Electric-field induced structural transition in vertical MoTe2- and Mo1–xWxTe2-based resistive memories


Transition metal dichalcogenides have attracted attention as potential building blocks for various electronic applications due to their atomically thin nature and polymorphism. Here, we report an electric-field-induced structural transition from a 2H semiconducting to a distorted transient structure (2Hd) and orthorhombic Td conducting phase in vertical 2H-MoTe2- and Mo1−xWxTe2-based resistive random access memory (RRAM) devices. RRAM programming voltages are tunable by the transition metal dichalcogenide thickness and show a distinctive trend of requiring lower electric fields for Mo1−xWxTe2 alloys versus MoTe2 compounds. Devices showed reproducible resistive switching within 10 ns between a high resistive state and a low resistive state. Moreover, using an Al2O3/MoTe2 stack, On/off current ratios of 106 with programming currents lower than 1 μA were achieved in a selectorless RRAM architecture. The sum of these findings demonstrates that controlled electrical state switching in two-dimensional materials is achievable and highlights the potential of transition metal dichalcogenides for memory applications.


Many applications, such as memristors1, micro-motors2, electronic oscillators3 and sensors4, greatly benefit from recent trends in the area of ‘phase engineering’. The most prominent materials that are explored in this context are VO2 and NbO2, which can both undergo a Mott metal-to-insulator transition5, and amorphous-to-crystalline phase change materials such as Ge2Sb2Te5 (ref. 6). Recently, transition metal dichalcogenides (TMDs) have attracted considerable attention in the field of two-dimensional (2D) phase engineering due to their polymorphism. TMDs exist in various crystalline phases, which exhibit semiconducting, semimetallic and metallic properties. In some TMD compounds the energy difference between the various phases is rather moderate7,8, so several groups are working on phase engineering in TMDs. For example, an in situ 2H to 1T phase transition in MoS2 has been introduced by means of electron beam irradiation9, and this transition has also been achieved through lithium intercalation10. However, MoS2 might not be the ideal candidate for TMD phase engineering. Because it has the lowest energy difference between the 2H and 1T′ phases among all TMDs7,8, MoTe2 appears to be the most promising compound for phase change applications. Experimental results for MoTe2 include a strain-induced semiconductor-to-metal transition11 and growth-controlled stabilization of different MoTe2 phases by choosing a specific substrate12 and through manipulation of the tellurization rate13 and growth temperature14,15,16. So far, the closest to a demonstration of electrically assisted phase change switching is a 2H to 1T′ transformation in monolayer MoTe2 induced by electrostatic liquid gating17, albeit without fabricating two-terminal electrical devices to utilize the semiconductor–metal transition. Therefore, device-compatible methods to enable reversible switching between the different crystalline phases in MoTe2 have yet to be reported.

Ultimately, for device applications, controlling the crystal structure of a TMD compound with an electric field and introducing a reproducible phase transition between two distinctly different resistive states is most desirable. Here, we experimentally demonstrate an electric-field-induced reversible structural transition in vertical devices made of MoTe2 and Mo1−xWxTe2 layers sandwiched between metal or graphene electrodes. A combination of electrical measurements with scanning transmission electron microscopy (STEM) revealed the formation of a conductive filament in the TMD layer, made of a transient structure (named 2Hd) that can be viewed as a distorted 2H state, after applying an electrical field. Temperature-dependent measurements further confirmed that the newly formed 2Hd structure exhibits electrical properties that range from semiconducting to metallic, with devices showing temperature trends that are consistent with a metallic behaviour for the lowest observed resistance values, indicating that the 2Hd state is a transient state between the semiconducting 2H phase and the metallic 1T′ or Td phase.

2H-MoTe2- and 2H-Mo1−xWxTe2-based RRAM

Figure 1a shows a schematic as well as optical microscopy and SEM images of a typical vertical TMD resistive random access memory (RRAM) device under investigation. The top contact area is 0.1 μm2. Our device design ensures that only vertical transport occurs between the two electrodes, with no lateral transport contribution. Because of the large aspect ratio between the top contact area and the flake thickness, spreading resistance contributions can be ignored and the active device area is defined by the top contact area. Area-normalized IV curves of exemplary vertical MoTe2 devices are shown in Fig. 1b. For all measurements the bottom electrode was grounded. Experimental current densities follow the expected trend with thickness. Device characteristics are reproducible and do not change substantially after multiple scans between −1 V and +1 V.

