Synthesis of Antimonene on Germanium - Nano Letters (ACS

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Synthesis of Antimonene on Germanium Matthieu Fortin-Deschênes, Olga Waller, Tevfik Onur Mentes, Andrea Locatelli, Samik Mukherjee, Francesca Genuzio, Pierre L. Levesque, Amanda Hebert, Richard Martel, and Oussama Moutanabbir Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.7b02111 • Publication Date (Web): 05 Jul 2017 Downloaded from http://pubs.acs.org on July 6, 2017

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Synthesis of Antimonene on Germanium §

§

Matthieu Fortin-Deschênes†, Olga Waller†, Tevfik O. Menteş , Andrea Locatelli , Samik § Mukherjee†, Francesca Genuzio , Pierre Levesque¥, Amanda Hébert†, Richard Martel¥, Oussama Moutanabbir*,† †

§

¥

Department of Engineering Physics, École Polytechnique de Montréal, C. P. 6079, Succ. Centre-Ville, Montréal, Québec H3C 3A7, Canada Elettra - Sincrotrone Trieste S.C.p.A., S.S. 14 - km 163, 5 in AREA Science Park, 34149 Basovizza, Trieste, Italy

Département de Chimie, Université de Montréal, 2900 boulevard Edouard Montpetit, Montréal, Québec, H3T 1J4 Canada

* Corresponding Author. Email: [email protected]

The lack of large area synthesis processes on substrates compatible with industry requirements has been one of the major hurdles facing the integration of 2D materials in mainstream technologies. This is particularly the case for the recently discovered monoelemental group V 2D materials which can only be produced by exfoliation or growth on exotic substrates. Herein, to overcome this limitation, we demonstrate a scalable method to synthesize antimonene on germanium substrates using solid-source molecular beam epitaxy. This emerging 2D material has been attracting a great deal of attention due to its high environmental stability and its outstanding optical and electronic properties. in situ low energy electron microscopy allowed the real time investigation and optimization of the 2D growth. Theoretical calculations combined with atomic-scale microscopic and spectroscopic measurements demonstrated that the grown antimonene sheets are of high crystalline quality, interact weakly with germanium, exhibit semi-metallic characteristics, and remain stable under ambient conditions. This achievement paves the way for the integration of antimonene in innovative nanoscale and quantum technologies compatible with the current semiconductor manufacturing. Keywords: 2D materials; Antimony; Germanium; Epitaxial growth; in situ electron microscopy; Density Functional Theory. 1 ACS Paragon Plus Environment

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The isolation of graphene in 20041 gave birth to the field of two-dimensional (2D) materials and devices. The unique properties exhibited by this new class of materials make them appealing for both fundamental studies and technological applications2-4. Nowadays, established 2D materials include monoelemental group IV materials and their derivatives, binary alloys of group III-V elements, metal chalcogenides, and layered oxides, in addition to several other newly discovered materials5, 6. The latter comprise the recently reported group V 2D materials7-16. The canonical group V 2D material has been black phosphorus (bP), which displays a thickness-dependent band-gap varying between 0.3 eV for bulk to 2.05 eV for monolayer in addition to its attractive anisotropic transport properties7. However, the use of bP in electronic and optoelectronic devices can be hindered by its rapid degradation under ambient conditions8 and the lack of scalable synthesis methods9. Other theoretically predicted group V 2D materials include several allotropes and compounds10 of nitrogen11, phosphorus12, arsenic13, antimony (Sb)14,15 and bismuth16. Amongst them, we find β-2D-Sb (or simply 2D-Sb), with a buckled honeycomb structure shared with silicene and germanene17. Bulk Sb is semi-metallic and predicted to become a topological insulator at a thickness under 22 layers, to exhibit a quantum spin Hall phase below 8 layers18, and finally to be a semiconductor with a 2.28 eV indirect gap and high electron and hole mobilities at the monolayer thickness14, 19. 2D-Sb oxides are also predicted to be 2D topological insulators with sizable global bandgaps20. Moreover, multiscale simulations of metal-oxidesemiconductor field-effect transistors with ultra-short (≤7 nm) 2D-Sb channels predict high device performances, such as a 60mV/decade subthreshold swing, which is long-sought-for in the sub-10 nm technology nodes of low-power electronics critical for the next generation of ultra-large scale integration circuits19. Recently, 2D-Sb has been produced by exfoliation21,22 and 2 ACS Paragon Plus Environment

