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Dative epitaxy of commensurate monocrystalline covalent-van der Waals moir\'e supercrystal

Realizing van der Waals (vdW) epitaxy in the 80s represents a breakthrough that circumvents the stringent lattice matching and processing compatibility requirements in conventional covalent heteroepitaxy. However, due to the weak vdW interactions, there is little control over film qualities by the substrate. Typically, discrete domains with a spread of misorientation angles are formed, limiting the applicability of vdW epitaxy. Here we report the epitaxial growth of monocrystalline, covalent Cr5Te8 2D crystals on monolayer vdW WSe2 by chemical vapor deposition, driven by interfacial dative bond formation. The lattice of Cr5Te8, with a lateral dimension of a few ten microns, is fully commensurate with that of WSe2 via 3 x 3 (Cr5Te8)-7 x 7 (WSe2) supercell matching, forming a single crystalline moire superlattice. Our work has established a conceptually distinct paradigm of thin film epitaxy termed dative epitaxy, which takes full advantage of covalent epitaxy with chemical bonding for fixing the atomic registry and crystal orientation, while circumventing its stringent lattice matching and processing compatibility requirements; conversely, it ensures the full flexibility of vdW epitaxy, while avoiding its poor orientation control. Cr5Te8 2D crystals grown by dative epitaxy exhibit square magnetic hysteresis, suggesting minimized interfacial defects that can serve as pinning sites.

However, the conventional exfoliation and stacking approach lacks scalability for practical applications.
Recently, 2D vdW heterostructures have been realized by chemical vapor deposition (CVD), such as graphene/hBN [9] , and transition metal dichalcogenide (TMD) heterostructures (e.g., WS2/MoS2, SnS2/MoS2, and NbTe2/WSe2) [10] . VdW epitaxy overcomes the constraints of lattice matching and processing compatibility requirements in conventional covalent heteroepitaxy [11] . It is particularly suitable for synthesizing 2D heterostructures owing to the atomically smooth and dangling bond-free vdW surface. However, because vdW surfaces are chemically inactive, chemical or plasma treatment may be needed to facilitate nucleation, which leads to a defective interface [12] . Moreover, limited success has been achieved for vdW epitaxy of continuous thin films of materials with 3D crystal structures on vdW substrates. Instead, discrete nanowires or domains with a spread of misorientation alignment were typically obtained [13] . This is because the weak vdW interaction and the resulting energy landscape as a function of the in-plane orientation angle may not exhibit clearly defined minima, required for high-quality epitaxy [13][14] .
In this work, we report dative epitaxy of a high-quality single-crystalline layer of 3Dbonded material on a 2D vdW template with large lattice mismatch. This unexplored regime of epitaxy exploits the bonding duality at the interface to realize epitaxial growth: the weak vdW interactions allow facile surface diffusion of precursor molecules for the growth of large area continuous layers, while the formation of dative bond (a special covalent bond where the bonding electrons derive from the same atom) fixes atomic registry and crystal orientation.
Specifically, we focus on the 2D heterostructure of Cr5Te8/WSe2, where Cr5Te8 is a non-vdW ferromagnet that can be considered as Cr atoms self-intercalated in-between the CrTe2 layers [15] (as shown schematically in Figure S1a, supporting information). First, a large-scale WSe2 monolayer millimeter in size was grown on a sapphire or Si/SiO2 substrate by CVD. Highly aligned 2D Cr5Te8 crystals with thicknesses down to a single unit cell and yet sizes of tens of microns were then achieved by dative epitaxy, with WSe2 as the vdW template. As a result, a globally commensurate, monocrystalline 3 × 3/7 × 7 Cr5Te8/WSe2 moiré supercrystal is achieved, which differs from conventional moiré superlattices with spatially varying rigid moiré patterns [16] or local commensurate domain reconstruction [17] . Three decades after the realization of vdW epitaxy [11a] , our work redefines the scope and applicability of epitaxy with unprecedented opportunities for applications.

