Chemical Selectivity at Grain Boundary Dislocations in Monolayer Mo

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Letter

Chemical Selectivity at Grain Boundary Dislocations in Monolayer Mo WS Transition Metal Dichalcogenides 1-x

x

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Ziqian Wang, Yuhao Shen, Shoucong Ning, Yoshikazu Ito, Pan Liu, Zheng Tang, Takeshi Fujita, Akihiko Hirata, and Mingwei Chen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b08945 • Publication Date (Web): 18 Aug 2017 Downloaded from http://pubs.acs.org on August 22, 2017

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ACS Applied Materials & Interfaces

Chemical

Selectivity

at

Grain

Boundary

Dislocations in Monolayer Mo1-xWxS2 Transition Metal Dichalcogenides Ziqian Wang,1,2 Yuhao Shen,3 Shoucong Ning,4 Yoshikazu Ito,5 Pan Liu,6 Zheng Tang,3 Takeshi Fujita,2 Akihiko Hirata,2 Mingwei Chen,1,2,6,7* 1

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

21218, USA; Japan;

3

2

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

Key Laboratory of Polar Materials and Devices, East China Normal University,

Shanghai 200062, P. R. China; 4 Department of Mechanical and Aerospace Engineering, School of Engineering, Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong SAR; 5 Institute of Applied Physics, Graduate School of Pure and Applied Sciences, University of Tsukuba, Tsukuba 305-8573, Japan;

6

State Key Laboratory of Metal Matrix

Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, P.R. China;

7

CREST, JST, 4-1-8 Honcho Kawaguchi, Saitama 332-0012,

Japan. *Address correspondence to: [email protected]

ACS Paragon Plus Environment

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KEYWORDS:

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Two dimensional materials, Transition metal dichalcogenide, Dislocation,

Chemical selectivity, Grain boundary

ABSTRACT: Grain boundaries (GBs) are unavoidable crystal defects in polycrystalline materials and significantly influence their properties. However, the structure and chemistry of GBs in 2D transition metal dichalcogenide alloys have not been well established. Here we report significant chemical selectivity of transition metal atoms at GB dislocation cores in Mo1-xWxS2 monolayers. Different from classical elastic field-driven dislocation segregation in bulk crystals, the chemical selectivity in the 2D crystals originates prominently from variation of atomic coordination numbers at dislocation cores. This observation provides atomic insights into the topological effect on the chemistry of crystal defects in 2D materials.

Two-dimensional (2D) transition metal dicalcogenides (TMDs) have gained tremendous attention owing to the unconventional physical and chemical properties emerging from their ultimate thickness.1–4 Similar to bulk materials, doping or alloying has been demonstrated to be an effective way to manipulate the physical and chemical properties of the 2D crystals, including band gap engineering,5,6 modulation of carrier type and phase transition,7,8 reduction of contact resistance in 2D transistors,9 and catalytic activity enhancement.10 On the other hand, analogous to bulk crystals, the properties of 2D TMDs could be regulated by crystal defects, such as vacancies, dislocations, grain boundaries (GBs) and heterostructured interfaces, because of local symmetry breaking and chemical variation.11–15 Limited by the current synthesis techniques, it is still impossible to grow large-scale 2D TMD single crystals and, hence, GBs and dislocations are


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the unavoidable crystal defects in TMD monolayers. As a result, intensive studies have been carried out on the atomic structures of dislocations and GBs of pristine TMDs,16–18 as well as their important roles in mechanical,18,19 transport,13,20 magnetic,14 and optical properties.12,21 However, structure and chemistry of dislocations and GBs in 2D TMDs, particularly monolayer alloys, have not been well understood. In this study we employed the state-of-the-art spherical aberration corrected transmission electron microscopy (TEM) to investigate the atomic structure and chemistry of GB dislocations in monolayer Mo1-xWxS2 grown by low-pressure chemical vapor deposition (CVD).22 Obvious chemical selectivity at the cores of GB dislocations was observed, which provides insights into the chemical effects in 2D alloy materials as well as structure and chemistry of dislocations in the low-dimensional crystals. The monolayer Mo1-xWxS2 (x=0.3 and 0.5) samples, grown by low-pressure CVD,22 are transferred to holey carbon coated copper grids for TEM characterization. Figure 1 shows two sets of dark-field (DF) TEM images of monolayer Mo0.5W0.5S2 and Mo0.7W0.3S2 with representative GBs (Figure 1a,b and c,d) along with the corresponding selected area electron diffraction (SAED) patterns taken from the area including abutting grains (insets of Figure 1a-d). From the angle difference between diffraction vectors of two neighboring grains, the GB angles are measured to be approximately 12º and 22º for Figure 1a,b and c,d, respectively. DF-TEM micrographs shown in Figure 1a-d are imaged by selecting diffraction vectors indicated by the yellow circles in corresponding SAED patterns. By altering diffraction spots for imaging, the contrast of individual grains can be selectively changed as dark or bright and, consequently, the morphology of GBs can be visualized as shown in Figure 1a,b and c,d. In the microscopic images, the GBs are nearly straight. The directions of GBs are marked as the yellow dashed lines in the SAED patterns as shown in the inset of Figure 1b and d. It is notable that the GB