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by epitaxial growth on van der Waals (vdW) layered materials including mica23 and PdTe224. These studies pointed out to the outstanding environmental stability of 2D-Sb. Nonetheless, it remains crucial to develop scalable synthesis methods relevant for the integration of 2D-Sb in mainstream electronic technologies. With this perspective, we demonstrate that 2D-Sb can be synthesized on germanium (Ge) (111) surfaces, which are nowadays ubiquitous in complementary metal-oxide-semiconductor (CMOS) processing. The growth of 2D-Sb was achieved using solid source molecular beam epitaxy (MBE). The growth kinetics was studied in real time by low-energy electron microscopy (LEEM) and diffraction (LEED) leading to the identification of the optimal parameters for the 2D growth. The stability of epitaxial 2D-Sb and the influence of substrate-layer interactions on its electronic properties were elucidated by abinitio calculations for different interfacial atomic configurations. Additionally, cross-sectional scanning transmission electron-microscopy (STEM) observations were combined with DFT calculations to investigate and discuss the nature of the Ge-Sb interface. Finally, the atomic structure, crystalline quality, and environmental stability of 2D-Sb were determined by in situ scanning tunnelling microscopy (STM), Raman spectroscopy, and X-ray photoemission electron microscopy (XPEEM). 2D-Sb was grown in a LEEM microscope under ultra-high vacuum (UHV). Sb crystals at 99.99999% purity is evaporated using a Knudsen cell at evaporation rate varying between 2-700 Å/minute. The substrate temperature (T) is varied between room temperature (RT) and 330°C, as measured with an infrared pyrometer. Ge(111) surfaces were prepared by two different methods. In both cases, the substrates are cleaved from an undoped wafer then sonicated in acetone for 5 minutes, rinsed with isopropanol and dried with a nitrogen flow. The substrates are then introduced in the UHV system and annealed between 600-700°C for at least one hour, then 3 ACS Paragon Plus Environment

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flashed above 800 °C for a few seconds. The second method contains an additional Ar+ sputtering step at 2kV before annealing. Sputtering and annealing is repeated until a sharp c(2×8) LEED pattern is observed. Without the sputtering step, only a diffuse c(2×8) pattern is visible (Fig. S1). Both sample preparation methods allowed for 2D-Sb growth. LEEM energy is kept between 0-5eV during growth and no evidence of an influence of the electron beam on growth was detected by comparing irradiated and non-irradiated regions post-growth. After growth, samples are transferred under UHV to a STM equipped with a polycrystalline tungsten tip at RT. WsXM was used for STM data analysis25 and FFT filtering of noise at frequencies higher than atomic periodicity was done to improve image quality in some cases. ex situ LEEM and XPEEM measurements were carried out at the Nanospectroscopy beamline26 at the Elettra Synchrotron laboratory using an incident x-ray beam at 16° incidence angle and photon flux between 109-1011 ph/sec/µm2. The photoelectron energy resolution is better than 300 meV and the XPEEM spatial resolution is better than 30nm. DFT calculations were carried out with Quantum Espresso27 using ultrasoft pseudopotentials28. The PBE functional29 was used for electronic structure calculations and the optB86b-vdW30-32 functional was used for a more precise estimation of the van der Waals interactions between 2D-Sb and Ge. Slab geometries consisting of 6 Ge bilayers (12 atoms) with the 2 bottom bilayers fixed at bulk position and an hexagonal unit cell were used to simulate the Ge(111) surface. Single layer, bilayer and trilayer 2D-Sb were added on top of the Ge(111) slab with at least 20 Å vacuum separating periodic images. For Ge, the calculated lattice constant is 5.75 Å (PBE) and 5.69 Å (optB86b-vdW), in good agreement with previous calculations30. For freestanding single layer 2D-Sb, the relaxed lattice constant is 4.05(PBE) and 4.00 (optB86bvdW), yielding lattice mismatches smaller than 0.6% on the Ge(111) surface. A 680 eV cut-off 4 ACS Paragon Plus Environment

