In-plane template induced selective growth of H-phase CrSe2

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In-plane template induced selective growth of H-phase CrSe2

https://doi.org/10.21203/rs.3.rs-3456540/v1

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Phase engineering in two-dimensional transition metal dichalcogenides (2D-TMDs) offers opportunities for exploring their unique properties of different phases and achieving new desired functionalities. Here, we report a selective synthetic strategy for the molecular beam epitaxy growth of the H-phase CrSe₂ induced by an in-plane heteroepitaxial template. The lattice-matched MoSe₂ nanoribbons are utilized as the in-plane heteroepitaxial template to seed the growth of H-phase CrSe₂ thin layers with the consequential formation of MoSe₂-CrSe₂ lateral heterostructures. Scanning tunneling microscopy and non-contact atomic force microscopy studies reveal the atomically sharp interfaces in the lateral heterostructures and the characteristic defects of mirror twin boundaries emerging in the H-phase CrSe₂monolayers. Distinct from the metallic T-phase CrSe₂, the semiconducting character of the H-phase CrSe₂ is detected using the scanning tunneling spectroscopy. The type-I straddling band alignments of MoSe₂-CrSe₂ lateral heterostructures are directly visualized in atomic precision with band bending at the interfaces. The mirror twin boundaries in the H-phase CrSe₂ exhibit the Tomonaga-Luttinger liquid behavior in the confined one-dimensional electronic system. This in-plane heteroepitaxial template induced phase-selective growth provides a promising approach for phase engineering of 2D TMDs, thereby further expanding the library of crystal phases and promoting the versatile device applications of specific phases.

Physical sciences/Materials science/Nanoscale materials/Two-dimensional materials

Physical sciences/Materials science/Nanoscale materials/Electronic properties and materials

2D TMDs possess various polymorphic structures including 2H (trigonal prismatic), 1T (octahedral), 1T′ and T_d phases, which are determined by the different chalcogen coordination geometries of metal atoms and stacking orders of atomic layers¹. The polymorphic phases give rise to a variety of intriguing properties such as superconductivity^2,3, ferroelectricity⁴, ferromagnetism^5,6, Weyl semimetal⁷ and quantum spin Hall effect⁸. The physical and chemical properties of TMDs can be modulated by changing the structural phases. The so-called phase engineering of TMDs makes it possible to achieve the precise control of structural phases in a desired manner^9-11. The phase engineering can be implemented through two main strategies of phase transformation and phase-selective synthesis. The phase transformation strategies including alkali metal intercalation^12,13, alloying¹⁴, thermal treatment^15,16, electric field¹⁷, mechanical strain^18,19, and external irradiation^20,21have been employed to trigger the structural transformation to effectively obtain the target phases. The aforementioned chemical and physical methods might encounter the issue of impurities and defects introduced during the phase transformation process as well as the metastable phases recovering to their original structures when the external stimuli are removed²². In comparison, phase-selective synthesis has the ability to directly fabricate 2D TMDs with the specific phase and high phase purity. The strategies of precursor design, atmosphere regulation and temperature control have been utilized in chemical vapor deposition (CVD) growth to achieve the phase-selective synthesis^23-26. As for the MBE growth, it is highly desirous to develop general and controllable strategies to fabricate 2D TMDs with excellent phase selectivity.

Heteroepitaxial templates have been employed to seed the growth of crystals with unconventional phases^27-30. The heteroepitaxial template usually provides a constrained environment to guide the crystal growth with the formation of lattice-matched interfaces. Under the effect of heteroepitaxial templates, a stabilized nonequilibrium phase can be formed within a critical coherent thickness, for instance, the metastable cubic phase of aluminium nitride (AlN) was stabilized in AlN/TiN superlattice films^27,31. The formation of a lattice-matched interface with a lower interfacial energy is the essential prerequisite for the stable existence of metastable phases. The heteroepitaxial templates play a prominent role in phase control of noble metal nanocrystals, for example, the unconventional 4H-phase Ag, Pd, Pt, Ir, Rh, Os, Ru and Cu were synthesized using 4H-Au nanoribbons as the epitaxial templates^30,32. In two-dimensional TMD materials, due to the existence of atomically flat dangling-bond-free surfaces, in-plane heteroepitaxial template effects are considered to seed the growth of 2D TMDs for achieving the selective growth of metastable phase structures in MBE growth.

