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One-dimensional atomic segregation at semiconductor-metal interfaces of polymorphic transition metal dichalcogenide monolayers Ziqian Wang, Min Luo, Shoucong Ning, Yoshikazu Ito, Hamzeh Kashani, Xuanyi Zhang, and Mingwei Chen Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b01839 • Publication Date (Web): 12 Sep 2018 Downloaded from http://pubs.acs.org on September 13, 2018

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One-dimensional atomic segregation at semiconductor-metal interfaces of polymorphic transition metal dichalcogenide monolayers Ziqian Wang,1,† Min Luo,2,† Shoucong Ning,3 Yoshikazu Ito,4 Hamzeh Kashani,1 Xuanyi Zhang,1 and Mingwei Chen1,5,*

1

Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, MD

21218, USA;

2

Department of Physics, Shanghai Second Polytechnic University, Shanghai

201209, P. R. China; 3 Department of Materials Science and Engineering, National University of Singapore, 9 Engineering Drive 1, 117575, Singapore;

4

Institute of Applied Physics, Graduate

School of Pure and Applied Sciences, University of Tsukuba, Tsukuba 305-8573, Japan;

5

Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan.

*Address correspondence to: [email protected] † These authors contribute equally to this work.

ACS Paragon Plus Environment

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ABSTRACT: Interface segregation is a powerful approach to tailor properties of bulk materials by interface engineering. Nevertheless, it is little to know about the chemical inhomogeneity at interfaces of polymorphic two-dimensional transition metal dichalcogenides (TMDs) and its influence on the properties of the 2D materials. Here we report one-dimensional monoatomic segregation at coherent semiconductor-metal 1H/1T interfaces of Mo-doped WS2 monolayers. The monoatomic interface segregation takes place at an intact transition metal plane and is associated with the topological defects caused by symmetry breaking at the 1T/1H interfaces and the weak difference in bonding strength between Mo-S and W-S. This finding enriches our understanding of the interaction between topological defects and impurities in 2D crystals and enlightens a potential approach to manipulate the properties of 2D TMDs by local chemical modification and interface engineering for applications in 2D TMD electronic devices.

KEYWORDS: interface segregation, two-dimensional materials, transition metal dichalcogenide, semiconductor-metal interface, transmission electron microscopy

Interface segregation caused by the interaction of interface defects with chemical dopants and impurities has been widely utilized to manipulate the structure and properties of various heterostructure materials. Notable examples include mechanical property improvement in steels by interface segregations,1 anisotropic grain growth controlling in ceramics by grain boundary segregation,2 and lowering Schottky barrier height in field effect transistors by dopant segregation at silicide-silicon interface,3 etc.. Recently, great attention has been paid to twodimensional (2D) crystals along with their unconventional chemical and physical properties emerging from atomic thickness.4–8 Analogous to the conventional bulk materials, the chemical


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and physical properties of 2D materials are dependent on their chemical compositions.9–15 Chemical heterogeneity has also been observed in 2D materials, such as monolayer MoSe2-WSe2 heterojunctions16 and WSe2-MoS2 lateral p-n junctions.17 However, to the best of our knowledge, significant interface segregation has not been observed in 2D materials up to date. Transition metal dichalcogenides (TMDs) are a unique family among all the 2D materials discovered so far due to their polymorphism even in monolayer thickness. Taking monolayer WS2 and MoS2 as examples, the trigonal prismatic 1H phase is thermodynamically stable and shows semiconducting properties, while the octahedral 1T phase is metastable and exhibits metallic behavior.18–21 Because semiconducting and metallic states can be alternatively obtained, the phase transition and thus interface structure have drawn much attention from both basic research and practical applications. Particularly, the introduction of 1H-to-1T phase transition in a 2D field effect transistor can largely reduce the metal-semiconductor contact resistance for device applications.20,22 The phase transition between 1H and 1T is accomplished by the slide of one S layer in the S-TM-S (TM=transition metals) unit-cell-thick crystals while the transition metal layer remains unchanged and is thus fully coherent at the 1T/1H interfaces. In principle, significant segregation is unlikely at such geometrically defect-free interfaces according to classical interface segregation theories.1,23 In this work, we report atomic-scale observations of one-dimensional segregation of monoatomic layer Mo at the fully coherent 1H/1T interfaces in Mo-doped WS2 monolayers. Different from conventional equilibrium and non-equilibrium interface segregations which are driven by lattice disorder and local elastic strains, the 1D atomic segregation in the 2D TMD arises from topological effect, i.e. the number variation of coordinated S atoms caused by the reflection symmetry breaking at the 1H/1T interfaces. Theoretical calculations suggest that the monoatomic layer segregation can modify the contact