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was used for plane-wave expansion along with (12 × 12 × 1) and (24 × 24 × 1) Г centered Monkhorst-Pack k-point meshes for structural relaxation and electronic properties calculations respectively. The criteria for electronic convergence and structural relaxation convergence were set to 1µeV and 0.01eV/ Å, respectively. After a proper Ge surface conditioning (Fig. S1), parameters favoring the 2D growth were determined by systematically varying the Sb deposition rate and Ge substrate temperature while monitoring in situ the growth morphology (Fig. 1a). At T200 °C), crystalline structures corresponding to β-2D-Sb are observed (Fig. 1a). Indeed, the corresponding LEED patterns show a 6-fold symmetry with a 4.28 ± 0.02 Å lattice parameter (Fig. 1b), slightly smaller than the 4.31 Å bulk Sb lattice constant33, but consistent with DFT predictions for few-layer antimonene22. It is noteworthy that the growth temperature influences the morphology of 2D-Sb. In fact, at 200 °C, islands are dendritic due to slow diffusion, whereas at higher T they exhibit a more regular shape with welldefined edges (Fig. 1(c-d)). The islands tend to have zigzag oriented edges, which are preferentially aligned along the Ge〈101〉 directions (Fig. 1d). Although the increase of substrate temperature seems at the first glance to improve the growth quality, it brings a major hurdle related to the sticking coefficient of the evaporated species (mainly Sb4), which decreases by several orders of magnitude above 300 °C. This limits and even prevents nucleation and growth. Besides the effects of substrate temperature, we also investigated the influence of deposition rates. For proper comparison, we compare rates at fixed T (T=270°C). We found that low rates (below 50 Å/minute) yield larger islands (up to a few µm in lateral dimension) with lower area density (Fig. 1(a,d)). The deposition rate also influences nucleation and morphology of the 2D 5 ACS Paragon Plus Environment

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islands. At deposition rates below 50 Å/minute range, 3D pyramidal nuclei first form and then 2D growth occurs around them, leading to clover-shaped islands with 3-fold symmetry (Fig. 1c). Interestingly, higher rates (e.g., 200 Å/minute) were found to facilitate the nucleation of flat, hexagonally shaped 2D islands, together with facetted 3D islands (Fig. S3&S4). Based on these observations and in order to achieve a better control of the growth morphology and quality of 2D-Sb, we established a two-step growth process with nucleation at high rate (200 Å/minute) and growth at much lower rates (Fig. S4). In this way, only a limited number of 2D and 3D islands nucleate in the first step. In the second step, 2D islands grow laterally at higher speed than 3D islands, without additional nucleation, therefore increasing the coverage of 2D islands compared to 3D islands. Finally, we examined the post-growth thermal stability of 2D-Sb and found that as-grown sheets are thermally stable up to T~400 °C. Above this temperature, edge sublimation begins (Fig. 1(e,f)) in a similar fashion to earlier observations on exfoliated bP34, 35. Different post-growth studies described below were carried out on these MBE grown 2D-Sb layers. The difference in 2D-Sb and Ge lattice constants (LEED, Fig. 1b) indicates the absence of coherency between the two lattices, suggesting a weak interaction at the interface. LEED shows two main 2D-Sb orientations with 3-fold symmetry (Fig. S5) and aligned with the Ge lattice (Fig. 1b). The two orientations are rotated 180° from each other, thus explaining the LEED 6-fold symmetry, as seen in Fig. 2a. In this figure, the two main orientations (red and green) form few µm grains which cover most of the surface. The rest of the layer (dark areas) is made of 2D-Sb at various rotational angles with respect to the Ge [111] direction. The differently oriented grains are expected to originate from two types of interfacial atomic configurations or stackings. Detailed DFT calculations were carried out to determine the energetics and the nature 6 ACS Paragon Plus Environment