Selective growth of H-phase CrSe₂ with the formation of lateral heterostructures

Atomically thin CrSe₂ films are usually observed in the metallic 1T phase which is demonstrated by atomic-resolution transmission electron microscopy and other experiments^6,33-35. As shown in the schematic illustration (Fig. 1a), the codeposition of Cr and Se atoms on a substrate of highly oriented pyrolytic graphite (HOPG) will produce the metallic T-phase CrSe₂ islands. The T-phase CrSe₂has an octahedral coordination structure with D_3d symmetry and degenerate and orbitals. The partially filled d orbitals give rise to the metallic electronic properties as illustrated in Fig. 1b. By utilizing the in-plane heteroepitaxial template of H-phase MoSe₂ nanoribbons, the growth of H-phase CrSe₂ monolayers can be guided with minimizing the interfacial energy (Fig. 1a). The atomically matched lattices guarantee the formation of lateral heterostructures with seamlessly connected interfaces. Unlike metallic T-phase CrSe₂, the H-phase CrSe₂ adopts the trigonal prismatic coordination structure with D_3hsymmetry, which splits the d orbitals into three groups , and with a sizeable bandgap (Fig. 1c).

In the STM image of the MBE-grown CrSe₂ samples (Fig. 1d), the T-phase CrSe₂islands exhibit approximately hexagonal shapes with straight edges. A nonlayered growth behavior is unraveled by the existence of CrSe thin layer with a height of ~0.3 nm shown in Fig. 1d. The lattice constant is calibrated to be 3.4 ± 0.1 Å which is consistent with that of previously reported 1T-CrSe₂ crystals (3.39 Å) synthesized through deintercalation of lithium from LiCrSe₂³⁵. The metallic characteristic of the T-phase CrSe₂monolayer is unveiled in the differential conductance (dI/dV) spectrum (Fig. 1g). In order to achieve the phase-selective growth of 1H CrSe₂, the in-plane epitaxial template of MoSe₂ nanoribbons are prepared beforehand with the typical width of 10 − 50 nm at a growth temperature of about 550 °C (Fig. 1e). The distinctive Mo- and Se-edges at the opposite sides of nanoribbons can be distinguished by different edge features in the atom-resolved STM images (Supplementary Fig. 1). After the growth of MoSe₂ nanoribbons, Cr and Se atoms are subsequently codeposited onto the same HOPG substrate. The diffusing atoms preferentially nucleate and aggregate at the active edges of MoSe₂ nanoribbons. The H-phase CrSe₂ rather than T-phase structure is formed with the lattice matching under the effect of the in-plane epitaxial template of MoSe₂ nanoribbons. The as-grown H-phase CrSe₂ segments are seamlessly fused to MoSe₂ nanoribbons with the formation of lattice-matched lateral heterostructures (Fig. 1f). In the lateral heterostructures, the continuous H-phase interface structures have a lower interfacial energy compared with the 1H-1T interface structures as revealed by density functional theory (DFT) calculations (Supplementary Fig. 2). The chemical states of Cr and Se elements in two distinct phases of CrSe₂ were investigated by X-ray photoelectron spectroscopy. The measured binding energies of Cr 2p and Se 3d electrons (Fig. 1h) show an energy shift of ~0.5 eV between H- and T-phase CrSe₂ due to the different electronic structures. The electronic properties of H-phase CrSe₂ are revealed by STS measurements which exhibit a semiconducting character with bandgap of 0.75 ± 0.05 eV. The increased peak intensity of local electronic states with the invariant bandgap can be observed in the dI/dV spectra taken at different tunneling currents (Fig. 1i). The semiconducting property of H-phase CrSe₂ monolayer is also verified by DFT calculations. The electronic structure of H-phase CrSe₂ manifests a direct bandgap of 0.72 eV with the valence band maximum and conduction band minimum derived mostly from the Cr 3d orbitals (Supplementary Fig. 2).

