Sb2Te3 growth study reveals: Formation of Nanoscale Charge Carrier

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SbTe growth study reveals: Formation of Nanoscale Charge Carrier Domains Is an Intrinsic Feature Relevant for Electronic Applications Martin Lewin, Lars Mester, Tobias Saltzmann, Seung-Jae Chong, Marvin Kaminski, Benedikt Hauer, Marc Pohlmann, Antonio M. Mio, Matti Wirtssohn, Peter Jost, Matthias Wuttig, Ulrich Simon, and Thomas Taubner ACS Appl. Nano Mater., Just Accepted Manuscript • DOI: 10.1021/acsanm.8b01660 • Publication Date (Web): 05 Nov 2018 Downloaded from http://pubs.acs.org on November 7, 2018

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ACS Applied Nano Materials

Sb2Te3 Growth Study Reveals: Formation of Nanoscale Charge Carrier Domains Is an Intrinsic Feature Relevant for Electronic Applications Martin Lewin1,‡,

Lars Mester1,‡,

Tobias Saltzmann2,‡,

Seung-Jae Chong2,

Marvin Kaminski1, Benedikt Hauer1, Marc Pohlmann1, Antonio M. Mio1, Matti Wirtssohn1, Peter Jost1, Matthias Wuttig1,3, Ulrich Simon2,3 and Thomas Taubner1,3*

1 Institute

2 Institute

3 JARA

of Physics (IA), RWTH Aachen University, 52056 Aachen, Germany.

of Inorganic Chemistry, RWTH Aachen University, 52056 Aachen, Germany

– Fundamentals of Future Information Technology, 52056 Aachen, Germany

Keywords: Antimony telluride, infrared near-field microscopy, domain formation, epitaxial growth, transition metal dichalcogenides, van der Waals materials

ACS Paragon Plus Environment

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Abstract: Sb2Te3 exhibits a plethora of fundamentally relevant electronic phenomena enabling electronic phase change memory cells, thermoelectric devices and threedimensional

topological

insulator

structures.

Thus,

the

controlled

growth

of

nanostructures and thin films with well-defined electronic properties is of uttermost importance. Previously, our group observed symmetric infrared domains in hexagonal Sb2Te3 nanoplatelets from a solvothermal chemical synthesis. The relative optical contrast observed was indirectly linked to the formation of regions with different defect densities (charge carrier concentrations). This raises two major questions, which we answer in this study: Is the domain formation restricted to the specific platelet growth process? No! Do the infrared spectra of both domains really follow a ‘Drude-like’ free charge carrier response? Yes!

By controlling the initial water concentration, we promote the growth of the nanoplatelets in c-direction and tune the morphology from platelet-like to octahedra-like. Although the growth mode changes from spiral growth to layer-by-layer, similar infrared domains are identified using scattering-type scanning near-field optical microscopy (s-SNOM).


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Furthermore, we also reproduced the formation of symmetric infrared domains in thin, high quality crystalline films grown using molecular beam epitaxy (MBE). Normalized infrared near-field spectra of smaller Sb2Te3 nanoparticles reveal a relative shift of the plasma frequency in both domains. These findings demonstrate that the formation of domains with different charge carrier properties is an intrinsic material property of Sb2Te3 and might strongly influence all of its electronic applications.


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Sb2Te3 is attractive for novel electronic applications due to its plethora of electronic phenomena. It holds promise e.g. as thermoelectric material for converting thermal into electric energy1 and as a topological insulator representing a new state of quantum matter2. Furthermore, Sb2Te3 shows two stable phases (crystalline / amorphous) at room temperature with significantly different physical properties (electrical resistance, optical reflectivity), which can be switched reversibly in a non-volatile manner by voltage pulses enabling energy-efficient phase change memory or active photonic applications.3,4 Device performance can be significantly enhanced, if Sb2Te3-GeTe superlattice structures are employed, which presumably do not rely just on a local heating mechanism.5,6,7 Hence, the interplay of the Sb2Te3 growth process, local defects and the resulting local electronic properties needs to be understood on the nm-scale to improve and develop novel nanodevice concepts.8

Natural, nominally undoped Sb2Te3 exhibits a strong p-type conductivity, which was linked by scanning tunneling microscopy (STM) and density functional theory to its intrinsic defect concentration of Sb vacancies and Sb-on-Te antisites,9,10,11 which vary spatially


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beyond a statistical distribution.9 The ex-situ STM analysis, especially on samples obtained from cost-effective processes, such as chemical synthesis, is often hindered by the formation of insulating oxide layers. In this case, infrared spectroscopy of the plasma frequency, which monitors the free charge carrier density, can still yield quantitative information about the electronic properties of Sb2Te3 under ambient conditions from below the surface. Albeit, the spatial resolution is limited by diffraction to approximately half the applied wavelength, which corresponds to several micrometers in the infrared spectral range.

