and related layered crystals.

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2. RESULTS AND DISCUSSION 2a. Au-nanoparticle seeded thickness-controlled growth of SnS plates Figure 1 shows scanning electron microscopy (SEM) images of the characteristic morphologies obtained during vapor deposition of SnS on thermally dewetted Au thin films with nominal thickness of 1-2 nm on Si (100). Initially small SnS plates nucleate sparsely on the Au nanoparticles formed on the Si surface by the dewetting of the Au film (Figure 1 (a)). The plates grow vertically and exhibit irregular shapes that are typically wide at the base and narrower at the tip. A particle with bright contrast can be seen at the tip of each plate, indicating that that the nucleation of the plates is associated with individual Au-rich nanoparticles. Continuing the deposition leads to nucleation and growth of more plates (Figure 1 (b)), until the entire substrate surface is ultimately covered with mostly square plates with significant thickness (up to 200 nm, Figure 1 (c)). Importantly, even at this late stage each of the plates contains a single Au-rich particle with bright contrast. X-ray diffraction confirms that the plates consist of crystalline orthorhombic SnS (Figure 1 (d)).39 The presence of a single Au-rich nanoparticle on each individual SnS plate, initially at the tip (Figures 1 (a – inset)) and later embedded in each plate (Figure 1 (c)), suggests that the nucleation proceeds via a vapor-liquid-solid (VLS) mechanism. In the VLS process, Au nanoparticles play the role of growth seeds and transport media whose interface to a growing crystal (usually a 3D crystalline40-42 or layered24-27, semiconductor nanowire) represents the growth front.

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Figure 1. SnS plates grown on a dewetted Au film on Si(100). (a)-(c) SEM images of the SnS plates grown for different time duration: 5, 10 and 20 minutes, respectively. (d) X-ray diffraction pattern of the sample in (c) identified as orthorhombic SnS. Scale marker for the inset in (a): 100 nm.

Here, the Au nanoparticles act as nucleation sites for SnS plates (Figure 1 (a)), but in contrast to nanowire growth44 they do not completely define the footprint of the emerging SnS nanostructure. After the initial nucleation on the Au-rich nanoparticles the SnS flakes expand laterally due to the enhanced reactivity of the exposed edges. Following this fast initial lateral expansion, thickening by nucleation of additional SnS layers ultimately competes with lateral growth and the flakes transform into thick plates that form a complete polycrystalline film (Figure 1 (c)). The observation that SnS plates nucleate on individual Au-rich nanoparticle growth seeds suggests a way for obtaining sparsely nucleated plates whose placement on the substrate is determined by isolated Au nanoparticles. In order to verify this scenario, we carried out growth using extremely dilute solutions of Au nanoparticles, atomized onto Si substrates to provide sparse arrays of isolated Au nanoparticles. Au nanoparticles

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with different sizes were used in these experiments. SEM images of the plates grown under identical conditions (i.e., in the same growth run) on substrates decorated with 5 nm, 20 nm, and 60 nm Au nanoparticles are shown in Figure 2.

Figure 2. SnS plates grown on size-selected Au nanoparticles on Si(100). (a)-(c) SEM images of the SnS plates seeded by Au nanoparticles with diameters of 5 nm, 20 nm and 60 nm, respectively. (d) High-magnification SEM images of individual SnS plates nucleated on 5 nm (top) and 60 nm Au nanoparticles (bottom).

The SEM images show the formation of isolated flakes, which again stand vertically at different angles on the substrate. The Au nanoparticles are clearly visible in the images, in particular for flakes grown on 60 nm Au nanoparticles (Figure 2 (c)). Higher magnification SEM images show that the flakes seeded and templated by the smallest Au nanoparticles also have one Au nanoparticle associated with each of them (Figure 2 (d) top). This suggests that the Au nanoparticles indeed act as nucleation sites for the SnS and the their VLS-nucleated growth produces isolated standing nanoflakes. The size of the flakes and their thickness can be correlated to the size of the templating Au nanoparticles. The flakes grown on substrates with 5 nm Au nanoparticles are the thinnest and largest, while those templated by 60 nm Au nanoparticles are significantly thicker (Figure 2 (d), bottom). This templating effect is reflected in histograms of the lateral size and thickness for SnS flakes grown on 5 nm and 20 nm Au nanoparticle seeds, as measured from the SEM images (shown in Figure S1). Thus, similar to VLS nanowire growth, where the size of the nanoparticle catalysts controls the diameter of the growing