Fig. 1: Vertical TMD-based device characterization.

a, Schematic diagram of a vertical TMD device and optical and SEM images showing the top (Ti/Ni) and bottom (Ti/Au) electrodes and the SiO2 isolation layer as well as the actual flake. b, Area-normalized IV curves of vertical MoTe2 devices before electroforming for different flake thicknesses. c, IV curves of a vertical MoTe2 device from a flake with a thickness of 24 nm and a contact area of 520 nm × 330 nm. Red circles show IV curves before memristive switching. Filled black dots show the current versus voltage dependence after forming. Arrows indicate the sweep direction of the applied d.c. voltage. The current compliance is set to 400 μA.

The situation changes when the voltage range is extended. MoTe2 devices can transition into a low resistive state (LRS), as illustrated in Fig. 1c, at a set voltage (here Vset = 2.3 V). Details about the forming process are provided in the Supplementary Information. After the forming event, the device characteristics can be cycled to exhibit typical bipolar RRAM behaviour. Note that, after the forming has occurred, the high resistive state (HRS) always remains more conductive than the original state of the device, indicating that a permanent electronic change has occurred. For the case shown in Fig. 1c, the current ratio between the HRS and the LRS is ~50 at Vread = 1 V when the compliance is set to 400 μA.

To further explore the switching mechanism in TMDs, RRAM cells with exfoliated MoTe2 layer thicknesses between 6 nm and 36 nm were fabricated. All cells were non-volatile and stable (see Supplementary Fig. 2b for read disturb measurements). Before the forming process, the current per unit area through the vertical structures scales in an approximately inversely proportional manner with the flake thickness for not too small voltages (Fig. 1b). However, once the system transitions into its LRS, the current levels (below compliance) are rather similar and do not show any coherent scaling trend with the flake thickness or active device area. This observation is consistent with the notion that the formation of the conductive filaments that enable the LRS is not uniform but gives rise to local current paths. For flake thicknesses from 6 nm to 36 nm, the set voltages can be tuned from 0.9 V to 2.3 V (Fig. 2d). It is noteworthy that the RRAM behaviour is independent of the contact metal used, which indicates that the switching mechanism does not involve tunnelling through the interfacial Schottky barriers or metal ion diffusion. For example, employing Ni or graphene instead of Ti/Ni as the top electrode resulted in the same RRAM performance as reported here. In addition, a compositional energy-dispersive X-ray spectroscopy (EDS) line scan analysis through the filament area in the Ti/MoTe2/Au cross-section STEM sample shows that metal ion diffusion is not the reason for the observed RRAM behaviour in MoTe2 devices (Supplementary Section 3).

Fig. 2: 2H-MoTe2- and 2H-Mo1−xWxTe2-based RRAM behaviour and their set voltages as a function of flake thickness.

a,b, Log (a) and linear (b) scale IV curves of vertical MoTe2 RRAM devices after electroforming. The active device areas of the 7, 9 and 24 nm MoTe2 flake devices are 542 nm × 360 nm, 542 nm × 360 nm and 518 nm × 332 nm, respectively. c, Log-scale IV curves of vertical Mo0.96W0.04Te2 RRAM devices after electroforming with a current compliance of 400 μA. The active device area of the 10, 12 and 20 nm Mo0.96W0.04Te2 flake devices are 500 nm × 380 nm, 522 nm × 400 nm and 510 nm × 384 nm, respectively. d, Set voltage values scale with the flake thickness of MoTe2, Mo0.97W0.03Te2, Mo0.96W0.04Te2, Mo0.93W0.07Te2 and Mo0.91W0.09Te2. The error bars for the set voltages and the flake thicknesses are in the range of the sizes of the dots. MoTe2-Cl and MoTe2-I denote crystals grown with TeCl4 and I2 transport agents, respectively.

Next, we explored the impact of the material preparation and composition on the RRAM characteristics by extending the experiments to 2H-MoTe2 samples obtained using different synthesis approaches, and also to 2H-Mo1−xWxTe2 alloys. The similar switching characteristics seen in MoTe2 devices fabricated using either commercially obtained material or single crystals synthesized in this work by chemical vapour transport indicate that the observed RRAM effect does not depend on the material processing conditions. Mo1−xWxTe2 devices exhibit very similar switching behaviour, and their set voltages also depend monotonically on the flake thickness (Fig. 2c,d). It is noteworthy that, although we are currently unable to resolve a quantitative trend of the set voltages as a function of composition x in Mo1−xWxTe2 alloys, the set voltages for the Mo1−xWxTe2 devices show a tendency of being smaller than the set voltages of the MoTe2 devices. This implies that the critical electric field needed to trigger the RRAM behaviour may have been reduced in the alloys. Density functional theory (DFT) calculations8,18,19 and recent experimental results20,21 indicate that the energy required to transform the semiconducting state into the metallic state is increasingly lower in Mo1−xWxTe2 alloys with increasing x, which in turn should reduce the set voltage in Mo1−xWxTe2 devices as compared to their MoTe2 counterpart. Figure 2d shows experimental evidence of this trend.

