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Controllable Growth and Formation Mechanisms of Dislocated WS2 Spirals Xiaopeng Fan, Yuzhou Zhao, Weihao Zheng, Honglai Li, Xueping Wu, Xuelu Hu, Xuehong Zhang, Xiaoli Zhu, Qinglin Zhang, Xiao Wang, Bin Yang, Jianghua Chen, Song Jin, and Anlian Pan Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.8b01210 • Publication Date (Web): 15 May 2018 Downloaded from http://pubs.acs.org on May 15, 2018

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Controllable Growth and Formation Mechanisms of Dislocated WS2 Spirals Xiaopeng Fan†||, Yuzhou Zhao‡||, Weihao Zheng†, Honglai Li†, Xueping Wu†, Xuelu Hu†, Xuehong Zhang†, Xiaoli Zhu†, Qinglin Zhang†, Xiao Wang†, Bin Yang†, Jianghua Chen†, Song Jin‡,*, Anlian Pan†,*

†

Key Laboratory for Micro-Nano Physics and Technology of Hunan Province, State Key

Laboratory of Chemo/Biosensing and Chemometrics, and College of Materials Science and Engineering, Hunan University, Changsha 410082, China ‡

Department of Chemistry, University of Wisconsin—Madison, 1101 University Avenue, Madison, Wisconsin 53706, United States *Corresponding authors. E-mail: [email protected]; [email protected]

ABSTRACT: Two-dimensional (2D) layered metal dichalcogenides can form spiral nanostructures by a screw-dislocation-driven mechanism, which leads to changes in crystal symmetry and layer stackings that introduce attractive physical properties different from their bulk and few-layer nanostructures. However, controllable growth of spirals is challenging and their growth mechanisms are poorly understood. Here, we report the controllable growth of WS2 spiral nanoplates with different stackings by a vapor phase deposition route and investigate their

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formation mechanisms by combining atomic force microscopy with second harmonic generation imaging. Previously not observed “spiral arm” features could be explained as covered dislocation spiral steps, and the number of spiral arms correlates with the number of screw dislocations initiated at the bottom plane. The supersaturation-dependent growth can generate new screw dislocations from the existing layers, or even new layers templated by existing screw dislocations. Different number of dislocations and orientation of new layers result in distinct morphologies and different layer stackings, and more complex nanostructures, such as triangular spiral nanoplates with hexagonal spiral pattern on top. This work provides the understanding and control of dislocation-driven growth of 2D nanostructures. These spiral nanostructures offer diverse candidates for probing the physical properties of layered materials and exploring new applications in functional nanoelectronic and optoelectronic devices.

KEYWORDS: Transition metal dichalcogenides, two-dimensional materials, screw-dislocationdriven growth, spiral structure, stackings, second harmonic generation imaging


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Two-dimensional (2D) layered transition metal dichalcogenides (TMDs) (e.g. MoS2, MoSe2, WS2 and WSe2, etc., generally expressed as MX2) have emerged as an exciting new class of nanomaterials for optoelectronic device applications.1-10 Monolayer and multilayer TMDs have sensitive layer-number-dependent optical and electronic properties, such as a transition from an indirect band gap in multilayers or bulk to a direct band gap in monolayer,1 an emerging valley degree of freedom11-16 and a much higher current on/off ratio in field-effect transistors based on monolayer structure compared to bulk counterparts at room temperature.17 In addition to the layer number, layer stacking behavior also plays a crucial role in determining the fundamental properties of TMDs.18-25 Spiral TMD structures obtained by screw-dislocation-driven growth (SDD) process display different and sometimes complex and non-centrosymmetric stacking behaviors26,27 compared to normal 2H stacking multilayer structures which were prepared from either mechanical exfoliation from bulk materials or chemical vapor deposition (CVD) growth.2,18-24,28,29 Screw dislocations, as a type of universal line defect, can act as the driving force to break the symmetry of crystal structure and promote the formation of spiral TMD structures with various stacking behaviors. Layered materials are generally prone to contain screw dislocations, and it has been understood that their crystal growth could be driven by screw dislocations.27,30-34 Only recently, careful structural and spectroscopic analysis revealed that screw-dislocation-driven pyramidal spiral WSe2 nanoplates with triangular, hexagonal, or mixed shapes have different stacking behaviors (sometimes non-centrosymmetric) that are determined by the layer orientation and number of the screw dislocations.26 Strong nonlinear optical properties have been observed in


