Deposition of topological insulator Sb₂Te₃ films by an MOCVD process

Abstract

Layered Sb₂Te₃ films were grown by a MOCVD process on Al₂O₃(0001) substrates at 400 °C by use of i-Pr₃Sb and Et₂Te₂ and characterized by SEM, AFM, XRD, EDX and Auger spectroscopy. The electrical sheet resistivity was measured in the range of 4 to 400 K, showing a monotonic increase with increasing temperature. The valence band structure probed by angle-resolved photoemission shows the detailed dispersions of the bulk valence band and the topological surface state of a quality no less than for optimized bulk single crystals. The surface state dispersion gives a Dirac point roughly 30 meV above the Fermi level leading to hole doping and the presence of bulk valence states at the Fermi energy.

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Introduction

Antimony telluride (Sb₂Te₃) belongs to the class of V₂VI₃ layered materials (V = Sb, Bi; VI = S, Se, Te) with a tetradymite structure. Sb₂Te₃ has been investigated for quite a long time since it is a narrow band-gap (E_gap = 0.26 eV) semiconductor with good thermoelectric characteristics near room temperature.^1–3 More recently, the interest in Sb₂Te₃ increased due to its capability to serve as a topological insulator^4–6 with promising technical applications in spintronics and topological quantum computation.^7,8 Topological insulators are bulk insulators with metallic surfaces. They are characterized by a bulk band inversion and strong spin–orbit interaction which lead to surface states that bridge the bulk band gap. The surface states display a linear dispersion in the band structure and are protected by time-reversal symmetry.^7,8 The tetradymite binary compounds such as Bi₂Te₃, Bi₂Se₃, and Sb₂Te₃ are prototypical topological insulators while Sb₂Se₃ lacks a sufficiently strong spin–orbit interaction to develop the topological surface state, according to density functional theory.⁹ Different from the theoretical predictions,⁶ Bi₂Te₃ and Bi₂Se₃ are intrinsically electron doped with the Fermi energy (E_F) strongly displaced (≈0.3–0.4 eV) from the theoretical value so that it lies in the bulk conduction band and transport is expected to be bulk dominated. This strong electron doping does not occur for Sb₂Te₃,^10,11 rendering this material very promising for reaching the topological transport regime.

Sb₂Te₃ thin films were prepared in the past by microwave heating,^12,13 vapour phase transport,^14–21 vapour–liquid–solid growth,^22,23 atomic layer epitaxy,²⁴ electrochemical deposition,^25–31 DC and RF sputtering,^32,33 and molecular beam epitaxy (MBE).^34,35 Unfortunately, these processes very often only produced polycrystalline or highly faceted Sb₂Te₃ films, which additionally often suffered from high antisite concentrations, which typically result in a strong shift of the Fermi energy.

Metal organic chemical vapour deposition (MOCVD) is a promising technical alternative, which has been shown to produce crystalline Sb₂Te₃ films.^36–44 In MOCVD processes, trialkylstibines SbR₃ (R = Me, Et) and dialkyltellanes TeR₂′ (R = Et, i-Pr, t-Bu) are typically used as molecular precursors.⁴⁵ Unfortunately, these precursors often require high pyrolysis temperatures, forcing the incorporation of impurities such as carbon into the material film. Therefore, there is a strong demand for alternative metal organic precursors. Single-source precursors are promising candidates for low-temperature deposition processes.⁴⁶ We recently demonstrated that (Et₂Sb)₂Te can be used as an MOCVD precursor for the synthesis of highly stoichiometric Sb₂Te₃ films with very low antisite defect concentrations at temperatures as low as 200 °C,⁴⁷ while O'Brien et al. and Chivers et al. deposited Sb₂Te₃ films in the temperature range from 375 to 475 °C using Sb[(TePi-Pr₂)₂N]₃ in an aerosol-assisted chemical vapour deposition (AACVD) process.⁴⁸ Unfortunately, the Sb₂Te₃ films obtained with both precursors were rather polycrystalline, consisting of hexagonal Sb₂Te₃ nanoplates. In the case of the (Et₂Sb)₂Te precursor, the low substrate temperature (200 °C) limited the diffusion mobility of the deposited Sb and Te atoms, hence avoiding the formation of a flat film. In contrast, higher substrate temperatures yielded flat Sb₂Te₃ films, but these were contaminated by elemental antimony.