Electric-field induced 2H to 2Hd transition

In general, the electroforming process in RRAM devices aims at creating a conductive filament by applying a sufficiently high electrical bias, which in turn results in an electric field and Joule heating inside the sample22. The bipolar RRAM behaviour of the MoTe2 and Mo1−xWxTe2 devices shown in Fig. 2 suggests that the main driving force for the switching is the electric field rather than Joule heating (Supplementary Section 4). To confirm the formation of conductive filaments in the case of MoTe2 RRAM cells, conductive atomic force microscopy (C-AFM) measurements were carried out. First, a fully functional MoTe2 device was biased to form the LRS, and then the top electrode was removed by wet chemical etching as described in Supplementary Section 14. This approach allows access to the TMD surface after filament formation to perform a local analysis of the surface resistivity after the forming process has occurred. As shown in the inset of Fig. 3a (indicated by an arrow), a bright spot, ~80 nm in diameter, which was formerly covered by the top electrode (red rectangle), is indicative of a higher conductivity path through the TMD layer. For comparison, those MoTe2 flakes that did not undergo a forming process show a uniform highly resistive surface (Fig. 3b).

Fig. 3: C-AFM and STEM measurements and analysis.

a, Current mapping, using C-AFM, of a MoTe2 flake after the set process and formation of the LRS. Left, topography image of the device before removal of the top electrode. Top-right, current map image after top-electrode wet-etching. The red dashed square denotes the active device area before removal of the top electrode. Note the bright spot marked with an arrow, which we interpret as a filament. Lower-right, experimental set-up schematic. b, C-AFM images of a pristine MoTe2 flake (left, topography; right, current map) showing no indication of the aforementioned highly conductive area. c, HAADF–STEM image showing cross-section of the Mo0.96W0.04Te2 device. d, Higher-magnification HAADF image from the region defined by a red box in c and showing coexistence of a distorted structure (2Hd) with 2H. e,f,g,h, Atomic-resolution HAADF images taken along the [110]2H zone axis (e,f) and [120]2H zone axis (g,h), showing the intact 2H and distorted 2Hd structures, respectively. i,j, Corresponding nanobeam diffraction pattern taken from the distorted 2Hd area, which is still indexed as the 2H structure. False colours are added to aid the eye.

To identify the exact nature of the observed filaments, STEM of cross-sectional samples was utilized for both MoTe2 and Mo1−xWxTe2 devices. Before performing the STEM analysis, RRAM devices underwent the same forming process as described above to create conducting filaments in the TMD layers. In total, more than 20 devices were carefully analysed by STEM. Figure 3c shows a high-angle annular dark field (HAADF) STEM image of a Mo0.96W0.04Te2 device cross-section. The RRAM structure is clearly visible from the HAADF contrast (which is similar to the MoTe2 device cross-section shown in Supplementary Fig. 10a,b). Note that the HAADF images display the TMD layer both in the active and non-active region, where the SiO2 isolation layer on top of the TMD is preventing RRAM operation. While only the original 2H phase is always observed in the non-active region, two structurally distinct domains can be clearly recognized in the active region; these domains are marked 2H and 2Hd, respectively, in Fig. 3d, which shows a magnified HAADF image of the red-marked box in Fig. 3c). The 2Hd region extends vertically throughout the whole TMD layer thickness in this case, but the cross-sectional view often displays the 2H and 2Hd structures simultaneously (Fig. 4a) when the cut through the device does not perfectly hit the filament core. The 2Hd region in Fig. 3d is ~80 nm wide and is separated from adjacent 2H regions by boundaries marked as white dash-dotted lines in the figure. The width of the 2Hd region is consistent with the diameter of the conducting filaments measured by C-AFM. Figure 3e shows an atomic HAADF image in the [110]2H zone axis of the 2H region, where well-resolved atomic columns of Mo/W and Te and interlayer van der Waals gaps of the hexagonal 2H structure are clearly visible. Figure 3f shows a structural HAADF image from the 2Hd domain in the corresponding orientation. Instead of the well-aligned atomic columns of the 2H structure, the atomic columns of the 2Hd structure are not clearly resolved (note that similar 2Hd structural features were observed for the MoTe2 device, as shown in Supplementary Fig. 10).