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wires,45-46 the Au nanoparticles here govern the thickness of the flake. Evidently, the thickness control via the Au nanoparticle diameter is optimal at the early growth stages. Extending the deposition time ultimately leads to an increase in both the thickness and the size of the flakes (Figure S2), due to onset of thickening competing with the lateral growth of the flakes.19 To further investigate the morphology and corroborate the inferred growth process, SnS flakes grown on individual 5 nm Au nanoparticles were transferred to lacey carbon films on TEM grids. A survey TEM image of several SnS flakes on the carbon support is shown in Figure 3 (a). Figure 3 (b) shows a HAADF-STEM image of an individual SnS flake. The flake appears uniform except for the clearly visible Au-rich nanoparticle, which shows brighter contrast as expected due to its higher average atomic number, Z. High-resolution TEM images (Figure 3 (c)) confirm that the flake is singlecrystalline SnS (orthorhombic, space group Pnma) imaged along the [001] zone axis (Figure 3 (d)). Note that we use a notation in which the individual SnS sheets are

Figure 3. TEM analysis of Au-nanoparticle seeded SnS flakes. (a) Overview TEM image of SnS plates dispersed on carbon film. (b) HAADF-STEM image of a flake. (c) High-resolution TEM image of the SnS flake shown in (b). (d) FFT of the SnS flake. (e) FFT of the Au-rich particle. (f) Schematic of the Au nanoparticle (NP) seeded nucleation and growth of SnS flakes: (i) Au seed nanoparticles (diameter d1); (ii) uptake of source material from the vapor at the growth

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temperature, causing a diameter increase to d2 governed by the nanoscale phase diagram;46-47 (iii) nucleation of a SnS flake with thickness d2, (iv) preferential lateral growth by SnS incorporation into the flake edges maintaining constant thickness d2.

spanned by unit vectors a and b, and are stacked with weak interlayer interaction along the c-axis (i.e., the [001] zone axis is parallel to c, see also Figure 4 (e) below). The Aurich nanoparticle is clearly larger than the original 5 nm Au seed particle (Figure 3 (c)). This is consistent with a VLS-like process, in which the Au nanoparticle catalyst transforms into a eutectic alloy melt by absorbing source material (here Sn(S)) from the vapor phase at the growth temperature. The lattice fringes of the nanoparticle observed in the high-resolution TEM image (Figure 3 (c)) and their fast-Fourier transform (FFT) analysis establish that its structure at room temperature corresponds to that of an AuSn alloy (Figure 3 (e)), whereas the adjacent material shows the characteristic structure of orthorhombic SnS (Figure 3 (d)). Based on the combined analysis shown in Figs. 1-3, we can reconstruct the pathway of Au nanoparticle seeded nucleation and growth of thickness controlled SnS flakes (Figure 3 (f)). Similar to VLS nanowire growth, the Au seed particle (with starting diameter d1, (i)) takes in source material when exposed to vapor at the growth temperature. In the present case, this causes the formation of a liquid Au-Sn(S) alloy melt drop with enlarged diameter d2, which is governed by the solubility of Sn in the Au-rich melt (ii).46-47 Once a sufficient supersaturation is established, a solid SnS flake nucleates on the drop (iii). In contrast to VLS nanowire growth, our results show that the further growth no longer involves transport through the melt but occurs primarily by direct incorporation of SnS vapor into the open (reactive) edges of the SnS nucleus (iv), which leads to a rapid lateral expansion. At this early stage, the nanoflake thickness is thus dictated by the size of the Au-rich seed drop (diameter d2). At later stages, beyond ~1 μm 9



lateral size, the addition of layers along the c-direction causes a progressive thickening of the flakes that blurs the initial thickness selection by the seed particles, even though plates grown on small particles tend to still be thinnest. To further assess the SnS flakes, they were characterized by Raman spectroscopy, photoluminescence and cathodoluminescence in scanning transmission electron microscopy (STEM-CL) (Figure S3). Typical -Raman spectra (Figures S3 (a) and S4) show peaks corresponding to the Ag and B3g Raman modes of single crystalline SnS,15, 19 without any detectable secondary phases such as SnS2 whose most intense (A1g; A1+E) optical phonon mode would lie at ~310-315 cm-1.48 The typical room temperature PL spectrum of the SnS plates (Figure S3 (b)) is dominated by a peak centered at a photon energy hν = 1.27 eV, corresponding to the fundamental bandgap of SnS along with two higher energy transitions at 1.65 eV and 1.75 eV, similar to previous reports for highquality SnS flakes.19,