To circumvent this limitation and obtain infrared spectral information on the true nmscale we apply scattering-type scanning near-field optical microscopy (s-SNOM).12 It is based on an atomic force microscope (AFM) with a metal coated tip, which is irradiated with focused light from a tunable infrared laser source. Images are taken by scanning the tip relative to the sample surface and interferometrically detecting amplitude and phase of the demodulated backscattered light. The spatial resolution depends no longer on the applied wavelength, but is given by the strongly confined near-fields at the tip apex


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(approximately 25 nm). Due to the evanescent nature of the fields, conductive nanostructures, e.g. consisting of a crystalline phase change material, can be analyzed even below an insulating capping layer.13 Using s-SNOM, our group observed in previous work the formation of strikingly symmetric infrared domains on Sb2Te3 nanoplatelets with a high aspect ratio grown by a solvothermal synthesis (see also Figure S1a)).14 Due to the large size of the platelets, only a relative spectral response of both domains could be analyzed, as a sufficiently large background-free reference with constant infrared response was not accessible. The corresponding relative infrared spectrum could be attributed to different local defect densities. As neither thin film samples prepared by metalorganic vapor phase epitaxy (MOVPE), sputtering or molecular beam epitaxy (MBE) showed similar features, the formation of the domains with different defect densities was ascribed to the specific bi-directional spiral growth process of the platelets along a single screw dislocation.15

In general, Sb2Te3 crystallizes in the tetradymite structure and is comprised of hexagonal close packed layers of Sb and Te running perpendicular to the trigonal axes


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(see Figure 1a)). Each five of the layers in the order of Te(1)-Sb-Te(2)-Sb-Te(1) form layer stacks (quintuple layer) of approximately 1 nm height, dominated by covalent binding. Between the Te-terminated layer stacks weak van der Waals-like forces prevail, whereas a covalent binding contribution in the inter-layer-stack binding has to be taken into account as well.17,18 Thus, the layered morphology of the Sb2Te3 crystal structure is very similar to the layered van der Waals structure of many transition metal dichalcogenides (TMDCs) like WS2 and WSe2 (see Figure S1). Interestingly, also in hexagonal monolayers of WS2 grown by chemical vapor deposition (CVD) photoluminescence analysis revealed similar symmetric domains with different electronic properties, which were tentatively attributed to different local defect densities (see Figure S1b)).19 Very recently, Cai et al. showed that the etching of multilayered WSe2 grown by CVD yielded a very similar trigonal symmetric structure in the morphology of etched layers (see Figure S1c)).20 Thus, regarding the study by Hauer et al. on Sb2Te3 nanoplatelets two eminent questions arise: i) Does the domain formation originate from the specific chemical growth process of the Sb2Te3 nanoplatelets? ii) Is the infrared


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signature of both individual domains really dominated by a’ Drude-like’ free charge carrier response?

To answer both questions, we modify the growth conditions to allow vertical growth of the platelets and perform a near-field optical study of the domain formation (see Figure 1bc)). Based on our previous work,21 we control the particle morphology by adjusting the amount of initially added water. Thus, we tune the crystal growth from dislocation-driven to a process dominated by layer-by-layer growth. We find that similar infrared domains can still be identified and conclude that the domain formation is not dictated by the growth process of the platelets, but is an intrinsic property of Sb2Te3. To verify this hypothesis, we manage to reproduce the formation of infrared domains also in a thin MBE grown film of Sb2Te3. Reducing the size of the nanoplatelets enables quantitative SNOM spectroscopy of the individual domains, directly proving a dominating ‘Drude-like’ free charge carrier response. Additionally, the normalized spectra reveal an absorption peak at =1250 cm-1 in both domains, which might stem from an additional layer or tip contamination. Our results show that the formation of electronic domains is an intrinsic


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property of Sb2Te3 and therefore can potentially strongly affect all of its electronic applications.

Figure 1. Sb2Te3 growth study a) Crystal structure of Sb2Te3 consisting of three quintuple layers prepared using VESTA.16 Te(1) and Te(2) note different chemical states. b-c) Different growth routes in order to investigate the infrared domain formation in Sb2Te3: b) Solvothermal synthesis from liquid phase allows to tune the morphology of Sb2Te3 nanoparticles from plateletlike to octahedra-like by adjusting the initially added amount of water. c) Molecular beam epitaxy from the vapor phases also allows the domain formation for a low partial pressure of Te in the vapor phase.


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RESULTS AND DISCUSSION

To vary the dominating growth mechanism, we modify the chemical formation process. Hexagonal Sb2Te3 particles are obtained from a solvothermal synthesis (See SI for experimental details and chemical reactions during the synthesis). TeO2 and Sb2O3 are mixed with anhydrous diethylene glycole (DEG), diethylene glycolate (NaDEG) and PVP in a Teflon lined autoclave and heated up to 200 °C for 22 h.21 Now, we promote crystal growth in c-direction by adjusting the initial water concentration during the synthesis systematically between 0.1 and 2 M. On the one hand water can lead to a higher supersaturation by faster dissolving the starting materials, while at the same time water can also dissolve the already formed Sb2Te322. In Figure 2 a-d) SEM micrographs of Sb2Te3 crystallites synthesized with a) 0.1, b) 0.4, c) 1.0 and d) 2.0 M water are depicted. With increasing water concentration the [0001] facets change from a regular hexagonal shape to a regular triangular shape. Simultaneously, [2111], [1211], and [1121] facets become more pronounced due to an increased growth in the crystal c-direction (Figure S2). The final crystal shape equals a truncated octahedra. In Figure 2e) the