Fig. 4: STEM images and resistance of vertical MoTe2-based devices versus temperature in their respective 2H phase, HRS, LRS and 1T′ phase.

a, Atomic-resolution HAADF image taken along the [110]2H zone axis from the filament area, showing the coexistence of 2Hd and 2H regions. b, Atomic-resolution HAADF image taken along [120]2H zone axis—an orthorhombic Td phase is clearly observed together with the 2H phase. Top right inset, Corresponding fast Fourier transform image. c, Structural projections of two corresponding variants of orthorhombic Td and monoclinic 1T′ phases of MoTe2 (left) and simulated electron diffraction patterns (right). d, Semiconducting behaviour is observed for the 2H intrinsic state and for the HRS state. The 1T′/Td device at the bottom shows clear metallic behaviour, as is apparent from the decrease in sample resistance towards lower temperatures. The LRS temperature-dependent results show a gradual change in their slope, indicating a transition from semiconducting to metallic behaviour. Note that each line is obtained from an individual device that was set to the corresponding state at room temperature before the temperature-dependent measurements were performed.

In an attempt to obtain additional structural information for the 2Hd structure, the STEM sample was tilted 30° around the c axis to the [120]2H zone axis. For the 2H structure, Fig. 3g shows the well-resolved atomic columns of this phase. For the 2Hd structure, Fig. 3h shows ‘smearing’ of the Te and Mo/W atomic columns, displacement along the c direction, and preserved interlayer van der Waals gaps. Nanobeam diffraction was performed to determine the crystallography of the 2Hd structure. Figure 3i,j presents nanobeam diffraction patterns taken from the 2Hd regions in the [110]2H and [120]2H zone axes, respectively. These diffraction patterns are very similar to the corresponding pattern from the 2H region, which is surprising considering the clear differences in the atomic images of these two regions (Fig. 3e,f and 3g,h, as well as Supplementary Fig. 10d-f for the MoTe2 case). A detailed analysis of the elemental distribution of Te and Mo in the 2Hd region (see EDS scan in Supplementary Fig. 4) shows a similar Mo:Te stoichiometric ratio in the filament and pristine regions. However, due to the limitations in the STEM–EDS line-scan, such as the beam damage and sample/beam drifting issues, the rapidly acquired EDS spectra cannot completely exclude the possibility of vacancies in the filament. Moreover, because the structural images from 2Hd do not show well-separated atomic columns, and the exact arrangement of atoms in the 2Hd structure varies from device to device and for different focused ion beam (FIB) cuts, with possible interference from overlapping with the 2H matrix, it is extremely difficult to derive a precise structural model for the 2Hd structure directly from the HAADF images by relying only on the positions and average composition of the columns (Supplementary Section 9). Instead, we employed a more detailed experimental approach to better understand the observed 2Hd structure.

Most of the devices studied in the way described above exhibited similar behaviour, with the 2H structure in the matrix and the 2Hd structure in the filament. However, one device, as shown in Fig. 4b, displayed a filament consisting of the 2H phase together with the orthorhombic Td phase. This is an unexpected finding, because 1T′ MoTe2 should be more stable than the Td phase at room temperature and above (although the Td phase can be formed from the 1T′ phase at temperatures below ~240 K; ref. 23). This observation, which is substantiated by the fast Fourier transform pattern in the inset of Fig. 4b and the matching atomic structure model in Fig. 4c, is a clear indication that phases that are typically not stable in a bulk state can be stabilized in a heterostructure due to the particular boundary conditions.

This finding and the fact that we had observed a span of LRS values for different compliance settings and scanning conditions prompted us to perform temperature-dependent electrical measurements. We carefully compared the temperature-dependent electrical properties of the intrinsic 2H phase with the HRS and LRS states, as well as with reference devices fabricated from 1T′ MoTe2. The test devices were set to the HRS and LRS, respectively, at room temperature, and then characterized as a function of temperature from 160 K to 400 K at a voltage of 0.5 V. Devices kept their original state (either LRS or HRS) during the full temperature cycle. Figure 4d shows the result of our temperature-dependent characterization. Both the intrinsic 2H state and the HRS show a distinctly higher resistance at lower temperatures, which is consistent with semiconducting behaviour. On the other hand, for the LRS devices there is a clear spread in the temperature-dependent resistance slopes with a decreasing slope for decreasing sample resistance in the LRS. Interestingly, the lowest resistance values that we observed in the LRS and the corresponding slopes as a function of inverse temperature are approaching what was measured for reference 1T′ MoTe2 samples. Moreover, as shown in Supplementary Section 11, the device with the Td phase shown in Fig. 4b exhibited the lowest observed resistance values after setting. The sum of these observations in conjunction with the observed 2Hd structure and Td phase from STEM after setting provides experimental evidence that the observed 2Hd structure is a transitional structural state between 2H and 1T′ or Td (Supplementary Section 9). Devices that exhibit the lowest LRS values are electrically more similar to the 1T′ or Td phase, while devices with higher LRS exhibit a transient 2Hd state.