We therefore became interested in suitable Sb- and Te-precursors for the MOCVD deposition at temperatures between 350 and 400 °C and report herein on the MOCVD deposition of Sb₂Te₃ films using i-Pr₃Sb as the Sb-source and Et₂Te₂ as the Te-source. Both precursors are stable liquids at ambient temperature with a good volatility. i-Pr₃Sb has been used in the past for the CVD-deposition of GeSb₂Te₄ phase change materials,^49–51 whereas Et₂Te₂ has not been used as a CVD precursor, to date. To the best of our knowledge, dimethyl ditellane Me₂Te₂ is the only ditellane that has been used as a MOCVD-precursor, to date,^52,53 but it suffers from a relatively low vapour pressure. We carefully investigated the distinctive role of different substrates and substrate temperatures on the growth of the antimony telluride films, which were analyzed by energy dispersive X-ray analysis (EDX) and Auger spectroscopy, powder X-ray diffraction (XRD), scanning electron microscopy (SEM) and atomic force microscopy (AFM). Moreover, the electrical resistance of these films was determined by a Physical Property Measurement System (PPMS) in the range of 4 to 400 K and angle-resolved photoemission spectroscopy (ARPES) experiments revealed that the surface electronic structure presents a slightly hole-doped Dirac cone with photon-energy dependent photoemission intensity and a more surface localized surface state at higher binding energy. The measured doping level of the Dirac cone indicates that bulk valence states are present at the Fermi energy where they dominate the transport properties.

Experimental section

Materials

i-Pr₃Sb and Et₂Te₂ were prepared by literature methods.^54,55

MOCVD deposition

MOCVD studies were performed in a cold-wall high-vacuum MOCVD reactor.⁴⁷ Sb₂Te₃ films were deposited on either Si(100) or Al₂O₃(0001) substrates at a working pressure of 10 mbar within 15 minutes. Si(100) substrates were previously degreased with acetone, treated with hydrofluoric acid and heated to 500 °C at 10⁻³ mbar for 1 hour in the reactor. Al₂O₃(0001) substrates were degreased with acetone and cleaned with a 3 : 1 mixture of sulphuric acid and phosphoric acid before use. The precursors were loaded into bubblers attached to the MOCVD-reactor under inert conditions (Ar) and the bubbler temperature was set to −5 °C (i-Pr₃Sb) and 20 °C (Et₂Te₂) throughout the experiment. Argon was used as the carrier gas (40 sccm Et₂Te₂, 5 sccm i-Pr₃Sb) and the precursor flow was controlled by use of a mass flow controller (MKS Instruments). After the film deposition was finished, the system was cooled to ambient temperature within 30 minutes under vacuum.

DSC analysis

A DSC 200 Phox (Netzsch Gerätebau) was used for differential scanning calorimetry (DSC) analysis.

X-ray analysis

XRD patterns were obtained using a Bruker D8 Advance powder diffractometer with Cu K_α radiation (λ: 1.5418 Å).

Surface-analysis

The surface morphology of the Sb₂Te₃-film was investigated by AFM using an AFM Veeco diInnova and by SEM using a Jeol JSM 6510 equipped with an energy dispersive X-ray spectroscopy (EDX) device (Bruker Quantax 400). A cross-sectional sample of the Sb₂Te₃ film was prepared by using a Jeol Cross-Section Polisher (IB-09010CP).

Auger spectroscopy

SAM/AES studies were carried out on a Perkin Elmer PHI 660 Scanning Auger Microprobe System.

Photoelectron spectroscopy

Angle-resolved photoelectron spectroscopy (ARPES) was performed with linearly polarized synchrotron radiation (s + p) at the UE112-PGM2a undulator beamline of BESSY II using a Scienta R8000 hemispherical electron energy analyzer. The samples were transferred in air and the surface was prepared in situ by sputtering for 1 h at 1.5 keV at 1.2 × 10⁻⁵ mbar Ar pressure followed by 30 min annealing at 230 °C. ARPES was subsequently performed in a base pressure of 2 × 10⁻¹⁰ mbar at ambient temperature.

