Growth of crystalline WSe2 and WS2 films on amorphous substrate by reactive (Van der Waals) rheotaxy

Solar ~ Materials and Solar Cells

ELSEVIER

Solar Energy Materials and Solar Cells 44 (1996) 457-470

Growth of crystalline W S e 2 and WS 2 films on amorphous substrate by reactive (Van der Waals) rheotaxy T. Tsirlina a, S. Cohen a, H. Cohen a, L. Sapir a, M. Peisach a, R. Tenne a,*, A. Matthaeus b, S. Tiefenbacher b W. Jaegermann b E.A. Ponomarev c, C. L6vy-C16ment c a Department of Materials and Interfaces, Department of Chemical SerL'ices, Department of Physical Services, Weizmann Institute, Rehovot 76100, Israel b Abteilung Grenzfliichen, Hahn-Meitner lnstitut, 14109 Berlin, Germany c LPSB, CNRS, UPR 1332, 1-Place Aristide-Briand, Meudon 92195, France

Abstract Thin films of metal dichalcogenide compounds with a layered structure, such as MoS 2 (WSe2), play an important role in a number of technologies, like solid lubrication, experimental photovoltaic cells, etc. Such films usually adopt a type-I texture, in which case the c-axis of the crystallites is parallel to the substrate plane. However, for the aforementioned applications, type-II texture, where the c-axis of the crystallite is perpendicular to the substrate, is required. We have recently demonstrated a novel growth technique (Van der Waals rheotaxy, VdWR) which yields a crystalline film having exclusively type-II texture on amorphous (quartz) substrate. In the present work superior crystalline, optical and electronic properties of the overlying WSe 2 (WS 2) film together with an improved adhesion of the film to the quartz substrate are obtained by replacing the ultra thin Ni film with a N i / C r film. Keywords: Layered compounds; Rheotaxy; Metal dichalcogenide

1. Introduction Layered metal-dichalcogenide compounds, e.g., 2H-MoS 2, possess a highly anisotropic structure. Each layer consists of a metal layer sandwiched between two

* Corresponding author. 0927-0248/96/$15.00 Copyright © 1996 Elsevier Science B.V. All rights reserved. PH S0927-0248(96)00048-7

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chalcogen layers in a trigonal prismatic structure. Strong covalent bonds exist between the metal and chalcogen atoms in the plane of the layers. The layers are stacked together by weak Van der Waals type interactions. The unit cell of 2H-MoS 2 consists of two MoS 2 layers as a repeat unit, and they are packed in a hexagonal lattice. This structure provides an atomically smooth and inert Van der Waals surface ( ± c), while the dangling bonds along the IIc (prismatic) face are highly reactive. Photovoltaic (PV) application of these materials was originally proposed by Tributsch [1], and encouraging results were reported in a number of publications, using single crystal materials [2-5]. Photovoltaic (photoelectrochemical) cells of respectable efficiencies (r/) could be obtained by using the Van der Waals face ( i c orientation-type II texture) of n-type WSe 2 (~7 = 13%) [3], p-type WSe 2 (7/= 8%) [4], and InSe (r/= 11%) [5] of single crystal materials. Since both crystallinity and texture of thin films of layered compounds could not be controlled thus far, poor photoeffects for thin film cells [6] and mediocre lubricating power were reported, especially in ambient conditions, where water adsorption to the dangling bonds on the (110) face could not be totally averted. Crystalline films can be grown layer by layer (Frank Van der Merwe mode) on a lattice matched crystalline substrate, using such techniques as molecular beam epitaxy (MBE). The price-tag for films prepared in this way is prohibitive for many applications, including photovoltaics, thin film transistors, etc. Therefore, growth of crystalline films on amorphous substrates is a problem of prime technological importance, and an intriguing scientific problem. For a general discussion of this issue one may consider Ref. [7]. Single crystals of layered compounds grow (at elevated temperatures), layer by layer with their c-axis perpendicular to the substrate. Contrarily, attempts to grow highly textured-type II films at moderately low temperatures ( < 750°C), have been unsuccessful [6]. A likely explanation is that, being a thermally activated process, surface diffusion is inhibited at low temperatures which favors nucleation of many embryonic crystallites. The reactive dangling bonds on the prismatic face of these microscopic crystallites bind chemically to the substrate, leading to a preferred type I texture. In support of this hypothesis, nanocrystallites of such compounds were recently shown to "self-passivate" the dangling bonds on the prismatic face by forming hollow fullerene-like clusters and nanotubes [8]. These observations emphasize the need to control surface reactivity in attempting to grow films with large crystallites and type II texture. Moreover, atomically smooth substrates have smaller barriers towards surface diffusion and hence they are expected to favor type II texture even at relatively low growth temperatures. Most thin films are grown on a rigid substrate and hence the solid-solid interface between the substrate and the film induces a large strain. An alternative approach, which is well documented in the literature [9], is to use rheotaxy, i.e., high temperature growth of solid films on molten substrates. Recrystallization of Si (Ge) thin films by annealing in the presence of A1-Si eutectic resulted in large crystallites [10]. It was also used before to grow highly oriented metal dichalcogenide crystals on Te melt (sometimes designated the Marangoni effect) [11]. Quasi-rheotaxy is a process in which, e.g., a CdTe film is grown on top of a thick, low melting point metal film (e.g., Bi, Pb) [12]. Taking advantage of the lower melting temperature for surfaces compared with the bulk phase, the sample was heated to a few degrees below the melting temperature of the