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resistance of the metal-semiconductor interfaces and show potential applications in manipulating the electronic properties of TMD based devices. As schematically shown in Figure 1a, the S atoms closer to the viewer are colored yellow and those S away from the viewer are colored cyan. A 1T phase region embedding in the 1H matrix is formed via the 1⁄2 0110 glide of cyan S atoms which is indicated by the double line arrows. Herein, this transition gives rise to two kinds of interfaces as the result of 3-fold rotational symmetry breaking by the glide of S (Figure 1b and c). The interface with one transition metal atom (Mo or W) coordinating with seven S atoms (7c structure, Figure 1b) has one additional S coordination compared with the six coordinated S atoms in perfect 1H or 1T phase, while each transition metal atom at another interface only coordinates with five S atoms (5c structure, Figure 1c). Such interface structures embedded in the three-atomic-layer crystals enable to study atomic-scale interfacial segregation in an intuitive manner. Mo-doped WS2 monolayer samples were grown on slide glass substrates by chemical vapor deposition (CVD).24 A portion of metastable 1T phase domains were generated in the monolayer 1H phase matrix as a consequence of thermal strains applied to the sample during cooling after CVD growth.25 The formation of 1T phase is also verified by X-ray photoelectron spectroscopy on the basis of binding states of Mo-S and W-S (Figure S1, Supporting Information). The metastable 1T phase domains with 7c and 5c 1H/1T interfaces can be frequently observed in alloyed Mo1-xWxS2 monolayers with various Mo to W ratios by transmission electron microscopy (TEM). Figure 1d and e show the bright-field (BF) and dark-field (DF) TEM images taken from a 1H/1T two-phase region of a Mo0.2W0.8S2 monolayer. Different from the BF-TEM image in which no contrast difference between 1H and 1T phase is observable, the DF-TEM micrograph imaged by the 1100 diffraction beam discriminates the two phases by showing


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higher intensity of 1H and lower intensity of 1T in Figure 1e. Such diffraction contrast in the DFTEM image is associated with the dynamic electron scattering and has been elucidated elsewhere.26 From DF-TEM images, two kinds of 1T/1H interfaces with an about 60º angle can be identified in all 1T phase domains observed in this study, which is consistent with the crystallographic relationship between 7c and 5c interfaces shown in Figure 1a. High-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) images corresponding to 7c and 5c interfaces are shown in Figure 2. The identification of the 7c and 5c interfaces is based on the integrated intensity of excessive solo sulfur background at the interfaces (Figure S2) because the 7c structure has an additional line of solo sulfur atoms at the 1T/1H interface while the 5c structure does not have (Figure 1a). In the low-magnification HAADF-STEM images (Figure 2a and b), both 7c and 5c interfaces show unusual contrast in the form of a monoatomic dark line. The contrast of HAADF-STEM is associated with the Z number (or mass) of imaged elements. Combining with the chemical information from XPS, both atomic structure and chemical species in the TMD lattices can be determined. The atoms with the brightest contrast are the heaviest W atoms (Z=74) while the Mo atoms (Z=42) are dimer.27 For the lightweight S atoms, the overlapping double S atoms in the 1H phase show visible intensity whereas solo S atoms in the 1T phase have a very low intensity,28 as illustrated by the intensity profiles (Figure 2e) along the dashed lines noted in Figure 2c and d. The quantitative intensity ratios of these constituent atoms in 1H, 1T and interfaces are summarized in Table S1. From the atomic images, the interfaces in Figure 2c and d are assigned to be the 7c and 5c structures, respectively (Figure S2), which are the derivatives of γ and α interfaces reported in a previous work.29 From the atomic contrast, it can be recognized that both interfaces are occupied by a monoatomic layer of Mo atoms with a high occupation fraction measured by the Mo fraction