Results and discussion

Ditellanes R₂Te₂ are promising low-temperature Te precursors for MOCVD deposition due to their low thermal stability. Me₂Te₂, whose decomposition temperature is ≈100 °C lower compared to that of Me₂Te, has been successfully used by Kisker et al. for the growth of CdTe films at 250 °C.⁵⁶ Moreover, ditellanes are easily accessible by reaction of Na₂Te₂ with alkyl halides in liquid ammonia or other standard reactions.^57–59

The thermal properties of Et₂Te₂ and i-Pr₃Sb were investigated by DSC to identify suitable substrate temperatures for the MOCVD process. The decomposition temperature of i-Pr₃Sb (255 °C) is significantly lower compared to that of Et₃Sb (310 °C) and Me₃Sb (370 °C), which are typically used for the MOCVD process (Fig. 1). Et₂Te₂ starts to disproportionate at 150 °C with subsequent formation of elemental tellurium and Et₂Te, which then decomposes at higher temperature (260 °C) as is shown by comparing the DSC curves of Et₂Te₂ (Fig. 2, black curve) and of pure Et₂Te (Fig. 2, red curve).

Fig. 1 DSC curves of i-Pr₃Sb (red curve), Et₃Sb (black curve) and Me₃Sb (blue curve).

Fig. 2 DSC curves of Et₂Te₂ (black curve) and Et₂Te (red curve).

Sb₂Te₃ films were grown on polycrystalline Al₂O₃(0001) substrates at substrate temperatures between 350 and 400 °C using a home-made MOCVD cold-wall reactor and a deposition time of 15 minutes. An almost linear increasing growth rate with increasing substrate temperature was found in the temperature range from 350 °C to 400 °C, reaching 750 nm in 15 min at 400 °C. These findings point to a kinetically-controlled growth process, which we have also recently reported for the MOCVD deposition of Sb₂Te₃ films using (Et₂Sb)₂Te as a single source precursor.⁵⁶ The chemical composition (EDX) and the crystallinity (XRD) of the resulting Sb₂Te₃ films were not affected by the increasing substrate temperature, whereas the morphology of the Sb₂Te₃ films changed significantly. Sb₂Te₃ films deposited at 350 °C (Fig. 3a) and 375 °C (Fig. 3b) show a polycrystalline surface, most likely resulting from the low diffusion rates of adatoms as was observed with (Et₂Sb)₂Te.⁴⁷ In contrast, very smooth Sb₂Te₃ films were obtained at 400 °C (Fig. 3c). The substrate material also has a distinguished influence on the film morphology since the deposition on Si(100) substrates at 400 °C only yielded polycrystalline Sb₂Te₃ films (Fig. 3d).

Fig. 3 SEM images of the Sb₂Te₃ film deposited at 350 °C (a), 375 °C (b) and 400 °C (c) on Al₂O₃(0001) and on Si(100) (d) substrates.

The substrate typically affects the morphology of the growing film. Epitaxial films are preferentially grown on substrates which exhibit a small lattice mismatch (<7%) in order to reduce the lattice strain (so-called hetero-epitaxy).⁶⁰ However, since Sb₂Te₃ adopts a rhombohedral crystal structure, in which five monoatomic sheets in a sequence of Te–Sb–Te–Sb–Te are bonded in the c-axis direction, the Sb₂Te₃ films are always Te-terminated due to the weak van der Waals Te–Te bond. As a consequence, epitaxial V₂VI₃ films such as Sb₂Te₃, Bi₂Se₃ or Bi₂Te₃ have been grown on substrates with large lattice mismatches such as GaAs, Si(111) and Al₂O₃(0001), respectively, since the merely attached Sb₂Te₃ epilayers can achieve their own lattice parameters from the beginning of the growth.⁶¹ This growth mechanism is typically referred to as van der Waals epitaxy,⁶² which explains the formation of epitaxial Sb₂Te₃ films on Al₂O₃(0001). In contrast, Sb₂Te₃ films obtained on Si(100) always showed Sb₂Te₃ nanodisks of hexagonal shape perpendicular to the surface (Fig. 3d). Comparable findings were previously reported for Bi₂Se₃ films grown on Si(001).⁶³ The authors suggested that generating the nanodisks on the Bi₂Se₃ film is energetically more favorable than thickening the film.