Results and Discussion
A schematic of the growth process of 2D Cr5Te8/WSe2 moiré superlattices is shown in Figure 1a, which consists of monomer adsorption, desorption, and surface diffusion. Further details of the growth process are shown in Methods. CVD grown monolayer WSe2 with lateral dimensions of 100 -2,000 m, used as templates for the growth of Cr5Te8/WSe2 heterostructures, are shown in Figure 1b, c [18] . A typical optical microscope image of the heterostructures is shown in Figure 1d. Strikingly, all the Cr5Te8 crystals grown on a single monolayer WSe2 are self-aligned, with one of their edges oriented either parallel or at a 60 angle to one of the edges of WSe2, in contrast to the randomly oriented Cr5Te8 crystals grown directly on sapphire (Figure S2a,supporting information). This strongly suggests that the Cr5Te8 crystals grow epitaxially on WSe2. While monolayer WSe2 are randomly oriented, the orientations of the Cr5Te8 crystals align with individual WSe2 crystals, suggesting the dominant role of monolayer WSe2 in the epitaxial growth. Therefore, such vdW templates also allow the synthesis of highly oriented 2D Cr5Te8 crystals independent of substrates, as evidenced by samples grown on amorphous SiO2 substrates (Figure S2g, h, supporting information). heterostructures; e, f Highly aligned Cr5Te8 crystals with a thickness of ~ 10 nm (e) and 1.4 to 2.8 nm (the dashed lines serve as visual aids to discern the boundaries of 2D crystals) (f) on a single monolayer WSe2; g, An AFM image of an area of one unit cell thick Cr5Te8 crystals; the dashed line shows the boundary between monolayer WSe2 and sapphire substrate.
Optical microscope images of Cr5Te8/WSe2 heterostructures with relatively thick (~ 10 nm) and thin (1.4 to 2.8 nm) Cr5Te8 crystals are shown in Figure 1e and 1f, respectively. They were achieved by controlling the distance between the CrCl3 precursor and the substrate. The atomically thin Cr5Te8 crystals in Figure 1f exhibit extremely weak contrast. To help discern these crystals from the substrate, the boundaries of these crystals are highlighted by the dashed lines (the same image with enhanced contrast is shown in Figure S2b, supporting information).
A group of one unit cell thick, aligned Cr5Te8 crystals on WSe2 are further shown by the atomic force microscopy (AFM) image in Figure 1g and Figure S2d.
The a lattice constants of freestanding WSe2 and Cr5Te8 are 3.33 and 7.90 Å, respectively.
Conventional vdW epitaxy mechanism implies that with ~ 16% lattice mismatch, defined as ( 5 8 − 2 * 2 )/ 5 8 , the Cr5Te8 and WSe2 would form incommensurate moiré superlattices [19] . Indeed, a moiré pattern can be clearly seen in the atomically resolved planeview high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) image in Figure 2a. However, the strictly periodic moiré pattern suggests that the Cr5Te8 and WSe2 lattices are commensurate. The fast Fourier transform (FFT) pattern (Figure 2b) reveals three distinct sets of diffraction spots with six-fold symmetry, marked by yellow, red, and blue circles, respectively. The lattice spacing of the diffraction spots marked by yellow circles is 2.88 Å, which is consistent with the measured (100) lattice spacing of monolayer WSe2. The diffraction spots marked with red circles show identical orientation with that of the WSe2, with a lattice spacing of 3.36 Å in accordance with the Cr5Te8 (200) planes, which reduces by ~2% from the value of freestanding Cr5Te8 [20] . The inner, red-circled diffraction spots with a larger periodicity (6.72 Å) representing those ordered, self-intercalated Cr atoms is a structural fingerprint of trigonal Cr5Te8. Interestingly, the innermost diffraction spots highlighted by blue circles indicate a lattice spacing of 11.6 Å, which belongs to neither WSe2 nor Cr5Te8 alone, suggesting that it originates from the periodicity of the moiré superlattice of the Cr5Te8/WSe2 heterostructure. The individual lattices of Cr5Te8 and WSe2 are resolved by inverse FFT (iFFT) of the corresponding filtered diffraction spots ( Figure S3, supporting information), showing hexagonal lattices of both WSe2 and Cr5Te8 aligned in identical orientation, as seen in Figure 2c and 2d, respectively.