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aspect ratio is plotted versus the initially applied concentration of water during the synthesis (see Figure S3). A linear decrease of the aspect ratio from a value of 6.2 for hexagonal particles synthesized with a water concetration of 0.1 M to 2.1 for the octahedral crystallites synthesized with a water concentration of 2 M is observed. However, even at higher water concentrations hexagonal platelets can still be found, presumably due to the long reaction times and the complex forming mechanism.


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Figure 2. Promoting growth of the platelets in c-direction. a)-d) SEM micrographs of Sb2Te3 crystallites synthesized at different water concentrations, a) 0.1 M, b) 0.4 M, c) 1.0 M and d) 2.0 M. e) Statistical distribution of aspect ratios (diameter/thickness) found with the SEM in the modified synthesis for different amounts of water. Above the graph simplified models of the respective crystallite morphologies are depicted. For a better comparison, we chose to investigate only one sample with a medium initial water concentration and a larger spread of different aspect ratios. Such, we were able to investigate on one sample platelet-like as well as octahedra-like nanoparticles. In Figure 3 the domain formation in three nanoparticles with morphologies ranging from platelet-like


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to octahedra-like was imaged using s-SNOM (for parameters see SI). The topography images of the particles with aspect ratios of 5.9, 5.8 and 3.2 are shown in Figure 3a-c). Clearly, 1 nm step edges of the single quintuple layers can be identified in all three images, which can be directly linked to the dominating growth mechanism. The topography image of the platelet with the highest aspect ratio in Figure 3a) reveals the existence of a growth spiral (S1) with a trigonal symmetry around a defect centered on the [0001] facet. In contrast, the topographic image of the octahedral particle with the lowest aspect ratio in Figure 3c) shows that its surface texture consists of enclosed layers with triangular shape (L2) and can be attributed to layer-by-layer growth. This change in growth mechanism can be explained following the Burton-Cabrera-Frank (BCF) theory, which assigns the predominant growth mechanism to the supersaturation of the crystal building blocks in the growth media.23,24 While a low supersaturation leads to a spiral growth propagating via a step edge (scheme in Figure 3d)), an increased supersaturation leads to a more probable aggregation of crystal building blocks on an already existing crystal surfaces, thus leading to a layer–by-layer growth (scheme in Figure 3f)). An even more increased supersaturation would finally lead to a dendrite like growth.24 Water


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influences the dissolution and reduction of the starting material directly, or as OH- ions resulting from the hydrolysation of sodium glycolate (see Figure S5 for a detailed summary of the reactions taking place in the initial stage). At a low water concentration just a small amount of OH- ions is generated, thus leading to a slow dissolution of the starting materials and release of Sb3+ and Te2- species, respectively. According to BCF theory the dominating crystal growth can be tuned from screw dislocation dominated to layer-by-layer growth via the supersaturation of Te2- and Sb3+ and thus the water concentration. However, besides the two cases of growth, also a mixing of these mechanisms occurs (scheme in Figure 3e)). For instance, in the topography image of the platelet with a slightly lower aspect ratio in Figure 3b) a separate self-contained layer (L1) as well as also a spiral (S2) can be identified. Following Sheldrick et al. the dissolution of already formed Sb2Te3 by water under solvothermal conditions has to be taken into account.22 However, the [0001] facets are protected by polymeric PVP chains as this polymer tends to coordinate Te and hinders dissolution by water during the reaction.25 For the [1121] facets dangling bonds are reported by Shi et al. which results from a much weaker stabilization by PVP.26 At high water concentration Sb2Te3 material from these


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unprotected facets will be dissolved while the growth in the crystal c-direction is promoted. This controllable growth in the crystal c-direction is in contrast to the generally assumed growth mechanism of Sb2Te3 which is dominated by the intrinsic anisotropic crystal structure.27

The change in growth directly influences the formation of the infrared domains, as can be seen in the corresponding near-field optical phase 2-2Si images in Figure 3g-i). To visualize the influence, in Figure 3j-l) the infrared domains are schematically simplified and overlaid with the step edges, which can be identified in the topography images. In case of the thin platelet in Figure 3j) the infrared domains extend from the growth spiral across nearly the entire platelet. With the inset of multiple spirals and layer-by-layer growth in Figure 3k) a superposition of multiple trigonal infrared domains with sharp boundaries (x  100 nm) can be identified. With decreasing aspect ratio, the infrared domains become increasingly complex, as can be seen for the octahedra-like particle in Figure 3l). The complexity might arise from overgrown spirals and from the statistical attachment of new layers for layer-by-layer growth in contrast to a strict spiral growth. The