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directions appear to be the internal bisectors of the two groups of SAED patterns, indicating the formation of symmetrically tilted GBs in the CVD-grown Mo1-xWxS2 alloys, similar to previously observed GBs in CVD-grown MoS2.23 High resolution scanning TEM (STEM) images of the GBs in monolayer Mo0.5W0.5S2 and Mo0.7W0.3S2 are shown in Figure 2 a and c, which are taken from the GBs in Figure 1. In order to discriminate the transition metal atom W and Mo as well as the projected double S atoms simultaneously in the 1H phase Mo1-xWxS2 monolayers, high-angle annular dark-field (HAADF) STEM imaging was utilized for achieving atomic-scale spatial and chemical resolutions. Since atomic contrast in HAADF-STEM images is roughly proportional to the square of atomic number Z and the superposed two S atoms has comparable contrast to a Mo atom,24 the bright, less bright and dark spots in the images represent W, Mo and 2S atoms, respectively (Figure 2a,c). The random mixture of Mo and W atoms in atomic scale is consequently confirmed. From the atomic image of Figure 2a, the 12º GB shown in the DF-TEM image (Figure 1a,b) is determined to be a 0001235 0 symmetrical tilt boundary which consists of an array of dislocations at an interval of ~1.4 nm on average. Similarly, an array of dislocations with an interval of ~0.8 nm defines the 22º GB as a 0001123 0 symmetrical tilt boundary (Figure 2c). From the zoom-in HAADF-STEM images of the vicinity of single dislocation cores (Figure 2b,d), the dislocations at both GBs are found to have the same atomic configuration of a 5|7-fold ring.17 On the basis of the atomic images, the structure of the dislocations can be described by a model shown in Figure 3a. As indicated hereinbefore, the dislocation core is constituted by two interconnected 5- and 7-member atomic rings by sharing a pair of transition metal M atoms (M representing either Mo or W) as the common rim, analogous to the GB dislocations in pristine TMDs without dopants.14,17 In this study, we name the atomic


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(pair) sites as the shared M pair, upper M pair and lower M site for simplicity as indicated in Figure 3a. Except for the shared M pair, all other M atoms in the dislocation core are coordinated with six S atoms while each S atom bonds with three M atoms, i.e. each of the adjacent M or S atoms is intercepted by the others, which has the identical circumstance as an atom in perfect 1H crystal lattices.17 In contrast, each atom of the shared M pairs is coordinated with four S atoms, together with uneven M-M bonding lengths with the neighboring M atoms. Therefore, the shared M pairs are topologically defective. According to the dislocation structure, the experimental atomic images of the representative dislocation cores and their corresponding schematic atomic models are shown in Figure 2b,d. The Burgers vector 1 010 of dislocations are indicated by the purple arrows in Figure 2a,c under the definition of unit vectors in the models described in Figure 2b,d. Apparently, the GB angles between the neighboring grains originate from accumulation of the Burgers vectors from dislocation arrays. In accordance with the geometric model of classical dislocation GBs in bulk crystals,25 the GB angles in the monolayer crystals can be well described by the equation: = ⁄ , where b notes the length of Burgers vector and D notes the average interval between adjacent dislocation cores. By directly counting the component atoms based on contrast of atomic columns, the average chemical compositions of the 12º and 22º GBs are measured to be Mo0.5W0.5S2 and Mo0.7W0.3S2, respectively. However, careful observations can identify that the occupations of transition metal atoms at the 5|7-fold dislocation cores are chemically selective, different from the random distribution in perfect 1H lattices. We thus performed statistical analysis of chemical occupations at the shared M pairs, upper M pairs and lower M sites of dislocation cores in the 2D alloys. A series of HAADF-STEM images (containing over 100 dislocations in total) of each GB are captured for statistical analysis, and all the images used in statistics are of the same Mo/W