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spiral TMD nanostructures that lack inversion symmetry.27,35 For example, higher intensity second harmonic generation (SHG) can be observed from the spiral nanostructures of WS2 with the increase of layer number.35 However, the formation mechanism(s) of complex dislocated structures, such as the structural transition from triangular spirals to hexagonal spirals and the formation of nanostructures with multiple screw dislocations, remain elusive. Due to the lack of the understanding of the growth mechanisms of these complex dislocated nanostructures, it is still a significant challenge to controllably and reproducibly grow spiral TMDs with a variety of stacking behaviors and the corresponding physical properties. Therefore, it is essential to comprehensively understand the microstructural information of these 2D layered crystals. In this work, we demonstrate the controllable CVD growth of layer-by-layer WS2 nanoplates and a series of spiral WS2 nanostructures with increasing number of dislocations and complexity, including prototypic triangular spirals and triangular-hexagonal mixed spirals. On the basis of careful structural studies, we propose the formation mechanisms of these spiral nanostructures. We also explain the “spiral arm” features previously not observed and further correlate the number of spiral arms with the number of screw dislocations. Transition of morphologies and stacking behaviors is found in more complex spiral structures. SHG in combination with microscopic characterizations confirms the growth mechanisms of these spiral WS2 nanostructures and reveals that the number of screw dislocations, and the introduction and orientation of new layer determine the morphologies and stacking behaviors of the complex spiral nanostructures. These understandings will promote the fundamental research of layered materials and lead to potential


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applications in integrated nanoelectronics and nanophotonics. Controllable growth of wedding cake structures and spiral structures: WS2 nanostructures were synthesized by heating WS2 precursor in a flow of argon gas in a globally heated tube furnace as shown in Figure 1a and b (see experimental details in Methods in the Supporting Information). As illustrated in Figure 1c, both the deposition temperature and the concentration of precursor have significant influences on the formation of WS2 nanostructures. The wedding cake structures were formed via the layer-by-layer (LBL) growth mode (Figure 1d), while the spiral structures were grown via the SDD growth mode (Figure 1e). The wedding cake structures and spiral structures were controllably formed in different regions along the axis of the tube furnace. As shown in Figure 1b, the wedding cake structures generally form in the high temperature deposition region with a high molecular concentration of precursors, whereas spiral structures were grown in the low temperature deposition region further downstream with a low molecular concentration.26,31,33,36


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Figure 1. Controllable growth of wedding cake structures and spiral structures of WS2. (a) Illustration of the CVD experiment setup for TMD synthesis. (b) A structural schematic of nanoplates illustrate the distributions of wedding cake structures and spiral structures in different deposition regions. (c) Trends in the synthesis conditions for forming the wedding cake structures vs. the spiral structures. Schematic illustrations of (d) layer-by-layer (LBL) growth mechanism and (e) screw-dislocation-driven (SDD) growth mechanism.

According to the well-known Burton-Cabrera-Frank (BCF) crystal growth theory,37,38 the crystal growth process is governed by the supersaturation of local growth environment. The supersaturation is defined as σ = In(c/c0), where σ is the degree of supersaturation, the c and c0 are the local precursor concentration and equilibrium precursor concentration, respectively. Both c and c0 are affected by the temperature profile of the furnace. Specifically, c0 is determined by the thermodynamic equilibrium of the deposition reaction, whose equilibrium constant is essentially determined by the local temperature. The c is dynamically influenced by the temperature at which the precursor is heated at, the transport of precursor molecules along the reactor, as well as the


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concentration loss caused by deposition prior to the growth region. At a given local temperature, pressure, and flow rate, c is then mostly affected by the consumption of precursor due to deposition prior to the growth region, which is determined by the temperature gradient of the furnace. Hence, the local supersaturation is primarily affected by the deposition temperature, thus the growth of nanostructures is also sensitive to the deposition temperature. As summarized in Figure 1c, the wedding cake structures of WS2 prefer to form in the high temperature deposition region (around 950 oC) with a high supersaturation, following the LBL growth mode (Figure 1d).36 In contrast, the spiral structures with screw dislocations are more prone to form in low temperature deposition regions with a low supersaturation, following the SDD growth mode, where atoms can be added to the spiral step edges (Figure 1e).26,31,34,39,40 Thus, the spiral WS2 structures are the predominant structures to form in the lower temperature (600~800 oC) downstream end of the tube furnace. Furthermore, toward the lower temperature region (600 oC), more screw dislocations are found in each spiral nanostructure than the simple triangular single dislocation spirals predominantly found in the 800 oC temperature region (Figure 1b), suggesting that lower supersaturation condition tends to induce more complex spiral structures corresponding to multi-spiral pattern structures (see more discussion later). In order to obtain detailed morphological information about the various WS2 nanostructure products, optical microscopy and atomic force microscopy (AFM) characterizations were conducted. The typical optical image (Figure 2a inset) shows a large number of wedding cake structures were synthesized in the region with temperature of ~950 oC. The magnified optical