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XMo, i.e. 0.91 for the 7c and 0.80 for the 5c interfaces based on statistics over 800 atoms along both interfaces (Table 1), which are significantly higher than the average Mo fraction of 0.20 in the transition metal layer of the Mo0.2W0.8S2 sample. Note that the intensity of interfacial Mo and neighboring double S is relatively lower than that of the atoms in perfect lattices (Figure 2e and Table S1). However, the intensity ratio of interfacial Mo and double S is close to that derived from the perfect regions, indicating that the weak contrast may be associated with the symmetry breaking at the interfaces and resulting difference in vibration amplitudes of the interfacial atoms as well as possible local lattice bending (Figure S3). The relatively blurry contrast of the interfacial atoms may also be caused by the large interfacial vibration and lattice bending. Note that the Mo segregation at 1H/1T interfaces is not only observed in the W-rich Mo1-xWxS2 monolayers but also in Mo-rich ones, such as Mo0.6W0.4S2, Mo0.5W0.5S2, Mo0.8W0.2S2, as shown in Figure S4. According to our extensive TEM observations, the monolayer Mo segregation takes place in all the 1H/1T interfaces observed in this study and is a universal phenomenon in the alloyed two-phase TMD monolayers regardless of Mo/W ratios. We noticed that the 7c interfaces are relatively straighter and contain fewer steps while, as marked by the yellow squares in Figure 2b, there are a number of atomic steps in the 5c interface. The zoom-in HAADF-STEM image (Figure 2d) shows that the interface step has one atomic shift from the original plane with retained Mo segregation as marked by the double grey dot lines. Detectable lattice defects, such as dislocations, cannot be seen in the regions with rich atomic steps. Moreover, as indicated by the red dashed circles on Figure 2c and d, a small proportion of transition metal vacancies can be detected at the Mo segregated interfaces (Table 1). These vacancies are not likely caused by the knock-on damage of electron beam since the threshold for displacing a Mo atom in MoS2 is ~ 560 keV,30 which is much larger than 200 keV used in our


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study. This is also supported by the fact that the transition metal vacancies at the 1H/1T interfaces have not been observed in the monolithic MoS2 monolayers in this work and previous reports.20,28,29 Thus, it is reasonable to deduce that the interface segregation may accomplish via transition metal vacancies. The Mo/(Mo+W) ratio profiles in the vicinity of 7c and 5c interfaces, measured from over 200 Mo atoms at the interfaces, exhibit sharp Mo-enrichment exactly at the interface sites without observable diffusive layers as shown in Figure 2f. Thus, the interface segregation has a monoatomic thickness and only takes place at the interface layer between 1H and 1T phases. Moreover, the δ-function-like peaks and flat grounds in the Mo-proportion profiles indicate that the interface segregation of Mo in WS2 monolayers almost reaches the equilibrium state.31,32 As shown in Figure 1, the 1H/1T interfaces in monolayer TMDs are fully coherent in crystallography but are topologically defective in terms of the coordination number of nearest neighbored S atoms. Despite the fact that MoS2 and WS2 have almost the same lattice parameters and good miscibility in a single Mo1-xWxS2 phase,27,33,34 the W-S bond has relatively higher bonding energy and covalency compared with the Mo-S bond according to the theoretical calculations by Ref. 35. It can be inferred that the interface energy decreases when Mo atoms segregate at the 7c 1H/1T interface and the excessive metal-S bond is the lower-energy Mo-S, instead of the higher-energy W-S bond. Similarly, for the 5c interface with one deficient metal-S bond which results in the interface energy, replacing the higher energy W-S bonds by the lower energy Mo-S should be energetically favorable when Mo segregates at the interface. Therefore, the difference in bonding energy between Mo-S and W-S bonds, coupling with the topological specialness of the interface sites, could give rise to the monoatomic layer segregation by selective occupation of Mo atoms across the intact transition metal layer. To further understand