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of the bonding for these Sb-Ge stackings. We model each configuration with an ABC stack of six Ge bilayers covered by 2D-Sb (Fig. 2(b-d)). 2D-Sb and Ge bilayers share the same buckled honeycomb structure, which leads us to consider ABC-X and ABC-X’ stacks for the 2 orientations. X=A, B or C and X’ is analog X, but has an inverted buckling (top atom becomes bottom atom and vice versa). For multilayer 2D-Sb, we use the stacking of bulk Sb (BAC or A’B’C’), yielding stacks in the form ABC-AC and ABC-ACB for bi- and tri-layers for example. For simplicity, the unit cell used for calculations assumes identical 2D-Sb and Ge lattice constants. This is different from the experimental data in few-layer 2D-Sb. Nonetheless, the variety of studied configurations allows to understand the nature of the bonding for a noncoherent interface, where the type of stacking varies from place to place. Our calculations show that the nature of the interaction is sensitive to the stacking, as indicated by the charge density of epitaxial 2D-Sb minus the superposition isolated atoms charge densities (Fig. 2(e, f)). For ABC-B’C’A’ (Fig. 2e), the interfacial Sb-Ge bond (black arrow) is of similar magnitude to the weak interlayer Sb bonds (blue arrow), much smaller than the intralayer Sb-Sb bonds (red arrow) (Fig. S6 for other configurations). However, for ABC-C’A’B’ (Fig. 2f), the magnitude of intralayer Sb-Sb and Sb-Ge bonds is similar. Strong Sb-Ge bonds are correlated with short Sb-Ge interatomic distances, but not necessarily with more stable structures (Table S1). The interaction energy between Ge and 2D-Sb, calculated with the optB86b-vdW functional30-32 (Table S1), is relatively small and varies between 303-457 meV/atom, as defined

by  = − (  −  −  ) , with  and  are respectively the energies per unit cell of isolated Ge(111) and 2D-Sb, and   is the energy of epitaxial 2D-Sb. These energies are higher than the typical 50-100 meV/atom for epitaxial graphene36, 37, but lower than the ~700 meV/atom for other epitaxial 2D materials such as silicene on Ag(111)38. 7 ACS Paragon Plus Environment

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This weak interfacial interaction, crucial to preserve 2D materials’ properties, is further elucidated in Fig. 2(g-j). An STM image of multilayer 2D-Sb with 4-7 layers terraces (Fig. 2g) displays Ge step edges underneath the 4th (orange) and 5th (yellow) terraces, meaning that growth occurs across Ge steps, with bonds only within Sb layers. An intensity modulation attributed to a moiré pattern, similar to those reported for epitaxial graphene39, is observed. The moiré corrugation decreases with increasing layer thickness and the 6.3nm periodicity, together with the alignment of the two lattices, yields a 4.27 Å 2D-Sb lattice constant. The presence of a moiré pattern indicates the absence of coherent interface with the substrate, even for few-layer 2D-Sb. STEM diffraction and real-space images (Fig. 2(h-j)) confirm that the Ge-Sb interface is not coherent and that the 2D-Sb lattice parameter remains constant across the layer’s thickness. For example, Fig. 2h shows an ABC-X’ interface with no local coherence between Sb and Ge. Moreover, the 3.9±0.1 Å adsorption height agrees with the calculated 4.09 Å height for the ABC-A’ and ABC-B’ vdW-like stackings and is in opposition to the 4.55 Å adsorption height for the covalently bonded ABC-C’ stacking (Table S1). These observations are consistent with LEED and STM measurements and suggest the absence of Sb-Ge covalent bonds. STM measurements provide deeper insight into the antimonene’s atomic structure. Fig. 3(a-c) show single layer crystals, as determined from the step height. The layers in Fig. 3(a, b) are synthesized by Sb4 deposition on Ge(111)-c(2×8) at room temperature, followed by annealing at 300 °C. Chemisorbed Sb chains appear on the surface surrounding the islands. According to the structural model by Kendelewicz et al., the chains have a 2.58 Å height with respect to Ge40. A 2.66 Å step height is measured with respect to the chains (Fig. 3b). This indicates that single layer islands stand on bare Ge and are 5.24 Å thick, in good agreement with DFT calculations (Table S1). Furthermore, the buckled honeycomb structure of β-2D-Sb is 8 ACS Paragon Plus Environment