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facets of the nanoparticles (see e.g. ‘FA’ in Figure 3i)) also reveal a strong contrast. Here, the tip can couple in horizontal direction, affecting the measured contrasts by artifacts resulting from the asymmetric tip shape or illumination direction-dependent shadowing. Interestingly, a homogeneous outer frame of approximately 200 - 300 nm (see e.g. ‘FR’ in Figure 3i)) can be observed for all of the particles. Recently, ring-shaped electronic inhomogeneities were reported in polyol-synthesized Bi2Se3 and chemical vapor deposition-grown Sb2Te3 nanocrystals, which could not be attributed to elemental inhomogeneities.28 They concluded according to Hauer et al.14 that different defect densities during growth are causing the infrared inhomogeneities as the ring-shaped pattern was found to change after post-annealing. As the infrared domains do not vanish under the different growth conditions, we hypothesize that the domain formation is an intrinsic property of Sb2Te3 and thus is not limited to chemical syntheses.


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Figure 3. Transition of infrared domain formation in Sb2Te3 nanoparticles at different growth stages from hexagonal to octahedron shape a-c) Topography images of three Sb2Te3 nanoparticles with different morphologies (aspect ratios: diameter/height h, as indicated) obtained from the same modified synthesis run. The 1 nm high topographic steps can be directly linked to the growth process: a) Single growth spiral (S1). b) Growth spiral (S2) and an additional selfcontained layer (L1). c) Single self-contained layer (L2) d-f) Different growth mechanisms: d) Spiral growth, e) mixed growth, f) layer-by-layer growth. g-i) The corresponding near-


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field optical phase images 2-2Si at  = 1250 cm-1 normalized to the surrounding Si substrate. The plateau shows a homogeneous frame (FR), while the contrast of the facets (FA) is likely affected by measurement artefacts. j-l) Scheme of the infrared domains observed optically overlaid with the step edges detected in the topography images (red lines). Scale bar in all images: 1 µm.


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To verify our hypothesis, the domain formation has to be reproducible under well controlled conditions in a physical vapor deposited sample. Hauer et al. did not observe any features in the infrared response of a specific MBE grown nanofilm.14 However, the growth of MBE thin films strongly depends on the deposition parameters.29,30 Hence, we strived for MBE parameters that promote the formation of infrared domains (see Table S1). The specific model by Hauer et al. for the domain formation in the platelets is based on early work analyzing the formation of three-fold and six-fold layer morphologies during thin-film growth of Bi2Te3 using hot-wall-epitaxy.31 The growth of the shorter facets was attributed to a change in the surface stoichiometry because of Te atom desorption (Bi rich surfaces), 31 which is likely to cause also a change in the local defect densities.14 Due to the similar chemical structure the results may be directly transferred to Sb2Te3. In case of the previous MBE sample, which did not show any infrared domains,14 the higher partial pressure of Te in the vapor phase (Te/Sb ~ 10, TTe-cell = 330 °C) could have compensated any spatially inhomogeneous Te atom desorption during growth. Thus, we prepared a new MBE nanofilm with modified growth parameters at a significantly lower partial pressure of Te (Te/Sb ~ 3, TTe-cell = 245 °C).


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In Figure 4a) a topographic overview image of the MBE grown nanofilm is presented. Due to the clean deposition method, the step edges can be easily identified after a polynomial background was subtracted to correct the vertical drift during the timeconsuming large area scan. Hexagonal pyramids with different orientations have evolved, forming valleys at their respective intersections. A high-resolution topography zoom-in of the pyramid marked with a red-dashed rectangle is shown in Figure 4b). In the center of the pyramid a growth spiral is observed. However, in lower positions several starting and ending points of additional independent spirals can be identified (marked by black arrows: I-IV). In fact, the pyramid is formed by several over-grown growth spirals, which also seems to cause a trigonal pyramid structure with truncated edges (nearly hexagonal). In contrast, Hauer et al. observed a trigonal spiral structure with sharp edges of the nanoparticles (Figure S1).14 In Figure 4c) the corresponding overview image of the nearfield optical phase 3 at  = 1070 cm-1 is presented. Clearly, infrared domains, which appear to be centered around the growth pyramids, can be identified also on this unidirectionally (in contrast to the bi-directional growth of the nanoplatelets) grown MBE nano-film. Due to the interaction of several pyramids a very complex optical pattern with


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sharp boundaries (x  100 nm) similar to the investigated octahedral nanocrystal has evolved. A high-resolution near-field optical phase 3 image at  = 1070 cm-1 of the pyramid marked with a red-dashed rectangle is shown in Figure 4d). Additionally, the four independent growth spiral profiles are extracted from the corresponding topography image in Figure 4b) and overlaid with the optical image in Figure 4e). At each starting point of a new spiral, the formation of a new domain pattern can be observed. In contrast to the nanoplatelets, the topographic growth facets are strongly correlated with the infrared domain formation. The infrared domains marked by “B” start at the growth spiral and run parallel to the truncated edges /short facets of the growth pyramid. This observation matches our expectations as according to Ferhat et al. the growth of the shorter facets results from a different surface stoichiometry31, which might cause local variations of the defect densities. Thus, we could yield direct proof that the formation of infrared domains is not restricted to structures obtained from chemical syntheses, but reflects an intrinsic property of Sb2Te3.