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Comparison with the computed SnS band structure and

with absorption measurements19, 50-53 leads us to assign the lower energy peaks to radiative recombination across the indirect and direct bandgaps of SnS. Panchromatic STEM-CL maps of individual SnS plates (Figure S3 (c)-(d)) with spatial resolution below the diffraction limit show significant luminescence from the plates with pronounced maximum in emission intensity observed near the surface of the plates. 2b. Annealing and thermal decomposition of SnS flakes: Equilibrium versus growth shapes In order to investigate the thermal stability and decomposition pathways of thin SnS we carried out in-situ TEM observations of individual thin SnS flakes grown on 5

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nm Au nanoparticles (Figures 4 and 5) during annealing to high temperatures. TEM and HAADF-STEM images of a characteristic flake are shown in Figure 4 (a) and (b), respectively. The electron diffraction pattern confirms that the flake is single-crystalline SnS oriented along the [001] zone axis, i.e., exposing the (001) basal plane. A structural analysis combining electron diffraction with high-resolution TEM imaging (Figure 4 (c), (d)) identifies the main facets defining the shape of the SnS plate as low-index (100), (010) and {110} facets, which are perpendicular to the a- and b-axes and the basal-plane diagonals, respectively. In the case of the SnS flakes templated by Au nanoparticles one of the edges of the flake (unassigned in Figure 4 (b)) can be identified as the interface to the substrate supporting the flake during its standing growth. This edge is easily recognizable as it is usually rough. The side facets we find here are similar to the growth shapes of thin SnS plates, which usually exhibit extended {110} surfaces/edges,15 even in the case of growth of flat basal-plane oriented layered 2D SnS crystals on inert supports such as mica (see Figure S5).10, 15-16, 20 Ultrathin SnS plates are found to assume rounded shapes bounded by vicinal sections and smaller straight facets.6, 16

Figure 4. SnS flake prior to in-situ annealing in the TEM. (a) HAADF-STEM image of a characteristic SnS flake. (b)-(c) TEM image and electron diffraction pattern of the SnS flake. (d)

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High-resolution TEM image of the area of the SnS flake marked with rectangle in (b) viewed along the [001] zone axis. The inset in (d) shows a multi-slice image simulation. (e) Ball and stick model of the SnS crystal structure illustrating the (a, b, c) basis vector notation adopted here.

The SnS flakes were initially heated to 400ºC (Figure 5 (b)-(d)). Their side facets, which are originally straight (Figure 5 (a)) develop waviness almost immediately after reaching 400ºC, which suggests the onset of decomposition (Figure 5 (b)). At 400ºC the decomposition proceeds rapidly. Over ~30 s (Figure 5 (b)-(d)) the waviness increases and at several places along the edges of the facets (some of which are marked with dashed lines in Figure 5 (d)) removal of significant amount of material is observed along particular crystallographic directions. To allow decomposition to proceed more slowly and be followed in real time over tens of minutes rather than seconds, after ~1 minute at 400C the temperature was lowered to 380C. The SnS flakes were followed in real time during further annealing at 380C for about 30 minutes after which they were cooled down to room temperature. Below we summarize the distinct features of the decomposition process as established from the TEM observations of the SnS flakes at 380C.

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Figure 5. In-situ TEM of thin SnS flakes during annealing. (a), (a)’ Initial TEM image and electron diffraction pattern of a SnS flake at 200ºC. (b)-(d) Sequence of TEM images of the SnS flake during annealing at 400ºC over a period of 33 s. (e)-(f) Sequence of TEM images of the SnS flake during continued annealing at 380ºC over a period of 25 min. (g) The SnS flake after cooling to room temperature.

The decomposition by loss of SnS invariably starts at the edges (i.e., lowcoordination sites) of the plates and proceeds from the edges inward. The opening of holes in the basal planes is never observed. The removal of material starting at the edges proceeds in two distinct ways. In thin flakes (Figure 5) material is removed from the edges inwards uniformly over the entire thickness (i.e., involves all the layers of the flake as one entity), i.e., the projected area decreases in size but thinning of the flake via removal of consecutive layers does not occur. In-situ TEM observations for thicker flakes (Figures S6-S8) demonstrate richer and more complex microscopic decomposition pathways, which are however consistent with the overall behavior found for thin flakes. In thicker flakes, local thinning by removal of a part of the layers occurs within small