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clearly resolved (Fig. 3c, d). The high crystalline quality and homogeneity of the layers can be seen in Fig. 3d. Expectedly, lattice constants are thickness-dependent, as predicted by calculations for free-standing films22. We obtain a 4.07 Å lattice parameter for single layers and 4.25 Å for a 2.16nm-thick island (~5 layers). Scanning tunnelling spectroscopy (STS) on asgrown few-layer 2D-Sb (Fig. 4a) shows a semi-metallic behavior, which is consistent with our DFT results for single and few-layer epitaxial 2D-Sb (Fig. 4(b-d)). In fact, single layer 2D-Sb seems to be metallic or semi-metallic for all studied stacks (Fig. S7). For example, the ABC-B density of states projected onto Sb atomic orbitals (PDOS) becomes non-zero at the Fermi level (Fig. 4b), as compared to the DOS of freestanding 2D-Sb which is semi-conducting (Fig. 4d). Bilayers and trilayers are also metallic or semi-metallic.

Additonal ex situ studies were carried out using Raman spectroscopy and XPEEM (Fig. 4(e-g)) to investigate further the nature, crystalline quality, and environnemental stability of the grown antimonene. Previous studies report a loss of Raman signal with decreasing 2D-Sb thickness21-23. Here, we find only a small decrease in Raman intensity for layers of thicknesses down to 4 nm (Fig. 4g), even after a few days of air exposure. Furthermore, the Eg peak blue shifts from 115.6 cm-1 to 124.0 cm-1, as thickness decreases from 30nm (~80 layers) to 4nm (~10 layers), in agreement with previous measurements and calculations21-23. Interestingly, the Eg /A1g intensity ratio increases from 0.19 to 0.40 with thickness decreasing from 30nm to 4nm. Interestingly, no Sb oxidation is detected by XPEEM on single 2D-Sb island after one week of air exposure (Fig. 4(e, f)). In fact, the local XPS spectra (Fig. 4f) taken from nanscale regions on an individual island (A) in Fig. 4b rules out any possible oxidation of the grown antimonene. These spectra are composed of the contribution of the undelying bulk Ge (3d3/2, 3d5/2), Ge oxides, and Sb peaks (4d3/2 (33.0 eV), 4d5/2 (31.8 eV)). Note that X-ray induced Sb oxide 9 ACS Paragon Plus Environment

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reduction, which can occur under long irradiation times (Fig. S8), doesn’t affect our observations. This clearly indicates that as-grown 2D-Sb on Ge is highly stable under ambient conditions.

In summary, we have successfully demonstrated the epitaxial growth of both single layer and few-layer high-quality antimonene on Ge(111) and have optimized the 2D growth parameters using in situ electron microscopy. Different interfacial configurations of 2D-Sb on Ge(111) have been theoretically investigated and correlated with LEED, STM, and TEM measurements. The experimental results indicated a weak interaction between Sb and Ge, in agreement with DFT calculations. The as-grown 2D-Sb layers were found to be semi-metallic for all considered interfacial configurations and remain perfectly stable under ambient conditions. We believe that these results will stimulate further theoretical and experimental investigations of 2D-Sb materials and devices. Moreover, the achieved growth on a group IV substrate is a very promising approach for the development of scalable antimonene devices laying the groundwork for the integration of this novel 2D material in technologies compatible with the current infrastructure of CMOS processing.