An atomic model of the Cr5Te8/WSe2 heterostructure based on iterative refinements from HAADF-STEM image analysis and first principles calculations is shown in Figure 2e, where a 3 × 3 Cr5Te8 supercell is commensurate with the 7 × 7 WSe2 supercell. To generate the larger periodicity with six-fold symmetry that reproduces the observed moiré diffraction pattern shown by the blue circles in Figure 2b, the model requires the number of interfacial Cr atoms to be reduced to 3 per supercell from 9 self-intercalated in the interior of Cr5Te8, suggesting interfacial reconstruction. In slight variations of the atomic model, either without interfacial Cr or with 9 interfacial Cr per supercell, the moiré superlattice diffraction pattern is noticeably missing since it is symmetry forbidden ( Figure S5, supporting information). This further confirms the reconstructed interface with suitable Cr occupation is a necessary condition for the observed moiré diffraction pattern and the commensurate moiré superlattice. As can be seen from Figure 2f, the simulated diffraction pattern matches well with that of Figure   First principles calculations were carried out to understand the atomic structure, charge transfer, and chemical bonding at the Cr5Te8/WSe2 interface. In bulk, the self-intercalated Cr atoms are coordinated with 6 Te atoms arranged on the corners of a triangular prism, as seen in Figure 2e. The interfacial Cr atoms have excess electrons due to the reduced coordination (losing 3 nearest Te neighbors comparing to the bulk). After forming the heterostructure, the crystal symmetry is further lowered with these interfacial Cr atoms coordinating to Se atoms on the WSe2 side, forming a Te-Cr-Se Janus interface. Figure 3a compares the site-decomposed partial DOSs of interfacial Cr (left panel) and Se (right panel) sp states for isolated Cr5Te8 and WSe2 monolayer (red curves) with those of WSe2/Cr5Te8 heterostructure (black curves). As can be seen from the change of DOS of Cr in the left panel, the majority spin state peak near the Fermi level (the red peak) is suppressed. Accompanying the suppression, a band at ~ -7 to -4 eV emerges. The emergence of this low-lying Cr band is a result of hybridization with the Se sp states, whose energy is also lowered due to the hybridization, as can be seen in the right panel.
One can understand the above results based on a level repulsion picture in Figure 3b: the two states shown in red color are the interfacial Cr and Se sp states before Cr5Te8 and WSe2 form the heterostructure. When forming the heterostructure, the Cr and Se sp states hybridize: the high-lying Cr sp state is pushed up in energy forming an anti-bonding state so its spectra weight near the Fermi level is suppressed, while the low-lying Se sp state is pushed down forming a bonding state in line with the result in the right panel in Figure 3a. As this is a hybridization, Cr also takes a significant share in the newly-formed low-lying state, as evidenced by the band at ~ -7 to -4 eV in the left panel in Figure 3a. Interestingly, after the hybridization, the antibonding state is pushed up to be above the WSe2 conduction band edge, resulting in electron transfer from the interfacial Cr to WSe2 conduction band to reduce system energy. Figure 3c shows the differential (deformation) charge density Δρ, which reveals a bondingstate charge accumulation in-between the interfacial Cr and Se atoms. Based on the Bader analysis [21] (see the Table S2), the total amount of charge accumulated is about 1 electron per supercell. The interfacial Cr-Se bond length of 2.88 Å is noticeably longer than the 2.56 Å for Cr2Se3.The interfacial binding energy of ~ 1 eV/Cr is an order of magnitude larger than a typical vdW binding but is only half of the usual Cr-Se covalent bond. These results point consistently to the formation of dative bonds at the interface [13] , which originates from Coulomb attraction between anion lone pairs (i.e., the doubly occupied non-bonding states of interfacial Se atoms) and the empty orbitals of the metal cations (i.e., the nearly empty Cr sp states upon electron transfer to the WSe2 conduction band) and is intermediate in strength between the vdW binding and covalent bond. The formation of the dative bonds weakens adjacent Cr-Te and W-Se bonds, e.g., the W-Se bond length increases from 2.545 to 2.550 Å. interfacial strain that would appear otherwise for the epitaxial growth.
In conventional vdW heterostructures with large lattice mismatch, either incommensurate superlattices with spatially varying moiré patterns [16] or local commensurate domain reconstruction were observed [17] . In conventional covalent heteroepitaxy, on the other hand, interfacial strain would lead to defects such as dislocations. For the Cr5Te8/WSe2 system, however, nature optimizes the atomic structure with just the right number of dative bonds at the interface. This allows nearly strain-free commensurate moiré superlattices over the entire 2D heterostructure with minimum density of interfacial defects. As evidenced by the HAADF-STEM images taken at different spots of a single Cr5Te8/WSe2 heterostructure ( Figure S4, supporting information), identical, perfectly commensurate moiré patterns were observed across the heterostructure, suggesting that the Cr5Te8/WSe2 is a monocrystalline moiré supercrystal. To our knowledge, such a monocrystalline moiré supercrystal has not been reported before and provides strong evidence for the proposed dative epitaxy mechanism.