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the underlying mechanisms of the anomalous interface segregation in the 2D crystals, we calculated the charge density distributions at 7c and 5c interfaces of monolayer WS2 with and without Mo segregation by using density functional theory (DFT) on the basis of the models shown in Figure 3a-d. For the pure WS2 monolayers, the directional bonds with noticeable spread-out charge distribution can be seen at both 7c and 5c interfaces (Figure 3f,h), revealing the elevated energy due to W-occupation. The interface segregation of Mo enhances the charge density near the Fermi level and lowers the energy of directional bonds with a more localized charge distribution, particularly at 7c interface (Figure 3e,f). Such effect can also be seen in the calculated density of states as discussed in Figure S5. These trends are further supported by the calculations of segregation energies, showing that the Mo segregation at the 7c and 5c interfaces is associated with energy minimization by the coupling between the Mo-S bond and the corresponding topological defects at the 1H/1T interfaces. The calculated total energies of a 1H-1T two-phase monolayer WS2 system with and without Mo segregation reveals that the interface segregation leads to the interface energy reduction of 116 meV and 77 meV per Mo atom for 7c and 5c interfaces in comparison with the corresponding non-segregation systems with the same chemical composition (Figure S6 and Table S2). On the basis of the calculated segregation energies, we can use Langmuir-McLean isotherm: ⁄ − = ⁄1 − exp−Δ ⁄

(1)

to model the equilibrium segregation at a certain temperature T. Herein, notes the Mo (solute atom) fraction occupying the interfacial sites (Table 1), i.e. = 0.91 and 0.83 for the 7c and the 5c interfaces, respectively, and = 0.2 is the overall Mo fraction for the present composition Mo0.2W0.8S2. The segregation energy Δ uses the aforementioned calculated values of 116 meV


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and 77 meV for the 7c and the 5c interfaces, respectively. Temperature dependence of ∆G is considered to be negligibly small as discussed in the Supporting Information. The saturation fraction of solute atoms at the interfaces is taken as 1 because the interfacial transition metal sites are able to be fully occupied by Mo atoms in MoS2. Using these values, the temperature at which equilibrium segregation is reached can be roughly estimated to be around 90 °C and 50 °C for the 7c and the 5c interfaces, respectively, in accordance with the minimum temperature required for 1T to 1H phase transition in MoS2.19 Because the 1T phase is metastable and energetically unfavorable at the CVD temperature of 700°C, the 1T phase and thereby the 1H/1T interfaces and single-atom Mo segregation should be formed during the cooling stage after CVD growth where 1H-to-1T phase transition is assisted by thermal strains. Such behavior can be evidenced by the increase of 1T proportion as a consequence of enlarged thermal strains.25 At the estimated temperature 50~90 °C, the equilibrium of segregation is considered to be reached accompanied with the finish of 1H-to-1T phase transition. From the kinetics point of view, the interface segregation requires transportation of Mo atoms to form Mo segregated interfaces. Classical diffusion theory depicts the atomic diffusion in bulk materials as the exchange of an atom with an adjacent vacancy instead of direct transposition of two adjacent atoms as the consequence of energetic favorableness. However, transition metal vacancies have not been noticed in the 1H and 1T phase regions away from the phase interfaces, indicating that long-range diffusion from the inner 1H or 1T regions to the interfaces should not be the primary mechanism for the observed segregation especially when approaching equilibrium at relatively lower temperature. Such aspect is also supported by the absence of composition gradient in the vicinity of the Mo segregated interfaces (Figure 2f) which is required as the driving force for the long-range diffusion. Instead, transition metal vacancies at the