Wang et al. showed that the substrate temperature is a key parameter for the MBE growth of high quality Sb₂Te₃ films on Al₂O₃(0001) substrates,⁶⁰ since both the reaction rate and the reaction kinetics are strongly influenced. The composition and surface morphology of the Sb₂Te₃ films was found to strongly depend on the substrate temperature. High-quality MBE grown films were obtained at 330 °C, whereas lower substrate temperatures gave Sb-deficient films due to the higher sticking coefficient of Te at low temperatures compared to that of Sb. Since the sticking coefficient of Te decreases at higher substrate temperatures, stoichiometric Sb₂Te₃ films were formed. However, Periano et al. also reported on the MBE growth of stoichiometric Sb₂Te₃ films at a substrate temperature of 250 °C.⁶⁴ Herein, we deposited highly stoichiometric Sb₂Te₃ films at 350 °C and at 400 °C. The film obtained at 400 °C is very flat, consisting of large Sb₂Te₃ flakes, whereas that obtained at 350 °C also showed hexagonal Sb₂Te₃ nanocrystals, which adopt orientations perpendicular to the film surface. The higher substrate temperature not only influences the decomposition kinetics of the precursor molecules, resulting in higher decomposition rates, but also affects the mobility of the adatoms, which increases with increasing substrate temperature, yielding larger Sb₂Te₃ nanoflakes. Moreover, the surface morphology (flatness) of the film improved with the increasing substrate temperature.

SEM (Fig. 4) and AFM micrographs (Fig. 5) of the Sb₂Te₃ film deposited on Al₂O₃(0001) substrates at 400 °C within 15 minutes prove the formation of smooth Sb₂Te₃ films over a large range (substrate dims. 10 × 10 mm). These results were confirmed by cross-sectional SEM microscopy (Fig. 5), showing a uniformly grown Sb₂Te₃ film of about 750 nm thickness.

Fig. 4 SEM images of the Sb₂Te₃ film deposited at 400 °C on Al₂O₃(0001). Bright spots result from small holes in the film.

Fig. 5 Cross-sectional SEM image of a Sb₂Te₃ film deposited at 400 °C on an Al₂O₃(0001) substrate showing the layered structure of the Sb₂Te₃ film. The substrate orientation is tilted by about 15°.

The AFM micrograph (Fig. 6) shows regular terraces of 1 nm height, which correspond to one quintuple layer of the layered tetradymite-type structure of Sb₂Te₃.⁶⁵ These findings indicate a layered growth as was also observed in MBE-grown Sb₂Te₃ thin films.⁶⁶ Sb₂Te₃ crystallizes in the space group R3̄m-D_3d⁵ and forms a layered, rhombohedral crystal structure.⁶⁷ It is built of anisotropic layers, in which five atomic planes are covalently bonded to the quintuple layer, that is roughly 1 nm thick and consists of five mono-atomic planes of Te(1)–Sb–Te(2)–Sb–Te(1) with 1 and 2 designating the different positions within the fivefold layer. The conventional unit cell includes three quintuple layers, which are bound by weak van der Waals forces.

Fig. 6 AFM photograph of a Sb₂Te₃ film deposited at 400 °C on an Al₂O₃(0001) substrate.

The chemical composition of the material film was analyzed by EDX analysis. Almost perfect Sb₂Te₃ films with the expected 40 : 60 (Sb : Te) molar ratios were found (standard deviation ±0.5 (Sb) and ±1.0 (Te)). Fig. 7 shows an Sb₂Te₃-film at an early stage of growth. The SEM image shows single plates which are not fully coalesced. The corresponding element mapping by scanning Auger electron microscopy (SAES) shows a homogeneous distribution of Sb and Te within the Sb₂Te₃ plates. In addition, some carbon impurities are located between the Sb₂Te₃ plates on the substrate surface.

Fig. 7 SEM (a) and SAES study (mapping of C (b), Sb (c) and Te (d)) of an Sb₂Te₃ film deposited at 400 °C on (0001) Al₂O₃ substrates; high element concentrations are indicated by bright colour.