The cross-section of a relatively thick (~ 7 nm) Cr5Te8 layer grown on WSe2 were imaged by HAADF-STEM to reveal atomic details of the interface (atomically thin Cr5Te8/WSe2 got oxidized during cross-sectional sample preparation). To reveal the position of Cr atoms that are much lighter than Te, integrated differential-phase contrast (iDPC) imaging technique were employed. As can be seen in Figure 3e, the cross-section of Cr5Te8 viewed along (100) direction matches well with the atomic model in Figure S1a. Atomic columns with substantially weaker contrast inside the vdW-like gap can be seen from the zoomed-in view in Figure 3f, which is attributed to the reduced number of interfacial Cr atoms, consistent with the atomic model matching diffraction and theoretical predictions. The local valence states of Cr atoms across the interface were mapped using the integrated intensity ratio of the electron energy loss spectroscopy (EELS) L3 and L2 excitation peaks (the so-called "white line ratio") [22] . As shown in Figure 3g, the Cr L3/L2 ratio increases towards the interface and becomes substantially larger than the value in the Cr5Te8 interior, indicating a lower Cr valence state near the interface (individual EELS spectrum at the bulk and interface is provided in Figure S6, supporting information). This is because after dative bond formation and electron donation, these interfacial Cr atoms still possess excessive charge due to lower coordination.
The predicted weakening of W-Se bonds in WSe2 was further investigated by Raman spectroscopy. As seen from the bottom panel in Figure 3h, a strong peak at ~ 250 cm -1 and a weak shoulder at ~ 260 cm -1 are observed for monolayer WSe2, which can be attributed to the degenerate out-of-plane A1g and in-plane E 2g 1 phonon modes of WSe2, and a second-order Raman mode due to LA phonons at the M point in the Brillouin zone labeled as 2LA(M), respectively [23] . These modes are also observed for the Cr5Te8/WSe2 moiré superlattice.
However, both peaks showed a small but measurable red shift. The average Raman peak positions measured at 9 different spots each for WSe2 and Cr5Te8/WSe2 are shown in Table 1 (all spectra and fittings are shown in Figure S7a, b and Table S1). A red shift (∆ ) of 1.3 cm -1 for the E 2g 1 /A 1g mode and 2.4 cm -1 for the 2LA(M) mode were observed, confirming the predicted W-Se bond softening due to dative bond formation.   Dative epitaxy enables nearly strain-free epitaxial growth of discrete monocrystalline 2D Cr5Te8 crystals on a single monolayer WSe2, which should lead to extremely low density of interfacial defects. The out-of-plane magnetic hysteresis of single 2D Cr5Te8 crystals were measured by reflective magnetic circular dichroism (RMCD), which is used to infer the crystallinity of Cr5Te8. Shown in Figure 4a-c are three representative 2D Cr5Te8 crystals with thicknesses of 8.4 nm (6 unit cells), 4.5 nm (3 unit cells) and 2.6 nm (2 unit cells), respectively.
The magnetic hysteresis loops for the three crystals measured at 5 K are shown in Figure 4d-f.
All three samples exhibit square hysteresis loops, with sharp transitions at the coercive fields (HC), and unity remanence at zero field. HC are 0.66, 0.36, and 0.74 T for 6, 3 and 2 unit cell thick crystals, respectively, which are noticeably smaller than the expected anisotropy field [24] , suggesting that the magnetization reversal proceeds by nucleation (e.g. at a sharp corner) followed by domain wall motion. The nearly perfect square hysteresis suggests nearly absence of domain wall pinning by defects, and thus once a magnetic domain is nucleated, the domain wall can propagate freely. On the other hand, the reported magnetic hysteresis of CrxTe1-x nanoplates on covalent substrates show skewed loops with a broad switching field distribution [25] . This might imply the presence of interfacial pinning sites. This comparison suggests that 2D Cr5Te8 crystals obtained by dative epitaxy does possess superior crystal quality and magnetic properties. For one unit cell thick Cr5Te8 crystals, however, no RMCD signal can be detected. The lack of magnetic signal may be due to surface oxidation.