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metal film, which resulted in recrystallization of the CdTe overlayer. However, the quality of the polycrystalline film was a far inferior to that of a single crystal. Koma and his co-workers have demonstrated [13], that MBE growth of perfectly crystalline films of layered compounds like WS% could be obtained on crystalline substrates with a large lattice mismatch, like InSe (Van der Waals epitaxy, VdWE). In this case the requirement for lattice match between the substrate and the film can be relaxed. This information suggested to us that a molten Se (S) alloy with an atomically smooth interface could serve as an ideal substrate for 2-D crystallization of layered compounds. Recently, 1 txm thick type II oriented films of WS%, albeit with a relatively small crystallite size ( < 1 Ixm), were prepared by us, using high temperature selenization of a thick (200-500 nm) WO 3 film, which was grown atop of a thin (3-10 nm) Ni-Se molten interlayer [14]. However, the film thickness was limited ( < 0.5 Ixm) and a poor adhesion to the quartz substrate was observed. The adhesion problems were attributed to two separate mechanisms: the large difference in thermal expansion coefficients between the quartz substrate and the overlayer film. Moreover, being a low energy surface, the Van der Waals face of the bottom of the WSe 2 film binds weakly to the substrate. The first problem could be partially remedied through better thermal management of the heating cycle of the selenization process. To resolve the second problem, a good binder was searched for. The melting point of Cr-Ni (Ni = 80%) [15]a and Cr-Se [15]b compounds exceeds the relevant process temperatures ( < 950°C). Therefore, the Cr~Sey crystallites were expected to pin the WSe 2 (WS 2) film to the quartz substrate and promote its adhesion. We report here, that by replacing the Ni with a Ni-Cr alloy as an ultrathin interlayer between the substrate and the WO 3 precursor, and through careful control of the heating cycle of the selenization (sulfidization) process, WSe 2 (WS 2) films with a high degree of crystallinity and texture is obtained on amorphous quartz substrates. In this process, a molten Ni 3Se2 (Ni3 S 2) interlayer [ 15]c serves as an atomically smooth substrate, which induces ordering of the first WSe 2 (WS 2) layer, on top of which the entire WSe 2 film grows. Furthermore, the Ni-containing molten film also plays the role of a catalyst for the initiation of the conversion of the oxide precursor into WSe 2 (WS2). Finally, the Cr-Se (Cr-S) crystallites pin the WS% (WS 2) film to the substrate, leading to a very robust adhesion between the film and the substrate.

2. Experimental Fig. 1 shows a schematic illustration of the growth process (accronimed Van der Waals rheotaxy, VdWR). The precursor consisits of a thin (ca. 20 nm), probably discontinuous Ni-Cr (80/20%) film, covered by a 3-6 Ixm thick WO 3 film (Fig. la). A stream of forming gas (5%H2/95%N2; 100 ml/min) is used during the annealing process. Se pressure is maintained in the reactor by preheating Se shot at 350°C in an adjoining upstream chamber. The sample is first heated to 760°C, during which time the selenium (sulfur) vapors initially react with the nickel to form the eutectic compositionNi3Se 2 [9]. The sample is further annealed for 30 min at this temperature, which leads to

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the melting of the Ni3Se 2 film and initiation of the tungsten selenide synthesis, where the WO 3 film is converted into WOxSe3_x, as shown in Fig. lb. Subsequently, the temperature is increased to 950°C and is maintained at this temperature for another 60 rain. At this temperature the synthesis of the WSe 2 film is completed (Fig. lc). The WSe 2 molecules assemble on the molten Ni3S % substrate in a mode which is reminiscent of the self-assembly of organic monolayers on water surface [16]. From that point on more WSe 2 layers grow and a crystalline and textured film is obtained. The conversion of the mixed oxide into WSe 2 is catalyzed by the underlying eutectic. Therefore the conversion of WO 3 into WSe 2 starts at the Ni3Se2/WO 3 interface, rather than from the top surface of the oxide film down to the substrate. The sample is allowed to cool-down at a rate of 50°C/h under the same atmosphere. The WS 2 films were synthesized in a sealed quartz ampoule in the presence of excess sulfur for 15 min at 750°C followed by 90 min anneal at 950°C. In order to obtain information on the sample composition at depths greater than the range which is accessible to 2 MeV 4He+ ion beams (ca. 1-2 Ixm), the films were analyzed by Rutherford backscattering of protons (PB). Measurements were made at a