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Figure 4. Identification of infrared domains on MBE grown Sb2Te3 nanofilm a) Topography image of a 230 nm thick Sb2Te3 film grown by MBE. Scale bar: 2 µm. b) High resolution zoomed-in topography image of the region marked by a red dashed rectangle in a) Scale bar: 1 µm. c) Overview optical phase image 3 at  = 1070 cm-1 of the corresponding topography image in a). d) High resolution zoomed-in optical phase image of the region marked by a red dashed rectangle in c). e) Scheme of the infrared domains observed optically overlaid with the step edges (marked by green, red and white solid lines) detected in the topography image. f) Schematic top-view (above) of the crystal structure of Sb2Te3. Depending on the crystal orientation the lateral surfaces A and B (side-view below) exhibit different Sb contents, as argued in ref. [14]. We cannot unambiguously link the Te flux to the formation of the infrared domains as several growth parameters have been changed compared to the study by Hauer et al. (see Table S1). Instead of using a Te-terminated Si substrate we deposited the films on


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an Sb-terminated Si(111) substrate, which was proven to yield high-quality films at slightly reduced temperature.32 Thus, further experiments are necessary to separate both contributions. However, we note that several studies report a strong influence of the Te flux on the morphology and the electronic properties. In general, comparable MBE growth studies for Bi2Te3 have noted that growth has to take place in a Te rich regime.29,30 Furthermore, the smoothness of the layer surfaces can be improved by increasing the Te excess in the vapor during growth.29 STM analysis revealed that the defect density in MBE grown Sb2Te3 films can be adjusted by the Te flux rate and even the dominating type of defect can be tuned from antimony vacancies VSb to Sb-on-Te antisites SbTe1.10

It remains questionable, whether regions with different defect densities can evolve with sufficiently sharp boundaries to reproduce such symmetric patterns. In particular, Hauer

et al. evaluated only relative spectra of both regions as due to the large size of the platelets a sufficiently large artefact-free reference region was not accessible. By fitting the relative spectral response at five wavelengths, they only indirectly derived a local shift of the free charge carrier resonance. This raised the question whether the infrared


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response in both regions is really dominated by a ‘Drude-like’ free charge carrier response? The modified synthesis enables us to answer this question as we can now spectroscopically image the domains in platelets with a much smaller diameter, which are surrounded by a sufficiently large artifact-free region of the surrounding Si substrate, with a constant infrared response. In Figure 5a)-c) the topographic image and the corresponding near-field optical amplitude and phase images s2/s2Si of the hexagonal platelet investigated at  =1100 cm-1 are shown. Although the roughness of the surface of this platelet hindered any identification of the 1 nm high step edges, a clear trigonal domain structure is observed in both, amplitude and phase. Spectroscopic data are collected by sequentially imaging the domains at different wavelengths of illumination (see Supporting Information for further details, Figure S9). In Figure 5d-e) the relative near-field amplitude s2B/s2A and phase spectra 2B-2A together with the spectral datapoints obtained by Hauer et al.(blue circles)14 are depicted. In Figure 5f-g) the normalized near-field amplitude s2/s2Si and phase spectra 2-2Si from region A (black) and region B (red) are shown separately. The Finite Dipole Model33 was applied to model


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the

expected

near-field

amplitude

s2((v))/

s2(Si)

and

phase

spectra

2(())-2(Si) (solid lines) in case of a pure free charge carrier response:

(

𝜈p

𝜀(𝜈) = 𝜀∞ 1 ― 𝜈2 + i𝜈𝛾

)

(1)

Literature values have been used for the high frequency offset ∞=29.534 and the carrier damping constant  = 1.8×300 cm-1 14. The plasma frequency p depends on the effective mass m* and the charge carrier density N:

𝜈p =

𝑁𝑒2 1 2𝜋c 𝜀0𝜀∞𝑚 *

(2)

Here, the plasma frequency p was adjusted manually in region A (pA  1454 cm-1) and B (pB  1115 cm-1) to best fit the experimentally observed amplitude and phase spectrum. Clearly, the normalized near-field amplitude and phase spectra of both domains show a broad resonance, which is blue shifted in region A compared to region B and can be well described by a ‘Drude-like’ free charge carrier resonance. Additionally, the detailed normalized spectra reveal also an additional smaller, much more confined resonant contribution at  = 1250 cm-1. In the relative spectra, on which the study by Hauer et al. is based, this peak is concealed (Figure 5 d-e)). In summary, the normalized


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spectra prove a dominating free charge carrier response and a local shift of the plasma frequency p in region A compared to region B.