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ratio x=0.50±0.02 for the series of 12º 0001235 0 GB in Figure 2a,b and x=0.30±0.02 for 22º 0001123 0 GB in Figure 2c,d. Statistical results on 12º 0001235 0 GB are shown in Figure 3b-d and 22º 0001123 0 GB in Figure S1 (Supporting Information). For the dislocations with the same geometric structure, there are twenty possible atomic configurations with various Mo and W occupations as displayed in Figure 3b. Herein, atomic configurations chiral to each other are classified as a single type. The configurations arranged in the same row have the same type of shared M pairs, i.e. Mo-Mo, Mo-W and W-W for the 1st, 2nd and 3rd row, respectively. Similar arrangements are simultaneously made with respect to upper M pairs, i.e. Mo-Mo, Mo-W and W-W for 1st, 2nd and 3rd column, respectively. The frequencies of these configurations from the statistical measurements are shown in Figure 3c,d by rows and by columns indicated by the purple and orange arrows in Figure 3b. Moreover, the grey or blue shaded area in each bar represents the constituent of Mo or W occupation at the lower M sites. The black scatters and lines in Figure 3c,d are the calculated standard values for each classification by using the composition x=0.50 and assuming equal possibility of chemical occupation of W and Mo at each M-site. Descriptions on the standard frequency of each dislocation core configuration are detailed in S3. There is good consistency between the observed frequency of each type of upper M pair (bar in Figure 3d) and the corresponding standard value (scatter in Figure 3d), suggesting the absence of partitioning occupation at the upper M pairs of the dislocations. In contrast, there is a large deviation between the standard values and the experimental measurements of the shared M pairs (Figure 3c). Particularly, the frequency of the Mo-Mo pair is experimentally measure to be 73%, much higher than 25% from the estimation from random chemical occupation, showing significantly preferential occupation of the Mo-Mo pair at the shared M pair site over Mo-W and W-W pairs. Furthermore, the blue


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bars in Figure 3c,d are all observed to be higher than corresponding grey ones, indicating slight preference of W occupation over Mo occupation at the lower M sites, i.e. 63% of W over 37% of Mo occupation at the lower M site. The same trends are observed from the statistics of 22º 0001123 0 GB as shown in Figure S1, confirming the generality of the observed chemical selectivity of the dislocations in the 2D TMD alloys. The statistical frequencies of all twenty configurations of dislocations in Figure 3b are shown in Figure S3a for 12º 0001235 0 and Figure S3b for 22º 0001123 0 GBs, respectively. On the basis of the experimental characterization, we carried out DFT calculations to explore the origins of the chemical selectivity at the GB dislocations. Figure 4a-f shows six models used in the calculations representing the cases of W-W, W-Mo, Mo-Mo occupation for the shared M pairs and the upper M pairs. Each unit cell contains twenty-eight Mo and two W atoms and two 5|7-fold dislocation sits at the middle to form a symmetric GB. Periodic boundary conditions are employed in the calculations.14 Figure 4g shows the comparison of the total energies of the six representative models. An energy increase of ~0.2 eV per W atom occupying the shared M pairs is predicted (left panel of Figure 4g), which is an order larger than the energy variation of ~0.01 eV per W-occupation at upper M pairs (right panel). Such energy difference indicates the preference of Mo-rich shared M pairs and the evenness of elemental distribution at upper M pairs in consensus with the trends of our experimental observations. However, no energy difference was predicted between Mo-occupation and W-occupation of the lower M sites by DFT calculations. The higher fraction of W-occupation at the lower M sites in our experimental observations may come from the finite temperature and complicated chemical environment around the lower M sites in real materials, which cannot be fully reproduced by DFT modelling. Discussion on the representativeness of the models and the results based on