Structural characterization of MoSe₂-CrSe₂ interfaces

The atomically sharp interfaces of MoSe₂-CrSe₂ lateral heterostructures are revealed both by STM and nc-AFM. In the STM image of the MoSe₂-CrSe₂ lateral hetetrostructure (Fig. 2a), CrSe₂ segments are seamlessly connected to the MoSe₂ nanoribbons in between. The different STM image contrast between CrSe₂ and MoSe₂ primarily originates from the local electronic states rather than topographic features, which varies with the applied bias voltages (Supplementary Fig. 3). There is a continuous linear defect of mirror twin boundary (MTB) crossing through the heterostructure interface, reflecting the in-plane heteroepitaxial template effect. The occurrence of MTB characteristic defects in CrSe₂ provides strong evidence for the claim of H-phase CrSe₂. The atomic structures and electronic properties of MTBs in CrSe₂ will be further elaborated in more detail later. The corresponding ball-and-stick model of the lateral heterostructure with a MTB line defect is exhibited in Fig. 2b. In the atomic row of MTB, four-fold rings shares a point at chalcogen site leading to each Se atom bound to four Cr atoms instead of three. The Mo- or Se-edge of the MoSe₂ nanoribbon can be determined by the orientation of MTB due to the constraint of mirror symmetry. To detect the strain at the heterostructure interfaces, the moiré pattern in the MoSe₂ region was studied since it can be used as a ‘magnifying glass’ to directly visualize the lattice-misfit strain at the heterostructure interfaces³⁶. In the high-resolution STM image (Fig. 2c), the MoSe₂region exhibits an undistorted 3 × 3 moiré pattern, indicating that there is no obvious lattice-misfit strain at the heterostructure interface. The nc-AFM technique was also employed to characterize the interface structures as it can provide the surface topography without being disturbed by electronic states. The constant-height nc-AFM frequency shift images (Fig. 2d and Fig. 2e) taken at Se- and Mo-edges of MoSe₂ nanoribbons respectively (labelled in Fig. 2a) both exhibit atomically sharp interface structures without any dislocations. The nc-AFM observations also revealed the atomically sharp MoSe₂-CrSe₂ interface structures at the adjacent Mo-edge and Se-edge with the included angle of 120° (Supplementary Fig. 4). This results imply that CrSe₂ in the lateral heterostructure adopts the same H-phase structure as MoSe₂. If we postulate that it were T-phase CrSe₂ connected to the different edges of MoSe₂ nanoribbons, dislocations and lattice misfit strain would emerge at the interfaces due to the different lattice symmetries and lattice constants²¹, however, they are not observed in our experiments. In the nc-AFM images, the MoSe₂ regions have a brighter image contrast (higher frequency shift) compared with CrSe₂ regions, which can be ascribed to the different Pauli repulsive forces at small tip-sample distances. We adopted the model of MoSe₂-CrSe₂ lateral heterostructure with continuous H-phase structures at the interfaces to simulate the STM and nc-AFM images. The observed STM and nc-AFM image contrast between the MoSe₂ and CrSe₂ regions are reproduced well by the DFT simulation (Fig. 2f).

Band alignments in the lateral and vertical heterostructures

STS measurements were carried out to explore the spatial evolution of electronic structures of both lateral and vertical heterostructures at the atomic scale, which are of fundamental importance for understanding the interface properties and evaluating their application potentials. The close-up STM image (Fig. 3a) displays the MoSe₂-CrSe₂lateral heterostructures with 12 nm-width MoSe₂ nanoribbon in between. The type-I band alignment with a straddling gap is found for the MoSe₂-CrSe₂ lateral heterostructures, which endows the lateral heterostructures with the potential for light-emitting applications and studying light-matter interaction. The band profile across the interface is directly visualized by the colour rendering of the dI/dV mapping (Fig. 3b). A bandgap of ~2.30 eV with Fermi level located 1.60 eV above the valence band maximum (n-type) is detected in the MoSe₂ region away from the interface. The MoSe₂ region exhibits upward band bending for both valence and conduction band near the interface due to the charge transfer from MoSe₂ to CrSe₂. The band profile at the interface shows a sharp transition into the CrSe₂ region which has a smaller bandgap of ~0.75 eV. A slight downward band bending can be seen in the magnified band profile of CrSe₂ region (Fig. 3c), which is attributed to electron injection from MoSe₂. The valence and conduction band offset (VBO and CBO) in the MoSe₂-CrSe₂ lateral heterostructures are determined to be ~0.84 eV and ~0.78 eV, respectively. Besides the band bending, the effect of bandgap narrowing emerges at the heterostructure interface probably due to strain at the interface. As illustrated in Fig. 3d, the downward shift of the conduction band is larger than that of valence band, resulting in a bandgap narrowing of ~0.06 eV for CrSe₂. The inconsistent upward band shifts in valence and conduction bands of MoSe₂ also lead to a bandgap narrowing of ~0.10 eV.