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Acknowledgements The authors thank P. Desjardins and R. Jacobberger for fruitful discussions. O.M. acknowledges support from NSERC Canada (Discovery: RGPIN421837-12 and SPG Grants: STPGP 494031), Canada Research Chair (Award No. 950-228250), Fondation de l’École Polytechnique de Montréal, and MRIF Québec. Computations were made on the supercomputer Briarée from Université de Montréal, managed by Calcul Québec and Compute Canada. The operation of this supercomputer is funded by the Canada Foundation for Innovation (CFI), ministère de l’Économie, de la Science et de l’Innovation du Québec (MESI) and the Fonds de recherche du Québec Nature et technologies (FRQ-NT). Supporting Information Available Supporting information include: effect of sputtering on surface preparation; growth of amorphous film at low temperature; AFM and STM characterization of 2D and 3D islands; 2 step growth; micro-LEED on single 2D-Sb island; single layer 2D-Sb islands; DFT calculation for all interfacial configurations; x-ray induced Sb oxide reduction; and LEEM movie of 2D-Sb growth.

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Horcas, I.; Fernández, R.; Gomez-Rodriguez, J.; Colchero, J.; Gómez-Herrero, J.; Baro, A. Review of Scientific Instruments 2007, 78, (1), 013705. Menteş, T. O.; Zamborlini, G.; Sala, A.; Locatelli, A. Beilstein Journal of Nanotechnology 2014, 5, 1873–1886. Giannozzi, P.; Baroni, S.; Bonini, N.; Calandra, M.; Car, R.; Cavazzoni, C.; Ceresoli, D.; Chiarotti, G. L.; Cococcioni, M.; Dabo, I. Journal of physics: Condensed matter 2009, 21, (39), 395502. Vanderbilt, D. Physical Review B 1990, 41, (11), 7892. Perdew, J. P.; Burke, K.; Ernzerhof, M. Physical review letters 1996, 77, (18), 3865. Klimeš, J.; Bowler, D. R.; Michaelides, A. Physical Review B 2011, 83, (19), 195131. Thonhauser, T.; Cooper, V. R.; Li, S.; Puzder, A.; Hyldgaard, P.; Langreth, D. C. Physical Review B 2007, 76, (12), 125112. Langreth, D.; Lundqvist, B. I.; Chakarova-Käck, S. D.; Cooper, V.; Dion, M.; Hyldgaard, P.; Kelkkanen, A.; Kleis, J.; Kong, L.; Li, S. Journal of Physics: Condensed Matter 2009, 21, (8), 084203. Schiferl, D.; Barrett, C. Journal of Applied Crystallography 1969, 2, (1), 30-36. Fortin-Deschenes, M.; Levesque, P. L.; Martel, R.; Moutanabbir, O. The journal of physical chemistry letters 2016, 7, (9), 1667-1674. Liu, X.; Wood, J. D.; Chen, K.-S.; Cho, E.; Hersam, M. C. The journal of physical chemistry letters 2015, 6, (5), 773-778. Lončarić, I.; Despoja, V. Physical Review B 2014, 90, (7), 075414. Sforzini, J.; Nemec, L.; Denig, T.; Stadtmüller, B.; Lee, T.-L.; Kumpf, C.; Soubatch, S.; Starke, U.; Rinke, P.; Blum, V. Physical Review Letters 2015, 114, (10), 106804. Guo, Z.-X.; Furuya, S.; Iwata, J.-I.; Oshiyama, A. Physical Review B 2013, 87, 235435. Tetlow, H.; de Boer, J. P.; Ford, I.; Vvedensky, D.; Coraux, J.; Kantorovich, L. Physics Reports 2014, 542, (3), 195-295. Kendelewicz, T.; Woicik, J.; Miyano, K.; Yoshikawa, S.; Pianetta, P.; Spicer, W. Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 1994, 12, (4), 18431847.