Conclusion
In conclusion, Cr5Te8/WSe2 moiré supercrystal has been achieved by dative epitaxy of covalent 2D Cr5Te8 crystal on monolayer WSe2. The Cr-Se dative bond formation at the interface drives the epitaxial growth of highly aligned 2D Cr5Te8 crystals. The dative epitaxy results in perfectly commensurate, monocrystalline moiré supercrystals, which are distinctly different from conventional moiré superlattices. The high crystal quality of the 2D Cr5Te8 crystals is further confirmed by square magnetic hysteresis loops nearly free of defect pinning sites. The dative epitaxy is likely not limited to the present material systems and synthesis method, but should be applicable to a wide range of covalent materials on vdW templates, using a variety of thin film deposition techniques. This newly realized paradigm of thin film epitaxy not only is attractive for a plethora of industrial applications, but also offers opportunities to explore emergent phenomena in previously unattainable heterostructures.

Synthesis of Cr5Te8/WSe2 heterostructures:
Cr5Te8/WSe2 heterostructures were synthesized through a two-step CVD process in a two-zone tube furnace with a 2" diameter. A schematic of the experimental setup and the heating profiles of the synthesis were shown in Figure S1b, c.
In the first step, WSe2 monolayer was grown on sapphire or SiO2/Si substrate. In a typical synthesis, 200 mg Se powder was placed in the first heating zone upstream of the furnace, which was kept at 400 ℃ during the growth. 5 mg WO3 was mixed with 0.5 mg NaCl and loaded in the second heating zone downstream. Sapphire substrates were placed close to the WO3 powder, while SiO2/Si substrates were placed face down directly above the WO3 powder. The second heating zone was heated with a ramping rate of 20 ℃/min to the growth temperature of 820 ℃ and held at that temperature for 20 min before cooling down. During the growth process, the flow rate of 2% H2/N2 was kept at 80 standard cubic centimeters (sccm), and the growth was at ambient pressure. In the second step, the as-grown WSe2 were used as the template for the epitaxial growth of 2D Cr5Te8 crystals to obtain Cr5Te8/WSe2 heterostructures. 40 mg Te powder was placed in the first heating zone and ramped to 540 °C at a rate of 13 °C/min, and kept at 540 °C for 10 min. 1.2 mg CrCl3 powder was placed in the second heating zone, and heated to 600 °C at a rate of 15 °C/min. The growth time was fixed at 10 min. A gas mixture containing 90% Ar and 10% H2 with a flow rate of 100 sccm was used to carry the precursor vapor species to the substrate. Once the reaction ended, the furnace was cooled down naturally to room temperature.
The CVD growth of 2D Cr5Te8 crystals on WSe2 template is dictated by monomer adsorption, desorption, and surface diffusion, as seen in Figure 1a. The weak bonding between the monomer and the WSe2 vdW template leads to low barriers for surface diffusion [26] . The monomers aggregate to form nuclei. Once a nucleus reaches a critical size, [26a] the growth proceeds primarily by surface diffusion of monomers and their attachment to the edges of the nucleated 2D islands [26b] . For Cr5Te8, the intralayer covalent bonding is substantially stronger than the interlayer bonding due to the presence of ordered vacancies. Such bonding character results in stronger adsorption energy of the monomers at the edges than that on the top surfaces.
Together with ease of diffusion of monomers on the surface of WSe2, atomically thin Cr5Te8 crystals can be achieved [26b] .
The Cr5Te8/WSe2 heterostructures with relatively thick (~ 10 nm) and thin (1.4 to 2.8 nm) Cr5Te8 crystals were achieved by controlling the distance between the CrCl3 precursor and the substrate, being ~ 0.2 mm for the samples shown in Figure 1e and ~ 2 mm for the ones in Figure   1f. Due to the high melting point (1,150 C)  Interestingly, the 2D Cr5Te8 crystals deposit highly selectively on WSe2 only, leaving sporadic nanocrystals on sapphire, as seen in Figure 1d. We attribute the absence of 2D crystal growth on these substrates to the covalent bonding and thus large surface diffusion barrier for monomers. This leads to 3D growth of nanoparticles which eventually dewet due to surface tension.