Owing to layer-by-layer growth mode, the CrSe₂/MoSe₂ heterostructures are formed with the first-layer MoSe₂ and second-layer CrSe₂ in the bilayer structures (Fig. 3e). This is representative of a CrSe₂/MoSe₂vertical heterostructure. The spatially resolved dI/dV spectra are acquired along the path with a step size of 8.5 Å labelled in Fig. 3e. The band offsets and edge states at the interface can be intuitively observed in the energy profile of the CrSe₂/MoSe₂vertical heterostructure (Fig. 3f). The valence and conduction bands of the first-layer MoSe₂ exhibits a joint upward band bending at the interface with a magnitude of ~0.14 eV and ~0.07 eV, respectively. Compared with the small band bending in the previously reported bilayer-monolayer MoSe₂ and WSe₂ homostructures³⁷, the band shifts in the first-layer MoSe₂ indicate a stronger interlayer interaction in the vertical CrSe₂/MoSe₂ heterostructures. Regarding the second-layer CrSe₂ region, the valence and conduction bands are bended upward by ~0.28 and ~0.36 eV, respectively (magnified in Supplementary Fig. 5). The edge state with a narrow gap (marked by orange lines in Fig. 3f) emerges at the edge termination of second-layer CrSe₂due to the existence of dangling bonds. The STS mapping is further sliced into selected individual spectra obtained at the vicinity of interface (Fig. 3g). In the dI/dV spectra of second-layer CrSe₂, the original characteristic peak around −0.6 eV turns to be a sloping bump with energy position ~0.18 eV lower than that of monolayer CrSe₂ on the semimetallic HOPG substrate. The increase of bandgap in the second-layer CrSe₂ can be attributed to the better electronic isolation provided by the underlying MoSe₂ interlayer. At the heterostructure interface, the localized dangling-bond states give rise to the new spectral features with smaller bandgaps in the dI/dV curves. The edge states and band shifts in the vertical heterostructure are further verified by our DFT-calculated local density of states (LDOS) plots of the Se atoms on the surface (Fig. 3h). The calculated differential charge density reveals the charge accumulation at the edge termination of second-layer CrSe₂ which induces the upward band bending of CrSe₂ near the edge.

Electronic properties of MTBs in the H-phase CrSe₂ monolayers

MTBs as quasi-one-dimensional metals provide an ideal platform for exploring the electronic behavior in the confined system. MTBs with a 4|4P structure have been extensively studied in H-phase molybdenum dichalcogenides such as MoSe₂ and MoTe₂^38-41, but rarely detected in other non-Mo-based TMD materials. Isolated MTBs with triangular shape (Fig. 4a) are obtained in the H-phase CrSe₂monolayers. The close-up high-resolution STM image of MTB (Fig. 4b) exhibits a typical feature of two bright parallel lines as previously observed in MoSe₂ and MoTe₂,but without an obvious spatial modulation period. The structure model of the MTB is shown in Fig. 4c and the corresponding atomic structure is revealed by the nc-AFM image (Fig. 4d), in which a surface feature that a row of Se atoms have a darker contrast than other Se atoms in the region can be observed. The darker contrast mainly results from the higher density of Cr atoms in the MTB which gives rise to stronger attractive forces. The dI/dV spectra (Fig. 4e) taken on the MTB with a length of ~12.8 nm exhibit a feature of gap opening at the Fermi level. The gap size increases with decreasing the length of MTBs, which is one of the signatures of Tomonaga-Luttinger liquid (TLL) behavior (Supplementary Fig. 6). In the magnified dI/dV spectrum (inset in Fig. 4e), four pronounced peaks located at 37.6 mV, 15 mV, −12 mV and −34.2 mV can be identified. The constant-height dI/dV conductance mappings (Fig. 4g) measured at the four different peak energies reveal a periodic charge density modulation at the period of ~3a (triple lattice constant of CrSe₂). The modulation period is also observed in the low-bias constant-height STM image (Fig. 4f), which reflects the spatial distribution of low-energy electronic states. As shown in Fig. 4g, the spatial modulation of highest occupied state (HOS) and lowest unoccupied state (LUS) is in phase with a symmetric two-lobe feature in the MTB, but it is partially out-of-phase for the second occupied and unoccupied states. The observed charge density modulation may not be a consequence of charge density waves (CDW) but the quantum confinement of electrons in the finite-length MTBs. In the dI/dV spectra along the MTB (Fig. 4h), the periodic fluctuation of peak positions at the negative bias from −13.2 mV to −5.7 mV and at the positive bias from 11.3 mV to 15.4 mV for the HOS and LUS respectively can be discerned. And the discrete quantized states below HOS and above LUS are directly visualized in the color plot of dI/dV spectra along the MTB (Fig. 4i). The phenomena may be described by TLL theory that low-energy excitations are standing waves with discrete energies in a finite-length MTB system⁴². Moreover, the TLL behavior of spin-charge separation is observed in the Fourier transform of the 2D STS map (Fig. 4j), in which two linear dispersion branches corresponding to the spin and charge density excitations have different slopes (marked with blue and yellow lines)⁴⁰.