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segregated interfaces, as observed in Figure 2c and d, may assist short-range atomic transposition closely neighboring to the interfaces. Direct position switch between one Mo and one W atom on the transition metal plane needs simultaneously breaking at least 10 Mo-S or W-S bonds. Whereas, only breaking 4 Mo-S or W-S bonds is required for the position exchange between a Mo or W atom and an interface vacancy, resulting in more than a half-drop of the activation energy, and thus benefits the interface segregation. A possible pathway of such short-range vacancy-assisted interface segregation in the 2D crystals is schematically elucidated in Figure S7. The observed atomic segregation at the 1T/1H interfaces has two unique features: (1) the occupation fraction of solute (Mo) atoms is very high at either 7c or 5c interface; and (2) the segregated area is only monoatomic thick without detectable composition gradient in the vicinity of the interfaces. Such features may lead to the interpretation of the Mo segregated 7c and 5c interface patterns as interface “complexions”,32,36–38 i.e. coupled solute-topological defect complexes at the interfaces. The interface complexions reported in bulk materials31,32,39,40 all result from the interactions between chemical impurities and interfacial defects and are structurally and compositionally different from abutting matrices and initial atomic configurations of equilibrium interfaces. In contrast, the interface complexions in the Mo doped WS2 monolayers remain the original 7c and 5c structures hosted by the abutting 1H and 1T phases in pure WS2 monolayers and are only compositionally different from the initial interfaces. In general, the interfacial energy principally comprises of the strain energy and the energy due to coordination number change, i.e. extra bonds or dangling bonds corresponding to 7c or 5c structure. Since the interfaces are highly coherent and the Mo segregated metal layers do not have any lattice mismatch, the proportion of strain energy could be negligibly small while the


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topological effect from the coordination number disparities at the interfaces should be responsible for the interfacial energy almost exclusively. The topological effect leads to the formation of a monoatomic chain by segregated Mo atoms exactly at the interface sites with less or excess coordination numbers in comparison with the perfect lattice sites of either 1T or 1H phase. Therefore, the formation of the 1D interfacial complexions in the heterostuctured 2D crystals is driven by topological effect solely. The interface segregation phenomenon should be ubiquitous in other 2D alloy systems which have the similar solute-matrix relation in chemical properties as the present circumstance of Mo1xWxS2

monolayers. It is well known that correlated d-electrons from transition metals in TMDs

play a central role in determining the electronic, magnetic, optical and mechanical properties, and electron correlation is very sensitive to d-electron density change especially under confinement in a low-dimensional regime.18,41,42 DFT calculations suggest that the 1D Mo segregation can potentially change the contact resistance between 1H/1T phase in the Mo1-xWxS2 monolayers (Figure S8). We have theoretically investigated the effect of the topological interface segregation on the contact resistance in different TMD systems, including Nb-, Re- and Tasegregated 1H/1T interfaces in chemically doped WS2 monolayers and obvious changes in electronic structure and properties caused by the interface segregation can be seen (Figure S9). Therefore, the monoatomic interface segregation could be utilized as a general approach, similar to the interface engineering in bulk crystals, to tailor the electronic properties of semiconductormetal interfaces in TMD monolayers with abundant choices of transition metals that can be used as solute and matrix. For example, designing interfacial complexions by changing the type of dopant atoms beyond the isoelectronic case might bring about novel properties of the semiconductor-metal interface, such as tunable interface dipole, interface magnetism, and so on.