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areas adjacent to the edges as evidenced by the lighter contrast that develops over time (Figures S6-S8) while the overall thickness in the flake interior is preserved. Also, the elimination of fast-moving facets clearly involves the formation of rounded (vicinal) segments (Figure S7). Independent of thickness, the removal of material along the interlayer c-direction, i.e., sublimation of the basal planes, is never observed. This is distinctly different from the decomposition pathways observed in few layer black phosphorus flakes which invariably initiate on the flat (001) facets36 and proceed via expansion of holes in the basal planes.35, 38 SnS decomposes by simultaneous loss of Sn and S in equal amounts,54 similar to the sublimation of GeS.27,

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Indeed, as in earlier experiments,49,

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we do not observe

decomposition and loss of S but a behavior consistent with the detachment and removal of SnS formula units during annealing. The detachment of SnS from the flakes during annealing is anisotropic, i.e., the loss of SnS is not uniform from the different facets that make up the edges but proceeds at different speeds along different crystallographic directions. Our in-situ TEM observations during annealing of the SnS flakes (Figures 5 and S6) suggest anisotropic decomposition with different detachment rates along the aand b-axes of the SnS sheets as well as along the directions. The SnS flake in Figure 5 is bounded by several different facets from the {100}, {010} and {110} families. SnS is removed from the {110} facets generally in the direction normal to the facet. After the roughening of the large (110) facet in Figure 5 (a), which occurs during the initial heating of the flake to 400ºC, several small segments of this facet lose SnS and move inwards so that the facet develops rectangular-shaped, concave cut-outs. With time (Figure 5 (c)-(e)) these segments expand both along the facet and perpendicular to it and

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the flake loses SnS by movement of the entire facet slowly inward parallel to itself. The large (200) facet in Figure 5 (a) undergoes drastic changes during annealing and SnS sublimation. In contrast to the {110} facets that lose material and move inwards by parallel translation, the loss of material from the (200) facet is such that it is gradually replaced by a set of two different facets. Already during heating to 400ºC it becomes rough and further annealing causes the appearance of several characteristic triangular cutouts from which SnS is predominantly removed. These segments are bounded by pairs of (110) and (110) facets (Figure 5 (d)). The pathway of decomposition of a (100) facet involves fast anisotropic removal of SnS along the [100] direction, which produces pairs of slow-moving (110) and (110) facets. Further rapid removal of material along the [100] direction causes the increase in size of these {110} facets (Figure 5 (d) – (g)). After ~30 minutes of annealing at 380ºC, large sections of the original (200) facet are replaced by (110) and (110) facets. Similar to the {110} facets that were part of the initial shape of the flake, these newly formed {110} facets appear to be relatively unaffected by annealing, i.e., they move slowly during SnS sublimation. Finally, in contrast to the fast removal of material along the [100] direction, almost no removal is observed along the [010] direction and slow-moving {010} facets are hence preserved even after 30 minutes of annealing at 380ºC (Figure S6).

3. CONCLUSIONS In summary, our in-situ observations of thermal sublimation from SnS flakes show distinct differences in the rate of material removal from different low-index facets. Particularly fast SnS removal is observed for (100) facets, so that these edges are

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progressively replaced by facets that retract more slowly during sublimation. Among these slow-moving facets we identify {010}, {110}, as well as the {001} basal facets. The same types of facets, namely primarily {110} facets rather than the zigzag and armchair edges along the a- and b-axes, also terminate SnS flakes obtained by vapor deposition6, 10, 15-16 and colloidal synthesis.55 Shapes consisting of {110} facets are likely kinetically stabilized, as these are not the lowest-energy facets. The SnS (100) facet, for instance, is predicted to have lower specific surface energy than (110).23 Moreover, SnS flakes frequently show large, slow growing {010} and small (if any) {100} facets (Figure S5). Our in-situ experiments show the fast SnS removal along the [100] direction; if the local expansion during growth is also fastest along the [100] direction, the (100) facets will be missing from the SnS flake shapes despite having the lowest surface energy. Based on a comparison of the shapes developed during growth and sublimation, our findings strongly suggest that the typical crystal habits observed for SnS and related group IV monochalcogenides are indeed growth shapes, determined by a kinetic Wulff construction involving anisotropic growth and decomposition rates rather than equilibrium shapes determined by different facet surface free energies. Finally, it is interesting to compare the decomposition of SnS to that of the isostructural black phosphorus. Our observations show both similarities and differences in the sublimation of SnS compared to P. Similar to the anisotropic detachment of SnS, anisotropic decomposition is reported for black phosphorus, but material removal occurs predominantly along the [010] direction. The decomposition of few layer black phosphorus flakes is invariably initiated on the flat basal (001) facet and proceeds via expansion of holes along the [010] direction, supposedly via removal of individual P