The X-ray diffractogram (XRD) of the Sb₂Te₃ film deposited at 400 °C (Fig. 8) clearly proved the formation of highly crystalline Sb₂Te₃ films. All reflection peaks can be indexed on the basis of the structure of rhombohedral Sb₂Te₃. In addition, a small reflex of the Al₂O₃(0001) substrate is visible at 41.6°. The peaks in the diffractogram correspond to (001) planes, pointing to a c-axis oriented growth as was previously reported for Sb₂Te₃ films grown by MBE,^68–70 and MOCVD,^40–42 respectively.

Fig. 8 2θ XRD pattern of the Sb₂Te₃ film deposited at 400 °C on an Al₂O₃(0001) substrate; peaks are labelled (JCPDS-ICDD card 15-0874); the asterisk (*) corresponds to the Al₂O₃(0001) substrate.

We have characterized the band structure of the surface by photoelectron spectroscopy (PES) after transfer to the photoelectron spectrometer and in situ preparation. To the best of our knowledge, this is the first time that angle-resolved photoemission spectroscopy (ARPES) studies have been performed on MOCVD-grown Sb₂Te₃ thin films. The PE spectrum in Fig. 9 shows Sb and Te core-levels without any shifted components similar to that of Bi₂Se₃.⁷¹

Fig. 9 Core-level photoemission characterization of the in situ prepared surface. Satellite peaks are not observed.

The overview ARPES spectra of the valence band in Fig. 10 show intense structures around 0.8 eV binding energy and more faint features near E_F which are well discernible at 16 and 17 eV photon energy but not at 21 eV, where their photoemission intensity is weak due to final state effects. Fig. 11 shows, therefore, the range near E_F for photon energies from 16 to 18 eV displaying the topological surface state. This surface state has previously been studied theoretically⁹ and in ARPES experiments at selected photon energies (21.2 eV),^10,72 55 eV,¹¹ 8.44 and 24 eV,⁷² respectively.

Fig. 10 Overview ARPES spectra of the valence band. The topological surface state forming a Dirac cone near the Fermi level (0 eV binding energy) is best observed at 16 eV photon energy. Another surface state with massive character appears as the higher binding energy and is best seen at 21 eV.

Fig. 11 Magnified spectra on the dependence of the topological surface state (Dirac cone) on the photon energy. The dashed line is a linear fit to the data at 16.5 eV and results in a Dirac point ≈30 meV above the Fermi level (hole doping). This Fermi level position means bulk transport according to band structure calculations.

While experiments on cleaved bulk samples led to the conclusion that the Fermi level does not cut the surface state,⁷³ epitaxial Sb₂Te₃/Si(111) films observed slight hole doping (Dirac point ≈100 meV above E_F) and, after deposition of 0.1 monolayer Cs, the undoped state with the Dirac point at E_F. The spin texture of the surface state has been probed by spin- and angle-resolved photoemission, and an additional, topologically protected, Rashba-split surface state was found.¹¹ These results have been confirmed recently for epitaxial Sb₂Te₃/Si(111) films.⁷² In Fig. 10, this additional surface state is marked at 0.4–0.8 eV binding energy in the data taken at 21 eV photon energy. Due to its high binding energy, it does not contribute to transport. On the other hand it has been found to be much more surface localized than the topological surface state that forms the Dirac cone and its sharp, intense dispersing photoemission features here indicate the presence of a well ordered and stoichiometric surface.