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other possible models is included in S4. Figure 4h plots the calculated density of state (DOS) based on the structure models shown in Figure 4a-c to show the effects of chemical selectivity on the electronic states of GB dislocations. Red and blue shaded regions in Figure 4h represent the DOS projected onto the atoms forming the GB dislocations for major and minor spins, respectively. The grey shaded regions are projected DOS on atoms at perfect lattices. Localized states of antibonding δ* from the GBs shifts towards higher energy as the shared Mo-Mo pair is replaced by Mo-W and W-W, accounting for the increasing repulsive interaction between shared M atoms with higher W occupation. Meanwhile, the bonding δ states keep nearly unchanging at the Fermi-level. As a complementary to the aforementioned calculations, we also calculated the case that two Mo atoms are embedded in twenty-eight W atoms in a unit cell. The same result that the total energy drops as the shared M pair changes from W-W and Mo-W to Mo-Mo is obtained (Figure S5). Therefore, chemical selection of Mo at GB dislocations found in our experimental observation is independent on the compositions of the TMD alloys and should be a universal phenomenon in the monolayer Mo1-xWxS2. As expounded by DFT calculations, the chemical selectivity of the shared M pairs at GB dislocations of monolayer Mo1-xWxS2 is energetically favorable and the selection of the shared Mo-Mo pair at dislocation cores results in the shift of the anti-bonding δ* localized states and thus the obvious decrease of the total energy. Therefore, the shared Mo-Mo pair shows the highest frequency in the experimental observations. From topological viewpoint, the most apparent difference between shared M pairs and other atomic sites of the dislocation core (upper M pairs, lower M sites or M sites in bulk area) is the coordination number. Each M atom in the shared pair is only coordinated with four S atoms, not six of the perfect lattice. Apparently, the topological defect of the shared M pair principally accounts for the excess energy of GBs. It can


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be deduced that the selection of the Mo-Mo pair may be associated with the bonding energy difference between W-S and Mo-S bonds along with the difference between Mo-Mo and Mo-W or W-W bonds. Although W and Mo have very similar chemical properties and a nearly identical atomic size, the W-S bond has higher bonding energy compared with the Mo-S bond26 and, on the other hand, the Mo-Mo bond is less repulsive than Mo-W or W-W bonds, as evidenced in our DFT calculations. In principle, the selection of the shared Mo-Mo pair with a relatively lower Mo-S bonding energy and a less repulsive Mo-Mo bonding is helpful in minimizing the excess dislocation energy without violating the topology of the GB dislocation core. For the weak preference of W at the lower M sites, it may result from local electron density variation or lattice distortion stimulated by neighboring shared M-M pairs with topological defects. The interaction of dislocations with chemical dopants and impurities has been long known because it significantly affects the dislocation dynamics and thereby many physical and mechanical properties. In general, the segregation of foreign atoms around dislocation cores is driven by the elastic stress fields and leads to the formation of Cottrell atmospheres.25,27 Intrinsically different from the classical dislocation segregation in bulk crystals as well as the segregation behavior at S-rich GBs in monolayer Mo1-xWxS2 system,28 the chemical selectivity of M-rich GB dislocations, discovered in this study, originates from the topological environment of the M atoms in the 2D crystals. The coordination number variation of the constituent atoms of dislocation cores, together with the diminutive bonding energy difference between W-S and Mo, results in the significant chemical selectivity. Apparently, the chemical effect will influence the dynamics and stability of dislocations and thus dislocations related kinetic processes of the 2D materials. Moreover, this study also demonstrated that TMD monolayers act as a platform for an unambiguous analysis of the occupation and partitioning of single atomic sites by different


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elements owing to the 2D nature, because of only one transition metal atom in a column when viewed by chemically-resolved STEM-HAADF. Finally, our study unveils the strong correlation between topological configurations and chemical selectivity at dislocation cores in monolayer TMD alloys, which provides atomic insights into the unique structure and dynamics of crystal defects in 2D materials.

ASSOCIATED CONTENT Supporting Information The following files are available free of charge. Experimental methods regarding CVD growth of Mo1-xWxS2 monolayers, microstructure characterization, and computational details; Supplementary statistical results; Definition of the standard frequencies of each dislocation configuration; Discussion on the representativeness of the models; Additional DFT calculation results (PDF)

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

Notes


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The authors declare no competing financial interest.

ACKNOWLEDGMENT This work is sponsored by JST-CREST “Phase Interface Science for Highly Efficient Energy Utilization”, JST, Japan; World Premier International (WPI) Research Center Initiative for Atoms, Molecules and Materials, MEXT, Japan, Japan Society for the Promotion of Science (JSPS) Grant-in-Aid for Scientific Research on Innovative Areas “Science of Atomic Layers” (Grant No. 26107504), and Interdepartmental Program for Multi-dimensional Materials Science Leaders in Tohoku University and Tohoku University Division for Interdisciplinary Advanced Research and Education. S. N is grateful for the financial support by the General Research Fund (Project number, 622813) from the Hong Kong Research Grants Council.