In this work, we successfully achieved the in-plane template induced selective growth of H-phase CrSe₂with the formation of MoSe₂-CrSe₂ lateral and MoSe₂/CrSe₂ vertical heterostructures. The effect of in-plane heteroepitaxial template plays a decisive role in the phase-selective growth of 1H CrSe₂ with minimizing the interfacial energy by lattice-matching. The atomically sharp interfaces in the MoSe₂-CrSe₂ lateral heterostructures and the characteristic defect of MTBs in the CrSe₂ monolayer are unveiled by the atomic-resolution STM and nc-AFM, verifying the H-phase structure in CrSe₂. The same phase structure and lattice constant of the H-phase CrSe₂ and MoSe₂ ensure that no obvious lattice-misfit strain exists at the interfaces of MoSe₂-CrSe₂ lateral heterostructures. The H-phase CrSe₂ has the semiconducting character with bandgap of 0.75 ± 0.05 eV for the monolayer on HOPG substrate and an increase of bandgap to a magnitude of ~0.93 eV for the second-layer CrSe₂ on the MoSe₂interlayer. Visualization of band profiles of the lateral and vertical heterostructures allows the identification of band alignment and band bending at the interfaces. In addition, the MTBs in H-phase CrSe₂ exhibit the quantum confined TLL behavior, which includes the charge density modulation, length-dependent band-gap opening and spin-charge separation. Our study paves the way for phase-selective growth of 2D-TMDs by in-plane heteroepitaxial templates, enabling further research and device applications on 2D TMDs.

Experimental measurement.

The MoSe₂-CrSe₂ heterostructures were fabricated by MBE growth in an ultrahigh vacuum chamber (base vacuum 8 \(\times {10}^{-10}\) mbar). The HOPG substrate was cleaved and then transferred into ultrahigh vacuum chamber to degas at 600 °C. Molybdenum and chromium atoms were evaporated from the electron-beam evaporators. The excessive selenium atoms were sublimed at the temperature of 160 °C to maintain the Se-rich environment. MoSe₂ nanoribbons are formed by codeposition of Mo and Se atoms with the substrate temperature kept at 550 °C. The heterostructures can be obtained through the subsequent growth of CrSe₂ at the substrate temperature of about 200 °C. The STM and nc-AFM measurements were performed on an Omicron low-temperature STM system operated at 78 K. The STM images were acquired in the constant-current mode and the nc-AFM measurement was carried out in constant-height frequency modulation mode with a native tungsten tip without CO functionalization. The AFM simulation is based on the online modelling software provided by Hapala et al⁴³. Differential conductance (dI/dV) spectra of MTBs were acquired at 4.3 K measured by a lock-in detection with a 963 Hz and 1 − 5 mV modulation superimposed on the sample bias. WSxM software was used to process all STM and AFM images⁴⁴.

Theoretical calculation.

Spin-polarized density functional theory calculations were performed using the projector augmented wave method for valence-core interactions, Perdew-Burke-Ernzerhof (PBE) functional of generalized gradient approximation, and a plane-wave basis set as implemented in the Vienna ab-initio Simulation Package (VASP).^45–47The plane-wave kinetic energy cutoff was set to 500 eV for geometric and electronic structures calculation. All the atoms were allowed to fully relax until the force on each atom was less than 0.01 eV\({\text{Å}}^{-1}\). The vacuum region was set to 20Å to avoid spurious interactions between periodic images. The Hubbard U was taken into account to describe the on-site Coulomb repulsion of Cr’s 3d electrons.

Acknowledgements

This work was supported by Guangdong Basic and Applied Basic Research Foundation (grant no. 2022A1515111001) and the National Research Foundation of Singapore (MOE Tier 2 grant WBS A-8000942-00-00). The computation part of the work was supported by National Supercomputer Center in Guangzhou.

Author contributions

M.L. and A.T.S.W. proposed and conceived this project. M.L. and J.G. performed the STM and nc-AFM experiments. M.L., Z.L. and J.X. conducted the MBE growth. M. L. and Y.Y. provided the theoretical support on the experiments. M.L., Z.C., X.X, D.Z., G.E. and A.T.S.W did the data analysis. M.L., D.Z. and A.T.S.W wrote the manuscript with the comments from all co-authors.

Competing financial interests

The authors declare no competing financial interests.

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