To explore whether the observed 2Hd structure is theoretically a (meta)stable new phase, DFT calculations were carried out based on the model proposed in Supplementary Section 9. Preliminary calculations show that the 2Hd structure is non-stable and relaxes towards the 2H phase (Supplementary Section 10). Note that this finding remains inconclusive, because the boundaries between the different structures as observed in the experiments; Figs. 3d and 4a,b) that can be expected to drastically impact the energy landscape for the various states in MoTe2 are not captured in this DFT simulation. Moreover, due to the limitation of the experiments, vacancy diffusion or other effects cannot be completely excluded for the formation of the observed 2Hd structure. Future studies are needed to investigate all plausible mechanisms.

2H-MoTe2-based RRAM under pulsed operation

Pulse measurements were used to explore the switching speed to toggle between the two states in MoTe2. Figure 5a shows the applied set and reset voltages across the device as a function of time. The set voltage of this device is 1 ± 0.1 V (Supplementary Fig. 6), and the effective set pulse width is accordingly smaller than 10 ns. Notice that the uncompensated parasitic capacitance contributions of the set-up did affect the shape of the pulse curve. From Fig. 5b it is apparent that the device resistance changed from the HRS to the LRS after the pulse set operation, and increased again after the reset pulse, in accordance with our expectations. Note that the current compliance settings in the continuous sweep and pulse modes are unavoidably different, resulting in different HRSs before set and after reset in Fig. 5b. This confirms that changing the 2H state into the 2Hd state can be achieved with appropriate voltage pulses of less than 10 ns. This value only represents an upper bound for the switching time since the experimental set-up did not permit performing faster pulse measurements. The intrinsic switching speed is believed to be even higher than this initial demonstration.

Fig. 5: Performance of 2H-MoTe2-based RRAM under pulsed operation.

a, Device voltage during set/reset operations of a device with 8 nm flake thickness. The set voltage is 1 ± 0.1 V and the effective set pulse width is less than 10 ns. The reset voltage is larger than −0.8 V and the effective reset pulse width is about 100 ns. b, Device read current at different states.

Al2O3/MoTe2-based low-programming-current RRAMs

For device applications, the LRS resistance should be sufficiently large that the voltage drop across the interconnects is negligible to ensure proper write/read operation. In addition, the current ratio between the LRS and HRS has to be as large as possible to allow the fabrication of sizable memory arrays. In this context the above-mentioned performance specifications are not ideal, and a modified TMD-based memory cell needs to be developed. To limit the current through the cell, a tunnelling barrier (Al2O3) was added into the stack.

Figure 6a shows multiple IV measurements on an Al2O3/MoTe2 RRAM cell, indicating a stable and reproducible memristive switching of this device. Unlike the previously discussed MoTe2-only RRAM cells, modified Al2O3/MoTe2 stacks immediately transitioned into an LRS when the applied bias reached 2.9 V, without the previously observed electroforming process. Moreover, no current compliance through external circuitry is needed—another desirable feature in RRAM cells. The stability of the RRAM cell is further underlined in Fig. 6b. The extended stack also gave rise to a much larger current ratio between the LRS and HRS of 105–106 as well as nonlinearities (IVoperation/I1/2Voperation based on a V/2 scheme) of ~100 for the LRS, and a resistance in the HRS state larger than 10 TΩ (limited by our measurement equipment). The latter is critical for reduced static leakage power consumption and allows the maximum RRAM array size to be increased. Moreover, the low current level in the HRS and LRS may eliminate the need for selector devices24. All of the above highlights the relevance of this structure for sub-1 μA programming currents for selectorless RRAM applications.

Fig. 6: Performance of 2H-Al2O3/MoTe2-based RRAM.

a, Multiple IV measurements on a vertical metal–Al2O3/MoTe2–metal RRAM cell with a flake thickness of 12 nm and an area of 360 nm × 390 nm. The set voltage is ~2.9 ± 0.25 V (shaded pink) and the reset voltage is ~−1.75 ± 0.25 V (shaded blue). A clear memristive behaviour is visible, with the arrows indicating the sweep direction. b, Read disturb measurement of the same RRAM cell at 2 V at room temperature (shaded purple in a).