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image of one representative WS2 nanoplate (Figure 2a) suggests a flat triangular top on a larger and thinner triangular flake. AFM (Figure 2e) further confirms a wedding cake structure which is formed by the stack of independent single layers and completely flat on top. The AFM line profile reveals the total height of the nanoplate is about 13 nm. In contrast, the representative optical images (Figures 2b-d) of the structures grown in the regions of the substrate with temperature ranging from 800 to 600 oC show increasingly more complex features that suggest spiral structures (as illustrated in Figure 1b). In region I (near 800 oC), the structures present simple triangular spiral pattern (Figure 2b).41 Compared to region I, mixed simple spiral patterns and multi-spiral patterns are found in region II (800 oC to 600 oC). Moreover, more complex multi-spiral patterns are frequently observed in region III (Figure 2d, near 600 oC). Several representative AFM images of the top regions of these nanostructures (Figure 2f-h) reveal triangular spiral nanoplates with increasing number of screw dislocation spirals (one in Figure 2f, two in Figure 2g, and three in Figure 2h) and different included angle of screw dislocations. The included angle is defined as the angle between the traces of the different dislocation spirals that share a common core, as illustrated by the dashed lines in Figure 2g near the dislocation core. Based on the 3-fold symmetry of the monolayer TMD structure and the orientations of each layer in the stacking, there can be two sets of included angles between screw dislocation spirals. When the included angles are only 0°, 120° or 240°, with the increase of the number of screw dislocations, the center structure of spiral nanoplates remains in the form of triangular spiral traces as shown in Figure 2f-h. Figure 2h displays a particular triple spiral pattern structure which not only contains 0° but also contains 120°


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included angles. In contrast, when we observe the other set of included angles of 60°, 180° and 300°, the center of those spiral nanoplates have hexagonal spiral traces. Note that the hexagonal nanostructures are much less commonly observed than the triangular structures and more details will be discussed later.

Figure 2. Optical and AFM characterizations of the wedding cake structures and spiral structures of WS2. (a) Optical images of wedding cake structures. (b-d) Representative optical images of various spiral structures found in different deposition region on the substrates. Insets show the distribution of nanoplates on the substrates, the total field of view of a is 100 micron wide, and 300 micron for b-d. (e) AFM characterization of a wedding cake structure. AFM line profile along the black line in e shows that the height of wedding cake structure is about 13 nm. (f-h) AFM images of several representative spiral structures, all of which are triangular dislocation spirals but the number of screw dislocation increases from one to three. The included angle of screw dislocations in g is 120°. There are two included angle in (h), 0° and 120°.

Formation mechanisms of spiral structures: While the growth kinetics of the clear screw dislocation spiral steps revealed by AFM has been well understood,37,42,43 another obvious and curious features are the “spiral arms”, resembling the spiral arms of a galaxy, that were found at


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the outer and thinner edge of these larger nanoplates. The nanoplates in Figures 2b, 2c and 2d display one, two and three spiral arms, respectively. These spiral arm features have not been previously observed and seem to be clearly visible only in the thin regions of the spiral nanoplates when they are very large (above 10 micron), therefore it is worthwhile to study them further to elucidate the potential connection to screw dislocations. We chose a representative nanoplate with a single spiral arm on the edge for detailed study (Figure 3a). AFM image of the dislocation center on top (Figure 3b) reveals a single triangular dislocation spiral. The AFM in Figure 3c shows the detailed topography of the spiral arm structure near the edge of the whole plate. Interestingly, the height of the spiral arm area is about one layer higher compared to the complementary area nearby. The line scan height profile (Figure 3d) following the blue trace in Figure 3c shows that the layer thickness decreases as we move from the spiral arm area to the complementary area, but then increases by approximately one layer (about 0.9 nm) when we move back to the spiral arm on the outer rim. This difference in layer thickness in the few-layer region of these large WS2 plates causes the clear optical contrast that gives rise to the spiral arm features observed.