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Figure 1: in situ LEEM/LEED observations of 2D-Sb growth. (a) Influence of temperature and deposition rate on 2D-Sb growth. In the low flux regime, the deposition rate (F) for each substrate temperature is: F(T=60°C) = 2 Å/minute, F(T=200°C) = 2 Å/minute, F(T=270°C) = 50 Å/minute. In the high flux regime, the deposition rates are: F(T=200°C) = 10 Å/minute, F(T=270°C) = 200 Å/minute. Note that the sticking of Sb species drops for T>200 °C. Images are 2×2 µm2 in size; (b) A typical post-growth LEED pattern (16 eV) showing the Sb (red) and Ge (blue) spots; (c, d) 2D-Sb growth snapshots at (c) T=280°C, F =50 Å/minute and (d) T=325°C, F=200 Å/minute; (e, f) First signs of 2D-Sb thermal decomposition.

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Figure 2: Sb-Ge interface. (a) Dark-field LEEM (15 eV) composite image of the two main 2D-Sb orientations. Inset: the corresponding LEED (30 eV) with the main orientations circled (green and red); (b-d) Structural model used in DFT calculations: (b, c) Buckled honeycomb structure shared by 2D-Sb

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and Ge. Darker atoms are lower than lighter atoms; (d) example of Sb-Ge stack (ABC-A’); (e, f) Charge density minus superposition of atomic charge densities at the Ge-Sb interface on the Ge(112) plane for: (e) ABC-B’C’A’ and (f) ABC-C’A’B’. Color scale is in e-×a0-3; (g) STM (Sample bias Vt=1.5V, tunnelling current It=0.20nA) of multilayer 2D-Sb island. Inset: moiré pattern (10 nm scale bar); (h-j) Cross-sectional TEM of 2D-Sb on Ge(111); (h) Calculated relaxed ABC-B’C’A’ geometry superposed to atomic resolution STEM of the Ge-Sb interface; (i) Low magnification TEM of 2D-Sb layer; (j) Epitaxial 2D-Sb diffraction. Inset: distinct Sb and Ge diffraction spots.

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Figure 3: STM of 2D-Sb. (a) Single layer 2D-Sb (Vt=2V, It=0.79nA). Insets: STM image showing the structure of single-layer 2D-Sb island (bottom left); Ge (2×1)-Sb chains (top right); (b) Top: model of 2D-Sb on bare Ge(111) surrounded by Ge (2×1)-Sb chains. Bottom: STM height profile measured on blue line in Fig. 3a; (c) Buckled honeycomb structure of monolayer 2D-Sb (Vt=2V, It=0.14nA) (see Fig. S7b for details); (c) Atomic resolution STM image of ~5 layers 2D-Sb (Vt=1.3V, It=2.2nA). Inset: A close-up image.

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Figure 4: Electronic structure and environmental stability. (a) (dI/dV)/(I/V) STS of ~ 5-layer 2D-Sb (100 °C). Feedback loop open at Vt=312mV, It=0.137nA. Inset: Corresponding STS I-V curve. Xaxis: Vt=[-1V, 1V], y-axis: It=[-1.5nA, 1.5nA]; (b) DOS of freestanding 2D-Sb and PDOS (5s, 5p) of epitaxial 2D-Sb (ABC-B); (c) Electronic band structure of epitaxial 2D-Sb (ABC-C); (d) Electronic band structure of freestanding 2D-Sb. (e-g) Measurements on samples exposed to air for a few days: (e) XPEEM map (photon energy 100eV) monitoring Sb 4d photoelectrons. (f) Local XPS spectra (photon energy 200eV) on island (A). Inset: XPS spectrum of the surface region indicated by (B). Red curve: raw data, brown curve: fit. (g) Raman spectra (633 nm excitation wavelength) for samples of average 2D-Sb layer height of 4nm, 8nm and 30nm.

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