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Azizi, A.; Wang, Y.; Stone, G.; Elias, A. L.; Lin, Z.; Terrones, M.; Crespi, V. H.; Alem, N. Defect Coupling and Sub-Angstrom Structural Distortions in W1–xMoxS2 Monolayers. Nano Lett. 2017, 17 (5), 2802–2808.


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SYNOPSIS


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Figures and captions

Figure 1. Microscopic view of grain boundaries in monolayer Mo1-xWxS2 under dark-field TEM imaging. (a,b) A 12º grain boundary in monolayer Mo0.5W0.5S2 imaged by using the circled diffraction spots shown in the insets. (c,d) A 22º grain boundary in monolayer Mo0.5W0.5S2 imaged by selecting different diffraction vectors. (Scale bar: 2 µm)


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Figure 2. Two representative symmetrical tilt dislocation grain boundaries in monolayer Mo1xWxS2.

(a) HAADF-STEM image of a 12º 0001235 0 symmetrical grain boundary in

Mo0.5W0.5S2. Dislocation cores are marked by the yellow arrow heads. Upper panel of (b): Zoom-in image and atomic model of a dislocation core marked by the yellow dashed square in (a). W, Mo and S atoms are represented by the blue, grey and yellow balls in the model. Lower panel of (b): intensity profiles along the two dashed lines with corresponding colors. The Burgers vector of the dislocation is noted by the purple arrow and determined by a Burgers loop shown in (a). (c) HAADF-STEM image of a 22º 0001123 0 symmetrical tilt grain boundary in Mo0.7W0.3S2 and (d) zoom-in image, atomic model and intensity profiles of a dislcoation.The


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dislocations have the same configuration and Burgers vector as those at 12º 0001235 0

symmetrical tilt grain boundary in Mo0.5W0.5S2.


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Figure 3. Statistical analysis on the atomic occupation at dislocation cores in 12º 0001235 0

symmetrical grain boundary of monolayer Mo0.5W0.5S2. (a) Atomic model of a grain boundary dislocation and the definition of different atomic sites in the vicinity of the dislocation core. (b) All possible atomic configurations with different occupations of Mo or W atom in each transition metal site. Color stipulations are the same as the atomic models in Figure 2b,d. Dislocation


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configurations in the same row have the same type of the shared M pair occupation and the ones in the same column have the same upper M pair occupation. Statistics on the frequency of dislocation core configurations are shown in (c,d) according to different (c) shared M pair occupations and (d) upper M pair occupations. Each statistic bar in (c) and (d) shows the summed frequency of configurations in the corresponding row and column in (b). Black scatters and lines in (c,d) denote the standard frequency from the assumption of random chemical occupation of each site. Blue and grey columns and bars in (c,d) show the constituent of Mo and W-occupied lower M site, respectively. From (c,d) a strong preference of Mo occupying the shared M pairs can be observed whereas there is almost even occupation of upper M pairs.


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Figure 4. DFT calculations of Mo-rich dislocation in monolayer Mo1-xWxS2. (a-f) Atomic models used for calculations: (a-c) considering the circumstances of W-W, W-Mo and Mo-Mo occupying the shared M pairs, and (d-f) considering W-W, W-Mo and Mo-Mo occupying the upper M pairs. Each of (a-f) displays a repeating period of the models containing 2 W and 28 Mo atoms. (g) From left to right the calculated total energies of the GB structures of (a-f). (h) Calculated density of states (DOS) showing the influence of chemical composition on the electronic states of dislocation cores. Projected DOS of GBs consisting of W-W occupied, WMo occupied and Mo-Mo occupied shared M pairs are shown in the upper, middle and lower panels in (h). The red and blue shaded areas are calculated DOS for major and minor spins projected onto the atoms forming the M-rich GB dislocations and the grey shaded areas are


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projected DOS on the bulk areas excluding M-rich and S-rich GBs. Upshift of the energy of antibonding δ* localized states can be observed with increasing fraction of W in the shared M pairs.


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Chemical Selectivity at Grain Boundary Dislocations in Monolayer Mo

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