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Figure 3. Formation mechanisms of spiral nanostructures. (a) AFM image of a single spiral structure. (b, c) Magnified AFM images corresponding to the marked areas in a. (d) Line profile marked with a blue and black line in c showing that the very outer rim of the spiral is thicker. (ef) Schematic diagram of the growth mechanism of single spiral pattern emphasizing the initial overlapping spiral layers and showing the formation progress of single spiral pattern with a “spiral arm”. (g) A representative AFM image showing two screw dislocations are terminated at the center and off center during growth progress. (h) The same AFM image in g but decorated with black and red dash lines to emphasize the centers of dislocations and the evolution of the spiral steps. (i) A schematic illustration from the side view of the pyramidal nanoplate of the two types of dislocation-driven growth processes, Type 1 and 2 to illustrate the formation mechanisms of complex spiral structures. The obvious next questions are why would there be an extra layer in a spiral plate and if and how this might be related to screw dislocations. One might imagine that the thicker spiral arm could be a layer covering the top surface; however, it actually results from an extra layer inserted at the bottom of the nanoplate (i.e. buried below the top layers), because the exposed edge is traced from the center of the plate to the edge without coming across another layer. Our interpretation is


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further verified by the optical and AFM characterizations of another spiral structure with two spiral arms (Figure S1). To understand the formation process of the spiral arm structure, a freestanding monolayer with a screw dislocation was introduced. If both sides are equally exposed to the precursor flux, a screw dislocation spiral would grow along both directions. As a result, a bipyramidal spiral nanoplate is formed, which has been demonstrated in the solution grown SDD free-standing Bi2Se3 nanoplates. However, in our case with the nanoplates grown on substrate, the bottom side of the spiral plate are not exposed to precursor flux during the growth. Thus, that hypothetical bipyramid would have to be compressed and the bottom of the pyramid pushed against the flat substrate, as shown in the side view in Figure 3e. Figures 3e and 3f (top view) illustrate the situation of growing a spiral structure on a flat substrate. The extra layer thickness arises from the overlapped region caused by screw dislocation. As marked by the red dashed line, the overlap of the bottom part leads to the complete coverage of one edge of the layer beneath, which is not exposed to the precursor vapor anymore during the growth process, and thus that step edge stops growing. However, the other edges are still exposed to precursor vapor and keep propagating during the growth, resulting in the formation of new overlapped edge and the eventual spiral arm structures. Since the spiral arms are directly caused by the screw dislocation introduced at the bottom layer, the number of spiral arms equals to the number of screw dislocation initiated from the bottom layer. Furthermore, we have found that the morphologies of spiral structures can be characterized by the number of screw dislocations, the included angle between screw dislocations spirals, and


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the spiral arm structures. Figure 2 demonstrates that these spiral nanoplates display the same number of spiral arms in the outer plane of the plates as the number of screw dislocations at the center, suggesting a one-to-one correspondence. However, the number of screw dislocation observed at the center of the nanoplates can be occasionally smaller than the number of spiral arms at the outer edge of the nanoplates (see an example in Figure S2). Based on the understanding of the single spiral structure, we propose the formation mechanisms of more complex spirals as illustrated in Figure 3i and classified into Type 1 and 2 growth process. For Type 1, the screw dislocation is introduced from the beginning of the growth process at the very bottom plane. The termination position of the spiral pattern is located on the top surface and finally resides at center of the nanoplate. This is what we see in Figure 3a-c and most of the spiral plates so far (such as those in Figure 2f-h). For Type 2, besides the screw dislocations in center, new screw dislocations are introduced at the edge of a nanoplate after growth for a certain time, resulting in a spiral trace with the termination position off the center of the plate. After screw dislocation is introduced, the screw dislocation grows along the out-of-plane direction with both sides terminating on the top and bottom surface as the nanoplate keeps growing laterally and vertically. Figure 3g displays a representative AFM image of a nanoplate that exhibits Type 2 growth. The spiral trace is marked with the red guideline (Figure 3h) that shows a screw dislocation in the center and the black guideline that shows another screw dislocation off center. So the number of screw dislocation observed on top is the same as the number of spiral arms at the outer edge of the nanoplates (Figure S2). It is noted that all the layers grown by Type 1 and Type 2 should have the same orientation