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atoms. The anisotropy of sublimation is determined by the removal of P atoms with one bond to the remaining black phosphorus flake.35 In contrast, in SnS flakes the decomposition starts always at the edges and is fastest along the [100] direction. The difference is likely due to the congruent sublimation of SnS units compared to the atomby-atom removal of P atoms from black phosphorus. Indeed, based on the different coordination of SnS units to the crystal along armchair, zigzag, and mixed {110} type edges (two and three Sn-S bonds, respectively, see Figure S9) one would expect a faster decomposition along [100] than along [010] or [110], as is observed in our in-situ experiments. Beyond this bond-counting argument, the detailed atomistic pathways of the decomposition of faceted SnS flakes could be analyzed by ab-initio calculations in future work.

4. MATERIALS AND METHODS SnS plates were synthesized using SnS powder (99.99%, Sigma Aldrich) in an experimental setup consisting of a pumped quartz tube furnace with two independently controlled temperature zones. The evaporation zone containing a quartz boat with the SnS powder was heated to 600C, while the zone containing the substrate was heated to growth temperatures of 300-500C. For temperatures up to 600C SnS does not undergo any structural phase transformations, i.e., source powders and SnS flakes are expected to be in the a-SnS phase (orthorhombic, Pnma) under the growth and annealing conditions reported here.56 Si (100) covered with 2-4 nm thick Au films deposited by sputtering at room temperature was used as substrate. Such Au films spontaneously dewet into Au nanoparticles at the growth temperature.40 These Au nanoparticles do not have any

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ligands or capping layers on their surface. Alternatively, citrate-capped Au nanoparticles with sizes of 5 nm, 20 nm, and 60 nm were used to seed the growth of the SnS nanoflakes. Dilute solutions with Au nanoparticles were deposited on the Si surface using a nebulizer to ensure a sparse nanoparticle coverage. We found no difference between SnS growth on the two types of particles, dewetted from Au films or colloidal. For comparison, SnS plates were grown on mica substrates using vapor transport from SnS powder. During growth, a carrier gas flow of H2 (2%)/Ar was maintained at 50 standard cubic centimeters per minute (sccm) at a pressure of 20 mTorr. The growth was typically performed for 520 minutes resulting in the formation of nanoplates with different sizes and thicknesses. Characterization of the as-grown SnS plates on the Si substrate was carried out using X-ray diffraction, micro-Raman, photoluminescence (PL) and scanning electron microscopy (SEM). SEM was performed in a FEI Helios Nanolab 660 with a field emission gun at 2 kV. Micro-Raman spectroscopy of the SnS nanoplates was carried out using a Raman microscope (Horiba Xplora plus) with an excitation wavelength of 532 nm and a lateral resolution of ~0.5 μm. Photoluminescence (PL) measurements were performed in the Raman microscope with a 100x objective at excitation wavelength of 532 nm and laser power 0.168 mW. The morphology of the nanoplates was investigated by transmission electron microscopy (TEM) and electron diffraction in an FEI Talos F200X microscope equipped with a Gatan 652 high-temperature sample holder. In-situ annealing experiments in TEM covered the temperature range between room temperature and 450C at pressures of ~5 10-8 torr. The electron irradiation intensity was kept intentionally low (< 0.1 A/cm2) during our observations to prevent any uncontrolled electron beam induced structural 18


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changes. Cathodoluminescence (CL) spectroscopy was performed in STEM mode (STEM-CL) using a Gatan Vulcan CL holder at 200 kV electron energy and incident beam current typically ~300 pA.

5. ACKNOWLEDGEMENTS This work was supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Award No. DE-SC0016343. The authors acknowledge technical assistance by C. Keiser.

6. SUPPORTING INFORMATION DESCRIPTION Supplementary figures: SEM on SnS plates; Raman, photoluminescence and STEM CL of SnS plates; TEM of SnS plates transferred from mica substrates; selected TEM images following the decomposition by sublimation of the SnS flakes; schematic of monolayer SnS showing the different coordination of SnS units (pdf).

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Nanoparticle-Templated Thickness Controlled Growth, Thermal

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