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In summary, we report 1D monoatomic segregation of Mo at 1H/1T interfaces of Mo-doped WS2 monolayers. Different from the interface segregations in bulk crystals, the significant monoatomic layer segregation in the 2D TMDs is solely driven by topological defects, as a consequence of energetic stabilization by forming interface complexions which are compositionally different from those of the clean 1H/1T interfaces. The influence of the interface segregation on the charge density and energy barrier height of 1H/1T interfaces shows the promise to tailor the metal/semiconductor contact resistance in 2D TMD electronic devices through the interface engineering. This finding shines new insights into the interface segregation phenomenon in 2D materials and may enlighten a new degree of freedom in engineering 2D materials for the realization of novel low-dimensional electronic devices.

METHODS CVD growth of Mo1-xWxS2 monolayers. The Mo-doped WS2 monolayers are synthesized using a low pressure chemical vapor deposition method as shown in the schematic diagram (Figure S10).24 MoO3 (Sigma Aldrich, purity >99.5%, 1~5 mg), WCl6 (Sigma Aldrich, purity 99.9%, 10 mg) and S (Wako Pure Chemical Industry, purity 99%, 1.0 g) powders are used as precursors, and micro slide glass (Matsunami Glass, S1214) as growth substrates. Temperatures of the sources and substrate were controlled separately by heating in different furnace zones. Growth chamber is pumped down to a base pressure of 5×10-2 mbar before growth. Ultra-pure Argon was used as the carrier gas as well as protection gas at low pressure (with pumping) throughout the growth. Sources and the substrate are heated to their designed values (35 °C for WCl6, 160 °C for sulfur, 500 °C for MoO3 and 700 °C for substrate) in 10 min under the Ar flow rate of 2000 sccm (corresponding to 5.1 mbar). Then, Ar flow rate was switched to 500 sccm (2.1 mbar)


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to initiate the supply of all the species to substrate for film growth. The temperatures of the precursor sources of sulfur and MoO3 and the substrate were kept constant for 40 min while the temperature of WCl6 gradually dropped from 35 °C to 20 °C. Finally, in the cooling down stage, WCl6 was rapidly cooled down to 10 °C using ice and the others to room temperature in 5 min using an electric fan to end the CVD growth. Vacuum and Ar flow are kept till the furnace completely cools down to protect samples from oxidization and impurities. Samples with different Mo compositions are controlled by changing the amount of MoO3 precursor.24 Microstructure characterization. As-grown samples are quickly transferred to holey carbon coated copper grids for TEM characterization. Note that as-grown samples can be lifted from glass substrates and float on the surface of small water droplets as releasing their CVD-induced thermal strains. Taking such advantage, only ultra-pure water was used throughout the transfer process, which enables clean transfer of the samples. TEM and STEM images and SAED patterns are captured using a JEOL JEM-2100F TEM equipped with double spherical aberration (Cs) correctors (CEOS) for imaging and probing lenses. Observations are performed at the acceleration voltage of 200 kV and the spatial resolution of the STEM is ~0.1 nm. The collecting angle for HAADF-STEM is between 64 and 171 mrad and a high scan rate of over 870 nm/s is used to acquire HAADF-STEM images. The estimated electron dose is lower than 5×104 e-/Å2 (corresponding to dwelling time of 19.1 µs/pixel, gun current of 50 pA and pixel area of 0.11 Å2). Supplementary data discussing the avoidance of knock-on damages of transition metal atoms are shown in Figure S11. DFT calculations. We performed first-principles calculations based on density functional theory (DFT) within the generalized gradient approximations (GGA-PBE)43 as implemented in the Vienna ab-initio simulation package (VASP).44 The projector augmented-wave (PAW)45


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pseudopotentials are used and the cut off energy is 450 eV. The 4×4×2 k-point grids are used in the 9×4×1 supercell. The optimized lattice constant of WS2 is 3.169 Å. For the atomic models used in the DFT calculations (Figure 3a-d and Figure S6a,b), periodical boundary conditons are used for the boundaries across the phase interfaces, whereas ribon configurations are used for the two sides parallel to the interfaces. The vacuum region separating two irrelevant layers in the ribbon configuration is set to be 20 Å. All the calculations are self-consistent and the total energy convergence criterion is set at the value of 10-5 eV.