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since they are the same connected layer. We believe the difference of Type 1 and Type 2 growth is caused by the sequence of the introduction of screw dislocations during the growth, which can be influenced by the surface roughness of the substrate 43 or the nuclei particles on the surface of the substrate,44 but at this stage, it is still challenging to precisely control these two types. The structural transition from triangular spirals to hexagonal spirals: The more complex spiral structures include cases where there seems to be structural transition from triangular spirals to hexagonal spirals. Figure 4a shows a plate that is found among the multi-dislocation plates toward the lower reaction temperature, which has a hexagonal double spiral located on top of a triangular spiral plate. Such abrupt structure change that makes hexagonal spiral to grow on the top of triangular spiral can only be explained by the introduction of new layer in the middle of growth process, as shown in Figure 4b in a cross-section view. As further illustrated by Figure 4ce, during the growth process of a single triangular dislocation spiral (Figure 4c), a new WS2 layer with opposite layer orientation could be introduced because of the local supersaturation change (Figure 4d), which is templated by the screw dislocation underneath and further evolves into another screw dislocation with a shared dislocation core but opposite orientation (Figure 4e).


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Figure 4. The introduction of new screw dislocations during SDD growth that results in more complex spirals. (a) The optical image of WS2 spiral nanostructure in which structural transition from triangular spiral to hexagonal spiral in the center of spiral nanostructure. (b) The schematic formation diagram of complex spiral structure, named Type 3 growth, in the cross-section view. New layer is introduced during the growth process and the triangular or hexagon spiral nanostructure show different staking behaviors. (c-e) The schematic illustration of the growth process from (c) a perfect single triangular spiral pattern, to (d) the initiation of a new dislocation spiral on top of the existing one, to (e) the combination of two processes leads to a more complex spiral plate with a hexagonal spiral on top of a triangular spiral. (f-h) AFM images of representative examples corresponding to the schematic figures: (f) A single spiral structure, (g) a multilayer structure on the top of bottom spiral structure, which confirms the existence of the new layer generation process shown in (d), (h) a hexagonal dislocation spiral atop of and at the center of a triangular dislocation spiral; which confirms the scenario in (e).

To provide supporting evidences for this proposed growth process, intermediate nanoplates were captured by quickly removing the WS2 source during the synthesis. Figure 4g shows a nanoplate where the bottom part is a single triangular spiral structure, but the top part is an emerging wedding cake structure caused by LBL growth which may result from a sudden increase of supersaturation. Furthermore, the orientation of this new layer is opposite to the bottom spiral structure, leading to the 2H “+-”stacking at the interface.26 Finally, Figure 4h shows a nanoplate


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that is consisted of a fully developed hexagonal dislocation spiral with two screw dislocations with an included angle of 60o sitting on top of a larger triangular single-dislocation plate. We call this Type 3 dislocation growth process. Figure S3 shows another example where the orientation of the new crystal layer is the same as the single spiral structure below. For all of these plates, the number of spiral arm observed is also less than the number of spiral traces at the top revealed by AFM, because a new layer and screw dislocation(s) is (are) introduced in the middle of the growth process. Previous study has shown that the number of screw dislocations and the rotation of the screw dislocations could impact the morphology of spirals and the layer stackings.26 The results and discussion above explain the mechanisms for the origin of the initial screw dislocation and how such dislocations result in the spiral arm features seen in the plates, furthermore, they illustrate how additional screw dislocations can emerge and lead to the formation of multiple and complex dislocation spiral patterns depending on the relative location and orientation of the new screw dislocations.

Verification of the stacking behaviors by SHG imaging: We further conducted SHG imaging in reflection mode (see details in the Supporting Information) to probe the structural symmetry and further verify the various stacking behaviors found in these diverse WS2 spiral plates. Because only noncentrosymmetric layer stacking in multilayer MX2 generates SHG signal, SHG imaging reveal the spatial dependence of the layer stacking and symmetry. Figures 5a-d reveal the structural information of a representative WS2 plate with a single triangular spiral. AFM (Figure 5a) clearly