According to Equation 2, a shift of the plasma frequencyp can result from a change of the effective mass m* or a change of the charge carrier density N. Hence, the symmetric domains could be caused by a local change of the band structure or a change of the local defect density, respectively. Ni et al. observed using s-SNOM infrared patterns, resulting from a local change of the band structure in form of a Moiré superlattice at structural defects in layered Graphene/hBN samples.35 However, to yield the infrared domains as observed in Sb2Te3 a complex triangular pattern of extended structural defects would be necessary, which does not appear to be very likely. Furthermore, transmission electron microscopy analysis of the MBE grown nanofilm did not yield evidence for stacking faults (see Figure S8). Recently the formation of similar shaped highly symmetric domains with a strong contrast in the photoluminescence signal in hexagonal monolayers of WS2 (see Figure S1b)) was attributed to different local electronic properties (defect densities).19 After controlled local laser exposure the formation (possibly due to oxidation) of small


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insulating triangles, whose density was strongly increased in only one of the two domains, was observed. This finding imposes an independent chemical evidence for different defect densities in both regions as defective regions should show a higher chemical reactivity. Possibly due to the same reason, Cai et al. observed the formation of similar triangular patterns in the topography of multilayered WSe2, which was grown in a chemical vapor deposition process and subsequently underwent an etching process (see Figure S1c)).20 Here, different defect densities might have caused different etching rates. For the charge carrier densities in region A and B in the Sb2Te3 nanoplatelet in Figure 5 we can derive values of NA  17×1019 cm-3 and NB  10×1019 cm-3, assuming a constant effective mass of m*=0.25 m034. The obtained charge carrier densities are of similar magnitude in comparison to spatially unresolved infrared studies on Sb2Te3 single crystals (7×1019 cm-3)36 or combined infrared far-field and local transport studies on other solvothermal colloidal nanoplates (2-3×1019 cm-3)37. The differences in the absolute defect density level might result from differences in the literature values for the effective mass (both studies used m*=0.1m0 instead). STM analysis of MBE grown thin films revealed p-type intrinsic defect densities in the range of 3×1019 cm-3.9 As the spectral


27


Page 28 of 46

response was found also to vary from platelet to platelet (not shown) the carrier density can be considered very comparable to the defect densities from literature14,36,37. The damping of the charge carriers of  = 540 cm-1 was adopted from Hauer et al.14 and is higher than comparable spatially unresolved bulk values ( = 200-300 cm-1).34,38 However, the increase in damping is reasonable, as recent local transport studies of single solvothermal colloidal nanoplates reported much lower mobilities µ, respectively higher damping values ( = e/(m*2cµ)  1900 cm-1).37 In conclusion, the formation of symmetric regions with different defect densities is the most reasonable explanation for the intrinsic segmentation of the plasma frequency observed here.

The additional peak around  = 1250 cm-1 could in principle be related to interband transitions observed for Sb2Te3 in infrared reflectivity spectra between 1250 cm-1 and 2000 cm-1.39 As these transitions were reported to show a broad peak in far-field spectra (several 100 cm-1) it appears unreasonable to link them to the sharp (several 10 cm-1) features observed in the near-field spectra. Very recently Schmidt et al. observed using s-SNOM inter-subband transitions due to vertical confinement in few-layer WSe2.40 The


28

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transitions were identified by a localized absorption around  = 1300 cm-1. While the spectral position corresponds very well to the peak observed on the Sb2Te3 nanoplatelet at  = 1250 cm-1, no such confinement effect should be observable for the comparably thick Sb2Te3 platelet consisting of many tens of layers. Furthermore, the absorption strength due to inter-subband transitions should scale linearly with the charge carrier density N, according to Schmidt et al., whereas the peak strength in our case is found to be similar in both domains with different N. As the peak was found to not vary locally on the sample (see Figure S10), it could also stem from a homogeneous additional layer on top of the nanoplatelet, e.g. from synthesis residues or from an oxide layer which might have formed under ambient conditions41. Another potential explanation for a peak close to  = 1250 cm-1 are tip contaminations from the PDMS storage boxes, which have recently been observed in synchrotron-based s-SNOM spectroscopic experiments.42,43


29


Page 30 of 46

Figure 5. Normalized near-field spectra proof local shift of the plasma frequency a) Topography image of a small hexagonal Sb2Te3 platelet (height h) from the modified synthesis. Scale bar: 1 µm. b-c) Corresponding near-field optical amplitude s2/s2Si and phase 2-2Si image at  = 1100 cm-1. Scale bar: 1µm. d-e): Infrared near-field amplitude and phase spectra of region B relative to region A (green squares). Additionally, the spectral data points obtained by Hauer et al.14 are included (blue circles). f-g): Normalized Infrared near-field amplitude and phase spectra of region A (black) and B (red) relative to the surrounding Si substrate. The near-field response of free charge carriers in Sb2Te3 was modeled (lines) in d-g) with the Finite Dipole Model33 and the plasma frequencies p were adjusted to best fit the normalized experimental spectra in region A and B separately (pA  1454 cm-1, NA  17×1019 cm-3) and region B (pB  1115 cm-1, NB  10×1019 cm-3). The normalized spectra reveal an additional resonant contribution at  = 1250 cm-1 (gray shaded).