ASSOCIATED CONTENT Supporting Information. The following files are available free of charge. Supporting Information including additional characterization data and calculation results, additional discussions. (PDF)

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Author Contributions † These authors contribute equally to this work. M.W.C. conceived and supervised this study. Z. W. performed sample fabrication and TEM characterization. S. N. contributed to STEM simulation and interpretation of STEM data. M. L. performed DFT calculations. Y. I., H. K. and


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X. Z. contributed to CVD. M. W. C. and Z. W. wrote the manuscript. All the authors discussed the results and commented on the manuscript. Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This work is sponsored by Whiting School of Engineering, Johns Hopkins University, and World Premier International (WPI) Research Center Initiative for Atoms, Molecules and Materials, MEXT, Japan.

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SYNOPSIS


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FIGURES

Figure 1. Illustration on the 1H/1T interfaces in TMD monolayers. (a) Schematic model of the formation of 1T phase and resultant two 1H/1T interfaces by 1⁄2 0110 glide of cyan S atoms. (b) The atomic structure of 7c interface. The transition metal atom has seven coordinated S atoms. (c) The atomic structure of 5c interface. The transition metal atom only has five coordinated S atoms. In the atomic models, blue represents transition metal atoms and yellow or cyan represents S atoms closer to or away from the viewer. Such phase transformed regions are also observed in the Mo-doped WS2 monolayer samples: (d) bright-field TEM image taken from a two-phase region of a Mo0.2W0.8S2 monolayer. No contrast difference between 1H and 1T


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phases can be seen in the bright-field TEM image; (e) dark-field TEM image using the 1100 diffraction beam, indicated by the yellow circle in the inserted diffraction pattern, distinguishes 1H and 1T phases as brighter and darker regions, respectively.

Figure 2. STEM characterization of the 1H/1T interfaces in Mo-doped WS2 monolayer. (a,b) HAADF-STEM images of the 7c and 5c phase interfaces corresponding to the yellow squares in the inserted DF-TEM images.. (c,d) are high-magnification HAADF-STEM images showing the


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atomic structures of 7c and 5c interfaces. Brighter and dimmer spots are W and Mo atoms. Overlapped double S atoms in 1H phase appear comparable intensity to Mo, and single S atoms in 1T phase are almost invisible. Single line intensity profiles along the colored dashed lines across the interfaces in (c,d) are shown in (e) with corresponding colors. Mo atoms segregate at both 7c and 5c interfaces by forming dark chains and a few vacancies are observed at the Mo segregated interfaces. From (a-d), most 7c interfaces are straight down to atomic scale, whereas 5c interfaces generally contain a high density of atomic steps and thus looks meandering. Regions indicated by yellow squares in (b) are step-rich 5c interface segments. (f) Mo-proportion profiles by surveying from the interface regions with the width of eighteen atomic layers parallel to the interfaces and the length of over 200 atomic columns. Upper and lower panels are corresponding to 7c and 5c interfaces, respectively.


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Figure 3. Differential charge density distributions at the 7c and 5c interfaces of Mo-doped and pristine WS2 monolayers. Structure models and calculated charge density distributions: (a,e) Modoped WS2 7c; (b,f) WS2 7c; (c,g) Mo-doped WS2 5c; and (d,h) WS2 5c interfaces. In the models of 7c and 5c interfaces in Mo-doped WS2 monolayers (a,c), Mo atoms are placed only at the interfaces for simplicity.


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TABLES Table 1. Statistical analysis on the segregation behavior based on HAADF-STEM observations. Interface type

XW

XMo

Xvac

Mo/(Mo+W)

7c

0.02

0.91

0.07

0.98

5c

0.12

0.80

0.08

0.87


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One-Dimensional Atomic Segregation at ... - ACS Publications

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