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shows the single screw dislocation in the center of the nanoplate and the spiral arm is clearly seen in the optical image (Figure 5c). With the screw dislocation, each layer of the whole nanoplate should have the same orientation as the deflection of the interlayer lattice is ignored,35 here we use the “+ +” notation to represent the simplified orientations of different layers (Figure 5b). The SHG result of a single spiral structure is totally different from the typical bulk 2H stacking TMDs with centrosymmetry, which has “+ ̶ ” notation and does not yield SHG signal.18,26,45 SHG mapping in Figure 5d also clearly illustrates that SHG intensity increases with increasing layer number of single dislocation spiral, confirming that the single spiral structure lacks centrosymmetry.35 Similar results were obtained from triangular nanoplates with more screw dislocations. Figure 5e-h and 5i-l show spirals with two or five dislocations with included angles of 120° and 0°, respectively. Such arrangements still yield the same orientation of each layer and thus are noncentrommetric. Indeed, SHG mapping of both nanoplates (Figure 5h, 5l) shows the SHG intensity increasing as thickness increases with the strongest intensity in the center.


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Figure 5. Verification of the stacking symmetry of various WS2 spiral structures by SHG imaging: (a) AFM image, (b) proposed noncentrosymmetric (or symmetric stacking) structure, (c) optical image and (d) SHG imaging of a single spiral structure (a-d), a double spiral structure with a 120° included angle (e-h), a five spiral structure with 0° included angles (i-l), a complex spiral structure with a double spiral structure with a 60° included angle at the center of a single triangular spiral (m-p), which is the same structure shown in Figure 4a,h. The left row shows the AFM images of the centers of the plates, the second row illustrates the top view of the structure and the layer stacking arrangements, and the third and fourth row show the optical and SHG images of these nanoplates.

In contrast, the complex double spiral structure examined in Figure 4 displays very different SHG features. Recall that AFM image (Figure 5m) shows hexagonal spiral at the center of a triangular spiral nanoplate. Multiple dislocations are seen in the center but only one spiral arm was 18 ACS Paragon Plus Environment

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found in the outside part (Figure 5o). The SHG intensity is modest for the thinner region on the outer rim of this large triangular plate, but increases quickly toward the center at the triangular top region as the number of layers increases; however, there is a clear hexagonal hole at the very center with no SHG signal (Figure 5p). Such more complex SHG behaviors are caused by the complex multiple dislocation structures in this plate as explained in Figure 4. The hexagonal spiral in the center has an included angle of 60° between two spirals traces, which leads to the 2H (+ ̶ ) centrosymmetric layer stacking structure (Figure 5n), thus the lack of SHG signal for this hexagonal center.

In summary, we have reported the controllable growth of dislocated WS2 nanoplates and revealed the underlying growth mechanisms. By tuning the CVD reaction conditions (especially the reaction temperature) and thus the supersaturation of the crystal growth, multilayer wedding cake structures due to layer-by-layer growth, and dislocated triangular spirals with increasing numbers of dislocations per plate, and further, triangular-hexagonal mixed spirals nanostructures can be grown. Microscopic characterizations combined with SHG measurements reveal the growth mechanisms of various spiral nanostructures and the origin of the initial screw dislocations and how they lead to spiral arm features visible in optical images. We elucidate that the structural transition from triangular spirals to hexagonal spirals is enabled by the introduction new screw dislocations to result in complex multi-dislocation spirals with complex layer stacking arrangements. Since these understandings are general to layered TMD materials, screw dislocation growth provides a simple and tunable method to build various TMD nanostructures with different 19 ACS Paragon Plus Environment

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morphologies and stackings in a controllable way. We expect that the fundamental understanding of growth mechanisms of layered TMD materials at the atomic level offers promising opportunities for designing new materials for novel functional nanodevices.

ASSOCIATED CONTENT Supporting Information Available: More AFM measurements for spiral structure with alterable number and included angle of screw dislocations is presented. This material is available free of charge on the ACS Publication website at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Authors * Email: (S. J.) [email protected]; (A. P.) [email protected] Author Contributions ||

X.F. and Y.Z. contributed equally to this work.

Notes The authors declare no competing financial interest. Acknowledgement


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The authors from Hunan University are grateful to the National Natural Science Foundation of China (No.51525202, 61505051, 61574054, 61635001), the Aid program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province (2017RS3027), Y. Z. and S. J. thank the support by the U.S. Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering, under award DE-FG0209ER46664. S. J. also thanks the support from K. C. Wong Education Foundation.


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