30

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In summary, we answer the two eminent questions, which arose from the observation of infrared domains in Sb2Te3 nanoplatelets by Hauer et al.14: 1) Is the formation of the domains limited to the applied chemical synthesis? No! 2) Is the infrared signature of both domains separately really dominated by a ‘Drude-like’ free charge carrier response? Yes! By controlling the super saturation of the growth species via the water concentration in the synthesis, we promote the growth of the nanoplatelets in c-direction. Although the growth mode shifts to a layer-by-layer growth resulting in octahedral particle morphologies, we observe similar domains using s-SNOM. Furthermore, we reproduce the formation of infrared domains under low Te excess conditions also on high quality films grown by MBE. Thus, we show that the formation of the symmetric infrared domains is an intrinsic material property of Sb2Te3. Normalized infrared spectra of both domains separately prove a ‘Drude-like’ free charge carrier response with a local shift of the plasma frequency. We gather strong evidence from literature that a binary segmentation of defect densities with sharp boundaries might intrinsically form in several layered V2VI3 compounds and TMDCs, which can strongly influence their electronic applications. Furthermore, the normalized spectra of both domains in the nanoparticle reveal an


31


Page 32 of 46

unknown peak at  = 1250 cm-1, which could be a signature of an additional layer or tip contamination. As the MBE grown samples allow now for a controlled growth and in-situ STM and chemical analysis without disturbing insulating oxide layers or chemical residues, we open up a new route to fully understand the domain formation. In future studies, using STM analysis, the single defects at least at the surface can be unambiguously identified10 and therefore the defect densities can be quantified on the nm-scale9 in different growth stages and subsequently correlated to chemical maps of the atomic composition and the observed infrared domains. Even alternative effects, which could influence the band structure, such as a Moiré pattern32,39 or topological signatures of the growth spiral40 can now be analyzed.


32

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ASSOCIATED CONTENT

Supporting Information. Figure S1:

Observation of trigonal domains and crystal structure for Sb2Te3, WS2 and

WSe2 Section:

Experimental details of the solvothermal synthesis

Figure S2:

Crystal morphology for Sb2Te3 single crystals

Section:

Growth characterization

Figure S3:

Statistical size distribution for the hexagonal platelets and octahedral

particles Section:

Detailed analysis of the chemical reaction process

Figure S4:

Schematic reaction cycle during dissolution and reduction of the starting

materials. Figure S5:

Reactions taking place during dissolution and reduction of starting materials

Figure S6:

X-ray diffraction data of the reaction product for different reaction times

Figure S7:

SEM and EDX analysis of a synthesis intermediate (Te nanowires)

Section:

Parameters of s-SNOM investigation

Table S1:

MBE growth parameters

Section:

MBE growth procedure

Figure S8:

Correlative s-SNOM and TEM characterization of MBE nanofilm

Figure S9:

Sequential spectroscopy: Single wavelength s-SNOM images

Figure S10: Sequential spectroscopy: Two-dimensional homogeneity of s-SNOM spectra AUTHOR INFORMATION

Corresponding Author *E-mail: [email protected]


33


Page 34 of 46

Author Contributions M.L., L.M., T.S., T.T., U.S., P.J. and M.Wu. conceived and designed the experiments. T.S. developed the synthesis and together with S.C. synthesized the nanoparticles. M.K., M.P., M.Wi. and P.J. optimized the MBE growth and prepared the MBE grown nanofilm. A.M.M. performed the TEM analysis. M.L., L.M. and B.H. performed the nearfield experiments. M.L. wrote the manuscript with support from T.S.. P.J., M.Wu., U.S. and T.T. supervised the research. All authors discussed the results and commented on the manuscript. All authors have given approval to the final version of the manuscript. ‡Martin Lewin, Lars Mester and Tobias Saltzmann contributed equally.

ACKNOWLEDGMENT The authors thank Gregor Mussler, Sebastian Peter, Markus Morgenstern and Marcus Liebmann for helpful discussions. Additionally, the authors thank the Ernst RuskaCentre (ER-C) Juelich for the chance to perform high-resolution TEM analyses at their facilities. Funding from the DFG (German Science Foundation) within the collaborative research center SFB 917 “Nanoswitches” is gratefully acknowledged. Furthermore, this


34

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work was financially supported by the Ministry of Innovation, Science, Research and Technology of the German State of North Rhine-Westphalia.

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Chem. Soc. 2018, 140, 10980-10987. (21) Saltzmann, T.; Bornhöfft, M.; Mayer, J.; Simon, U. Shape Without Structure: an Intriguing Formation Mechanism in the Solvothermal Synthesis of the Phase-Change Material Sb2Te3. Angew. Chem., Int. Ed. 2015, 54, 6632–6636. (22)

Sheldrick,

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M.