Notably, Al2O3 is one of the early materials investigated for RRAM applications. However, in our preliminary testing of vertical Ti/Au/Al2O3/Ti/Ni device structures, repeated voltage sweeps did not result in any filament formation and instead led to destructive breakdown at voltages in the 1.35 V range (Supplementary Fig. 17). While we cannot entirely exclude the possibility of filament formation in the Al2O3 layer of the Al2O3/MoTe2 stack, the electrical data can be qualitatively understood if filaments are only formed in MoTe2, but do not reach through the Al2O3 layer, in this way limiting the LRS conductivity to a much lower, more desirable value of ~10 MΩ as compared to the MoTe2-only RRAM cells. In fact, the current levels reached in the LRS of the Al2O3/MoTe2 RRAM cell are consistent with the maximum current levels achievable before breakdown in the Ti/Au/Al2O3/Ti/Ni stack.


In this Article, an electric-field-induced reversible transition from the semiconducting 2H to a higher conducting 2Hd phase in vertical MoTe2 and Mo1–xWxTe2 RRAM devices has been achieved. Cross-sectional STEM confirmed that the 2Hd state is formed after a set voltage is applied and that this structure can be described as a transient state between the semiconducting 2H and the metallic 1T′ or Td phase. Temperature-dependent electrical measurements indicate that the 2Hd state can exhibit semiconducting to metallic behaviour depending on the status of the transient state. Td and 2H co-existing phases in MoTe2 were observed at room temperature. Our work demonstates the possibility to locally and selectively engineer phase transitions in TMD layers with an electric field, and demonstrates the potential of TMDs for these types of resistive switching applications.


Device fabrication and electrical measurements

A layer of Ti/Au (10 nm/25 nm;acting as a bottom electrode) was deposited onto a 90 nm silicon dioxide (SiO2) layer located on top of a highly doped silicon wafer. Next, TMD flakes from either MoTe2 (2D Semiconductors and NIST) or Mo1–xWxTe2 (NIST) were exfoliated onto this electrode using a standard scotch tape technique, followed by thermal evaporation of 55 nm SiO2 (acting as an insulating layer). Device fabrication was completed by depositing a Ti/Ni (35 nm/50 nm) top electrode. For Ti/Au/Al2O3/MoTe2/Ti/Ni vertical devices, a 3-nm-thick Al layer was deposited onto the bottom electrode before its oxidation in an oxygen-rich environment at 250 °C for 6 h. This process formed a 4.5-nm-thick Al2O3 layer. MoTe2 flakes were then peeled onto the Al2O3-covered bottom electrodes. The isolation and top electrode formation were identical to the original process flow discussed above for vertical TMD RRAM devices.

Electrical characterization of the devices was performed at room temperature using a parameter analyser (Agilent 4156C). A Keysight B1500A semiconductor device analyser was used for switching speed measurements.

TMD synthesis and characterization

Both 1T′- and 2H-Mo1–xWxTe2 crystals (x = 0, 0.03, 0.04, 0.07 and 0.09, where x is the atomic fraction of W) were produced at NIST using the chemical vapour transport method. First, polycrystalline Mo1–xWxTe2 powders were synthesized by reacting stoichiometric amounts of Mo (99.999%), W (99.9%) and Te (99.9%) at 750 °C in a vacuum-sealed quartz ampoule. Next, Mo1–xWxTe2 crystals were grown at 950–1,000 °C using ~1 g of poly-Mo1–xWxTe2 charge and a small amount of I (99.8%, 5 mg cm−3) sealed in evacuated quartz ampoules. The ampoules were ice-water-quenched after 7 days of growth, yielding Mo1–xWxTe2 crystals in the metallic 1T′ phase. The 1T′-Mo1–xWxTe2 alloy crystals were then converted to the semiconducting 2H phase by annealing in vacuum-sealed ampoules at 950 °C for 24 h (or at 750 °C for 72 h) followed by cooling to room temperature at a 10 °C h−1 rate. 2H-MoTe2 crystals were also obtained by chemical vapour transport growth at 800 °C for 140 h using TeCl4 (99.9%, 5.7 mg cm−3) as a vapour transport agent. At this temperature, which, in accordance with the Mo-Te phase diagram25 is below the 1T′→2H phase transition temperature, MoTe2 crystals grow directly in the 2H phase.

The crystal structure and composition of the Mo1–xWxTe2 samples were determined by powder X-ray diffraction and energy-dispersive X-ray spectroscopy in SEM, respectively. More detailed information on crystal preparation and characterization can be found in ref. 20.