Changing the photon energy in ARPES allows us principally to vary the perpendicular component of the electron wave vector, k_z.⁷⁴ In this way, we can distinguish three-dimensional dispersing features from two-dimensional features. The identical E vs.k_|| dispersion in Fig. 11 proves the two-dimensional character. The intensity, however, is subject to the k_z-dependence of the parent bulk band structure, and this is seen at 18 eV as a shift of maximum intensity so that at 22 eV (not shown) the surface state is not discernible any longer. The intensity seen near the Fermi energy at 16 eV moves with increasing photon energy to lower binding energy and has crossed E_F at 18 eV. This intensity stems directly from bulk valence band states which exist in the complete lower Dirac cone up to the Dirac point as is evident from, e.g., surface-projected bulk band calculations.¹¹ Due to these states, the topological transport regime will require the Fermi level to be in the upper Dirac cone. The linear dispersion predicted by theory has been fitted to the data at 16.5 eV in Fig. 11. Extrapolation of the data yields a Fermi velocity of (2.72 ± 0.14) × 10⁵ m s⁻¹, roughly in agreement with previous measurements from single crystals [(3.8 ± 0.2) × 10⁵ m s⁻¹] and density functional theory (3.2 × 10⁵ m s⁻¹),¹¹ and a Dirac point ≈30 meV above E_F. The surface is, therefore, slightly hole doped. The calculation shows that bulk valence band states reach up to the Dirac point so that hole doping indicates that the Fermi level cuts through bulk valence band states rendering the transport bulk like.¹¹

Electrical measurements were performed with a PPMS-system (DynaCool QuantumDesign) using the standard Electronic Transport Option. By applying an alternate current of 1 mA with a frequency of 12.2 Hz and a sampling rate of 0.5 s⁻¹, the resistance was measured in the range of 4 to 400 K with a heating rate of 10 K min⁻¹. By making equally spaced 4-point contacts and choosing the contact spacing much larger than the film thickness, the sheet resistivity can be calculated by the following equation:⁷⁵

I is the applied current at the outer contacts and V is the measured voltage drop between the inner contacts.

The measured sheet resistivity (Fig. 12) shows a metallic like behaviour with a dominating contribution of the carrier transport in the bulk, and a monotonic increase with increasing temperature as was previously observed for ALD-grown Sb₂Te₃ thin films.²⁴ Here the sheet resistance increases by one order of magnitude (10×) over the temperature range from 2 to 400 K, whereas the resistance of the ALD film increased moderately by the factor of 1.5 over the same temperature range. The MOCVD films exhibit a higher carrier concentration (>10¹⁹ cm⁻³) correlated with intrinsic defects, most likely antisite defects as is common for Sb₂Te₃. Furthermore, we assume that the carrier concentration is nearly constant over the complete temperature range and the strong enhancement of the resistance over temperature is related to the reduction of the carrier mobility. Due to the high doping concentration the enhancement of the carrier concentration by thermal excitation of carriers over the bulk-bandgap is not observed by the resistance curve (Fig. 12) or has a negligible impact. In the case of ALD grown Sb₂Te₃ films a stabilisation of the resistance at least around 350 K has been reported.²⁴

Fig. 12 Resistivity against temperature. The inset graph shows the I–V-curve measured at room temperature with 1 mA at 12.2 Hz.

Conclusions

The layer-by-layer MOCVD deposition of highly crystalline Sb₂Te₃ films was achieved at substrate temperatures between 350 and 400 °C using i-Pr₃Sb and Et₂Te₂. Even though both precursors start to decompose at lower temperatures as was shown by DSC, higher temperatures are necessary in order to produce smooth and crystalline Sb₂Te₃ films. The film growth occurs under kinetic control, reaching deposition rates as high as 3 μm per hour at 400 °C. The electrical sheet resistivity was found to monotonically increase with increasing temperature. ARPES measurements for the first time showed the detailed dispersions of the bulk valence band and the topological surface state for MOCVD grown Sb₂Te₃ films with a quality no less than for optimized bulk single crystals. The surface state dispersion gives a Dirac point roughly 30 meV above the Fermi level leading to hole doping and the presence of bulk valence states at the Fermi energy. Additional Bi doping of the Sb₂Te₃ films is currently under investigation in order to further modify the valence band and the topological surface states.

Acknowledgements

Stephan Schulz (StS), Oliver Rader (OR) and Kornelius Nielsch (KN) acknowledge financial support by the DFG (priority program SPP 1666), the University of Duisburg-Essen (StS), the Karl-Vossloh-Stiftung (KN), and the University of Hamburg (KN). Partha Sarathi Mandal is supported by the Impuls- und Vernetzungsfonds der Helmholtz-Gemeinschaft. The authors like to thank PD Dr Nils Hartmann, University of Duisburg-Essen, for SAES studies.

Notes and references

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