Film transfer:
The as-grown Cr5Te8/WSe2 heterostructures were transferred onto TEM grids by dry transfer in a glovebox with a nitrogen atmosphere. The Cr5Te8/WSe2 heterostructures on SiO2/Si substrate were first covered by polymethylmethacrylate (PMMA) [27] . After baking at 80 °C for 5 min, the PMMA film with Cr5Te8/WSe2 heterostructures was peeled off from the SiO2/Si substrate and then transferred to a TEM grid in a home-built alignment stage integrated with an optical microscope, followed by 5 min baking at 80 °C. The PMMA was removed by immersing the sample in acetone for 30 min.

Cross-sectional STEM sample preparation:
The as-grown Cr5Te8/WSe2 heterostructure was exposed to a nitrogen atmosphere and subsequently covered by graphite through a routine drytransfer method in the glovebox to protect the surface from being oxidized. The cross-section STEM sample was prepared by using Focused Ion Beam (FIB) milling. It was thinned down to 70 nm thick at an accelerating voltage of 30 kV with a decreasing current from 0.79 nA to 80 pA, followed by a fine polish at an accelerating voltage of 2 kV with a small current of 21 pA to remove the amorphous layer.

HAADF-STEM characterization:
The atomically resolved HAADF-STEM images were carried out on an aberration-corrected scanning transmission electron microscope (FEI Tian Themis 60-300kV, operate at 300 kV). This TEM is equipped with a DCOR aberration corrector and a high-brightness field emission gun (X-FEG) with monochromator. The inner and outer collection angles for the STEM images (β1 and β2) were 38 and 200 mrad, respectively, with a semi convergence angle of 30 mrad.

Cross-sectional STEM imaging:
The cross-sectional HAADF-STEM reveals a clear vdW-like gap between the Cr5Te8 and WSe2 layers as shown in Figure S4h, but Cr sites are nearly invisible due to strong scattering of Te which has a much larger atomic weight. To reveal the position of Cr, we adopted integrated differential-phase contrast (iDPC) imaging, which measures the projected electrostatic potential instead of the integrated scattering signal of the atomic column.

STEM-EELS characterization:
EELS were acquired in the STEM mode and collected by setting the energy resolution to 1 eV at full width at half maximum (FWHM) of the zero-loss peak. The dispersion used is 0.5 eV/channel. EELS are acquired in the dual EELS mode to eliminate any systematic error due to the drift of the zero-loss peak.

STEM-EDS characterization: EDS was acquired in the STEM mode with a ChemiSTEM
technology (X-FEG and SuperX EDS with four windowless silicon drift detectors) operated at 300 kV.
Raman and photoluminescence spectra were measured using a confocal Renishaw inVia Raman microscope equipped with a 514 nm laser. A 50× objective lens was used to focus the excitation lasers onto the sample and collect the emitted signals.

X-ray diffraction (XRD) spectrum:
A Rigaku Ultima IV XRD system with an operational Xray tube power of 1.76 kW (40 kV, 44 mA) and Cu target source was used. The XRD measurements were performed under theta/2 theta scanning mode and continuous scanning type with a step size of 0.02.
X-ray photoelectron spectroscopy (XPS) spectrum: XPS was conducted on a PHI 5000 Versaprobe system using Al Kα X-ray radiation for excitation.

DFT-based ab-initio calculations were performed by using the Vienna ab initio Simulation
Package (VASP) package. We used the Perdew-Burke-Ernzerhof (PBE) form of the exchange correlation functional. Slab calculations were performed using supercell approach with a vacuum layer of ~15 Å (to remove interaction between periodically repeated layers). The supercell was constructed using the observed moiré superlattice. The in-plane lattice constant of Moiré superlattice was set at 23.3 Å. Plane-wave cut-off energy of 400 eV, and 4 × 4 × 1 Monkhorst-Pack k-point mesh. The atomic positions were optimized by the conjugate gradient method to have all forces less that 10 -2 eV/Å. Spin-orbit was added after the relaxation accuracy was achieved. Zero damping DFT-D3 method of Grimme models Van der Waals interactions.