Conductive AFM measurement

C-AFM was performed in a Veeco Dimension 3100 AFM system. All AFM images were taken in contact mode using SCM-PIT tips (Bruker). The conductive tip consisted of 0.01–0.025 Ω cm antimony-doped Si coated with PtIr.

TEM/STEM structural characterization

An FEI Nova NanoLab 600 DualBeam (SEM/FIB) system was used to prepare cross-sectional TEM samples. Electron-beam-induced deposition of 1-μm-thick carbon was initially deposited on top of the device to protect the sample surface, followed by 2-μm ion-beam induced Pt deposition. To reduce Ga ion damage, in the final step of preparation the TEM samples were thinned with 2 kV Ga ions using a low beam current of 29 pA and a small incident angle of 3°. An FEI Titan 80–300 probe-corrected STEM/TEM microscope operating at 300 keV was used to acquire both nanobeam diffraction patterns and TEM images in TEM mode as well as atomic-resolution high-angle annular dark field (HAADF) images in STEM mode.

Data availability

The data that support the plots within this paper are available from the corresponding author upon request.


  1. 1.

    Wuttig, M. & Yamada, N. Phase-change materials for rewriteable data storage. Nat. Mater. 6, 824–832 (2007).

    CAS  Article  Google Scholar 

  2. 2.

    Liu, K. et al. Powerful, multifunctional torsional micromuscles activated by phase transition. Adv. Mater. 26, 1746–1750 (2014).

    Article  Google Scholar 

  3. 3.

    Gu, Q., Falk, A., Wu, J., Ouyang, L. & Park, H. Current-driven phase oscillation and domain-wall propagation in WxV1 –xO2 nanobeams. Nano Lett. 7, 363–366 (2007).

    CAS  Article  Google Scholar 

  4. 4.

    Strelcov, E., Lilach, Y. & Kolmakov, A. Gas sensor based on metal−insulator transition in VO2 nanowire thermistor. Nano Lett. 9, 2322–2326 (2009).

    CAS  Article  Google Scholar 

  5. 5.

    Zhou, Y. & Ramanathan, S. Mott memory and neuromorphic devices. Proc. IEEE 103, 1289–1310 (2015).

    CAS  Article  Google Scholar 

  6. 6.

    Wong, H. S. P. et al. Phase change memory. Proc. IEEE 98, 2201–2227 (2010).

    Article  Google Scholar 

  7. 7.

    Duerloo, K.-A. N., Li, Y. & Reed, E. J. Structural phase transitions in two-dimensional Mo- and W-dichalcogenide monolayers. Nat. Commun. 5, 4214 (2014).

    CAS  Article  Google Scholar 

  8. 8.

    Duerloo, K.-A. N. & Reed, E. J. Structural phase transitions by design in monolayer alloys. ACS Nano 10, 289–297 (2016).

    CAS  Article  Google Scholar 

  9. 9.

    Lin, Y.-C., Dumcenco, D. O., Huang, Y.-S. & Suenaga, K. Atomic mechanism of the semiconducting-to-metallic phase transition in single-layered MoS2. Nat. Nanotech. 9, 391–396 (2014).

    CAS  Article  Google Scholar 

  10. 10.

    Py, M. A. & Haering, R. R. Structural destabilization induced by lithium intercalation in MoS2 and related compounds. Can. J. Phys. 61, 76–84 (1983).

    CAS  Article  Google Scholar 

  11. 11.

    Song, S. et al. Room temperature semiconductor–metal transition of MoTe2 thin films engineered by strain. Nano Lett. 16, 188–193 (2016).

    CAS  Article  Google Scholar 

  12. 12.

    Tsipas, P. et al. Direct observation at room temperature of the orthorhombic weyl semimetal phase in thin epitaxial MoTe2. Adv. Funct. Mater. 28, 1802084 (2018).

    Article  Google Scholar 

  13. 13.

    Park, J. C. et al. Phase-engineered synthesis of centimeter-scale 1T′- and 2H-molybdenum ditelluride thin films. ACS Nano 9, 6548–6554 (2015).

    CAS  Article  Google Scholar 

  14. 14.

    Vellinga, M. B., de Jonge, R. & Haas, C. Semiconductor to metal transition in MoTe2. J. Solid State Chem. 2, 299–302 (1970).

    CAS  Article  Google Scholar 

  15. 15.

    Keum, D. H. et al. Bandgap opening in few-layered monoclinic MoTe2. Nat. Phys. 11, 482–486 (2015).