Solvatothermale

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(42) Grasseschi, D.; Bahamon, D. A.; Maia, F. C. B.; Neto, A. H. C.; Freitas, R. O.; de Matos, C. J. S. Oxygen Impact on the Electronic and Vibrational Properties of Black Phosphorus Probed by Synchrotron Infrared Nanospectroscopy. 2D Mater. 2017, 4, 035028. (43) Freitas, R. O.; Deneke, C.; Maia, F. C. B.; Medeiros, H. G.; Moreno, T.; Dumas, P.; Petroff, Y.; Westfahl, H. Low-Aberration Beamline Optics for Synchrotron Infrared Nanospectroscopy. Opt. Express 2018, 26, 11238-11249. (44) Schouteden, K.; Li, Z.; Chen, T.; Song, F.; Partoens, B.; Van Haesendonck, C.; Park, K. Moiré Superlattices at the Topological Insulator Bi2Te3. Sci. Rep. 2016, 6, 20278. (45) Avdoshenko, S. M.; Koskinen, P.; Sevincli, H.; Popov, A.; Rocha, C. Topological Signatures in the Electronic Structure of Graphene Spirals. Sci. Rep. 2013, 3, 1632.

TABLE OF CONTENTS FIGURE


40

Page 41 of 46 a) Sb Te

Quintuple layer

Te(1) Sb Te(2) Sb Te(1)

c) c a

Sb

Te

b

Vapor

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15

b)

3

Liquid

2


Solvothermal synthesis Sb2O3 TeO2 DEG, PVP, NaDEG

c(H2O): 0.1 M Heated @ 200°C for 22 h c(H2O): 2 M

Molecular beam epitaxy Si(111) substrate Sb Sb termination Partial pressure: Te/Sb ~ 3 ACS Paragon Plus Environment Te Tsubstrate ~ 175 °C

a)

b)

ACS Applied Nano MaterialsPage 42 of 46

Aspect ratio (w/t)

1 2 3 4 1 µm 1 µm 5 c) d) 6 7 8 9 10 11 12 13 2 µm 2 µm 14 15 e) 16 17 10 18 8 19 6 20 4 21 2 22 0 ACS Paragon Plus Environment 23 -2 24 0.0 0.5 1.0 1.5 2.0 25 Water concentration [mol/L]

a) ratio: 5.9 Page 43 ofAspect 46 h = 320 nm 1 2 3 4 S1 5 6 7 Ø = 1900 nm 8 9 d) 10 11 12 13 14 g) 15 16 17 18 19 20 21 22 23 24 25 j)26 27 28 29 30 31 32 33 34 35 36

b)

Δz [nm] 10

Aspect ratio:Nano 5.8 Materialsc) Aspect ratio: 3.2 ACS Applied h = 310 nm h = 570 nm

8 S2

L2

6 4

L1

2 Ø = 1800 nm

Ø = 1800 nm

e)

f)

h)

i)

0

φ2- φ2Si [rad] 0.9 0.6 0.4

FR FA k)

l)


0.2 0.0 -0.2

a) 1 2 3 4 5 6 7 8 9 10 c) 11 12 13 14 15 16 17 18 19 20 21

Δz [a.u.] b) ACS Applied Nano Materials Δz [nm] max 13 I

e)

Page 44 of 46 B

A

10

II

φ3 [rad]

0.1

φ3 [rad]

0.0

-0.1

-0.1

-0.3

B

f)

B

0.1

0.0

-0.2

A

2 0

IV

d)

B

4

III

min

A

8 6


y x

-0.2 -0.3

z x

A

A

B

B

Te(1) Te(2) Sb

A

A

B

Page 45 of a) 46 h = 117 nm

8

d)

6

Ai

4

B

2

1 µm

1100 cm-1

0

12 11 10

9

8

7

1 µm

6

0.4 0.2

9

8

7

6

New data

0.7

Hauer et al.

Region B

1.6 1.4 1.2

g)

Opt. phase ϕ 2 - ϕ Si 2 [rad]

Region A

Hauer et al.

0.1 0.0 −0.1 - 0.2 - 0.3

0.6 1.8

5 New data

0.2

0.8

-0.3

λ [µm] 12 11 10

0.9

0.0

1 µm

0.5

5

Opt. phase ϕ B2 - ϕ A2

Opt. amplitude s B2 /s A2

0.6

1.0

1.0

1.0 800

Φ2-Φ2Si [rad] 0.9

1.5

1.1

f)

c)

e)

λ [µm]

Opt. amplitude s 2 /s Si 2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26

b)ACS Applied Nano Materials s2/s2Si 2.0 Si

∆z [nm] 10

0.5

Region A

0.4

Region B

0.3 0.2 0.1 0.0

ACS Paragon Plus Environment - 0.1 1000

1200

1400

1600

Wavenumber ν [cm- 1 ]

1800

2000

800

1000

1200

1400

1600

Wavenumber ν [cm−1 ]

1800

2000

ACS Applied Nano Materials Page 46 of 46

Amplitude

1 2 Si 3 Sb2Te3ACS Paragon Plus Environment 4 5 -1 800 2000 Wavenumber ν [cm ]

Sb2Te3 growth study reveals: Formation of Nanoscale Charge Carrier

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