RMCD measurements:
The samples for RMCD measurements were capped with 2 nm Al by sputter deposition to prevent oxidization. The RMCD is defined as where the ± are the intensities of the reflected right and left circularly polarized light. RMCD measurements were performed with the sample mounted on a custom microscope/nano positioner probe that was loaded into the variable-temperature helium insert of a 7 T magnetooptical cryostat (Oxford Instruments Spectramag). Light from a 633 nm HeNe laser was linearly polarized, and then modulated between left-and right-circular polarization at 50 kHz using a photoelastic modulator, before being focused to a 1-micron diameter spot on the sample. Light reflected from the sample was detected by an avalanche photodiode, and the normalized difference between the two polarizations was measured using a lock-in amplifier.
XMCD measurements: XMCD measurements were performed at beamline 11.0.1 of the Advanced Light Source at the Lawrence Berkeley National Laboratory.

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author. wrote the manuscript. All authors discussed the results and commented on the manuscript.

Conflict of Interest
The authors declare no conflict of interest.

Data Availability Statement
The data that support the plots within this paper and other finding of this study are available from the corresponding author on reasonable request.

X-ray absorption spectra (XAS) and XMCD spectra
The XAS and XMCD spectra of a Cr5Te8 2D crystal with a thickness of 3 unit cell was measured at the Cr L2,3 edges (560.0 -607.0 eV) at beamline 11.0.1 of the Advanced Light Source at the Lawrence Berkeley National Laboratory. Measurements were conducted at 105 K without an external magnetic field. The x-ray incident angle was 30° with respect to the film plane. Representative XAS with the corresponding XMCD spectra are shown in Figure S9a. To assist the magnetic moment calculations, the integrations of the XMCD and XAS spectra are shown in Figure S9b, c. The spin and orbital magnetic moments per atom, in the units of B/atom, can be determined from the XAS and XMCD spectra in Figure S9 by applying the XMCD sum rules [1] , specifically, for 3d transition metal element Cr 3+ according to the following equations: , 3+ = − ℎ, 3+ 3 − 2 × − < > , 3+ = − ℎ, 3+ 2 3 (1 ) where p, q, r are corresponding integrals read from Figure S9b, c for Cr 3+ ; the valence hole numbers of Cr 3+ (nh, Cr3+) is 7; SC is the spin correction factor estimated for Cr in Cr5Te8 [2] ; <Tz> is the expectation value of the magnetic dipole operator, which is negligible due to the small orbital moment of the 3d element [2][3] . From Figure S9, the calculated Mspin, Cr3+ and Morb, Cr3+ are 1.2 ± 0.3 μB and -0.02 ± 0.01 μB per Cr atom. Therefore, the sum gives MCr3+ = 1.2 ± 0.3 μB/atom at 105 K along the x-ray propagation direction, which is 30° away from the sample plane. The comparably large error bar mainly comes from the relatively large uncertainty in absorption background subtraction. Notice that the Cr5Te8 2D crystal possesses a perpendicular anisotropy based on RMCD measurements. Therefore, the estimated Mspin and Morb should be multiplied by a factor of 2 to project Mspin and Morb back to the out-of-plane direction, resulting Mspin = 2.4 ± 0.6 B/Cr 3+ and Morb = -0.04 ± 0.02 B/Cr 3+ . The estimated spin moment at 0 K is ~ 3 B, which is close to the average Cr 3+ spin moment of 3.03 B obtained from DFT calculations. In the fifth column, the first number is for the Se atoms at the interface, and the second number is for the atoms at the free surface. The Bader charge analysis shows that there is a charge transfer from Cr5Te8 towards WSe2. There is ~ 1e transferred across the interface. The largest charge transfer occurs for the intercalating Cr atom. It donates electrons to form a dative bond with Se in WSe2. Some charge is also transferred to W, although significantly smaller. Due to the asymmetry of the charge between Se sites at the interface and free surface, there is a polarity in charge distribution across the WSe2 monolayer.