    CAS  Article  Google Scholar 

  16. 16.

    Empante, T. A. et al. Chemical vapor deposition growth of few-layer MoTe2 in the 2H, 1T′ and 1T phases: tunable properties of MoTe2 films. ACS Nano 11, 900–905 (2017).

    CAS  Article  Google Scholar 

  17. 17.

    Wang, Y. et al. Structural phase transition in monolayer MoTe2 driven by electrostatic doping. Nature 550, 487–491 (2017).

    CAS  Article  Google Scholar 

  18. 18.

    Li, Y., Duerloo, K. A., Wauson, K. & Reed, E. J. Structural semiconductor-to-semimetal phase transition in two-dimensional materials induced by electrostatic gating. Nat. Commun. 7, 10671 (2016).

    CAS  Article  Google Scholar 

  19. 19.

    Zhang, C. et al. Charge mediated reversible metal–insulator transition in monolayer MoTe2 and WxMo1–xTe2 alloy. ACS Nano 10, 7370–7375 (2016).

    CAS  Article  Google Scholar 

  20. 20.

    Sean, M. O. et al. The structural phases and vibrational properties of Mo1−xWxTe2 alloys. 2D Mater. 4, 045008 (2017).

    Article  Google Scholar 

  21. 21.

    Lv, Y.-Y. et al. Composition and temperature-dependent phase transition in miscible Mo1−xWxTe2 single crystals. Sci. Rep. 7, 44587 (2017).

    Article  Google Scholar 

  22. 22.

    Wong, H. S. P. et al. Metal-oxide RRAM. Proc. IEEE 100, 1951–1970 (2012).

    CAS  Article  Google Scholar 

  23. 23.

    Berger, A. N. et al. Temperature-driven topological transition in 1T′-MoTe2. NPJ Quantum Mater. 3, 2 (2018).

    Article  Google Scholar 

  24. 24.

    Liang, J. & Wong, H. S. P. Cross-point memory array without cell selectors—device characteristics and data storage pattern dependencies. IEEE Trans. Electron. Dev. 57, 2531–2538 (2010).

    Article  Google Scholar 

  25. 25.

    Brewer, L. & Lamoreaux, R. H. in Binary Alloy Phase Diagrams 2nd edn, Vol. 3 (eds. Massalski, T. B., Okamoto, H., Subramanian, P. R. & Kacprzak, L.) 2675–2676 (ASM International, Russell Township, 1990).

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This work was supported in part by the Semiconductor Research Corporation (SRC) at the NEWLIMITS Center and National Institute of Standards and Technology (NIST) through award no. 70NANB17H041. S.K. acknowledges support from the US Department of Commerce, NIST under financial assistance award 70NANB16H043. H.Z. acknowledges support from the US Department of Commerce, NIST under financial assistance awards 70NANB15H025 and 70NANB17H249. A.V.D., S.K. and B.P.B. acknowledge support from Material Genome Initiative funding allocated to NIST. The authors thank I. Kalish (NIST) for conducting XRD and EDS measurements on Mo1–xWxTe2 samples. Certain commercial equipment, instruments or materials are identified in this paper to specify the experimental procedure adequately. Such identification is not intended to imply recommendation or endorsement by NIST, nor is it intended to imply that the materials or equipment identified are necessarily the best available for the purpose.

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F.Z. and J.A. designed the experiments. F.Z. fabricated, measured the devices and performed the conductive AFM measurement. S.K. and A.V.D. synthesized Mo1–xWxTe2 alloy samples. C.A.M. and D.Y.Z. performed the STM and STS measurements. D.Y.Z. contributed the STM surface analysis. H.Z. prepared TEM samples using SEM/FIB and performed TEM/STEM measurements. H.Z., L.A.B. and A.V.D. performed the TEM/STEM analysis. B.P.B. conducted ab initio modelling of energetics for the 2H, 2Hd and 1T′ MoTe2. Y.Z. and J.A. carried out the model simulation for vertical electrical transport. F.Z., H.Z., L.A.B., A.V.D. and J.A. wrote the manuscript and discussed the results at all stages.

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Correspondence to Huairuo Zhang or Joerg Appenzeller.

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Sections 1–14, Supplementary References 1–14, Supplementary Table 1, Supplementary Figures 1–17

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Zhang, F., Zhang, H., Krylyuk, S. et al. Electric-field induced structural transition in vertical MoTe2- and Mo1–xWxTe2-based resistive memories. Nature Mater 18, 55–61 (2019). https://doi.org/10.1038/s41563-018-0234-y

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