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backscattering angle of 140 ° and an incident proton beam energy of 2200 keV. To limit thermal deterioration of the films, relatively low currents of ~ 10 nA were used. Analysis could be achieved at depths exceeding 3 p~A. However, the N i / C r profile could not be accurately determined in this technique, as a result of the degradation of the bombarding beam energy by the overlying WSe 2 (WS 2) layer. A CAMECA IMS 4F with a cesium source was used for the secondary ion mass spectrometry (SIMS) measurements. Philips 515 scanning electron microscope (SEM) equipped with Tracor 5500 energy dispersive (EDS) analyzer was used for the imaging/chemical analysis of the samples. Kratos Analytical AXIS-HS X-ray photoelectron spectrometer with Ar + ion gun was used for surface analysis/depth profiling. Topometrix TMX2010 Discoverer atomic force microscope (AFM) was used for imaging with submicron down to atomic resolution, using microfabricated Si or Si3N 4 cantilevers with integrated tips. X-ray powder diflYactometry (XRD) was done by Rigaku Rotaflex RU-200B. For photoconductivity measurements, two contacts were established on the front surface of the film with silver paint. A 3 V bias was applied between the electrodes and the change in the resistance was measured by a lock-in technique. The modulation of the resistance due to the ac illumination was 0.4%.

3. Results and discussion Fig. 2a shows a comparison between the XRD spectrum of a typical W S e 2 film, prepared by VdWR and a single crystal. The remarkable intense and narrow (001) diffraction peaks of the film, attest to the high crystalline quality and the exclusive type II texture of the film. The inset shows the weak diffraction peaks of the two samples. The spectrum of the thin film shows, in addition to the weak WSe 2 diffraction peaks, those of Ni3Se 4 and Cr2Se 3. Fig. 2b shows a similar XRD spectrum of a typical WS 2 film. Here again, sharp and intense diffraction peaks, quite comparable to those obtained from a single crystal material, are observed. The intensity ratios of the various peaks (001/002) for 1 = 4,6,8 . . . . are quite similar to those of a single crystal sample for both films. Note that the finite thickness of the film may lead to differences between the peak intensities of films and the crystal. The lattice parameters of the WS 2 film, which was found from the XRD pattern, are: a = 3.160 + 0.002A and c = 12.353 + 0.007A, in agreement with the values for single crystal [17]. The lattice parameter c of the WSe 2 film was found to be 12.96 + 0.01A, in good agreemement with the literature value [18]. The morphology of the film investigated by AFM was found to be rather buckled, as demonstrated in Fig. 3a for WSe 2. Strain effects, due to mismatch in thermal expansion coefficients are commutative. Therefore the size of the sample was thought to influence the morphology of the film. For that reason, a sample with reduced size (0.3 × 0.3 cm 2 instead of the usual 1 × 1 cm 2 size) was prepared. Indeed this film was found to exhibit smoother morphology (Fig. 3b) than the film with larger surface area. Also, WS 2 films exhibited somewhat smoother morphology than the WSe 2 films (compare Fig. 3c to Fig. 3a). Atomic resolution could be easily obtained under ambient conditions on various flat locations of the WSe 2 and WS 2 surfaces. Furthermore, the orientation of the crystal

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planes, as r e v e a l e d by the A F M on the W S e 2 film, was p r e s e r v e d o v e r a few m m range, w h i c h suggests little m o s a i c i t y and a large crystallite size in this film. The lattice p a r a m e t e r (a) for the W S e 2 film was d e d u c e d to be 3.3A, in a g r e e m e n t with literature value for this p a r a m e t e r [18].

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Fig. 3 (continued). The film composition and the depth distribution of the elements were investigated. Fig. 4 shows typical PB spectra of the films. In addition to the expected observation of surface W and S (Se), Ni and oxygen were clearly observed on the film surface. Because the concentration of W and Se (S) largely exceeded that of Ni, the depth profile of the latter could not be accurately determined by the current method [19]. Evidence for the presence of a large concentration of Cr at the film/quartz interface (Fig. 5) was also obtained from the PB. Both Si and O from the quartz substrate were observed at depths corresponding to the approximate thickness of the tungsten compound. XPS analyses of a WS 2 film reveal that the surface is stoichiometric and not oxidized. No Ni or Cr could be detected by XPS on the film surface indicating surface concentration < 0.1%. The surface concentration of Ni varies from sample to sample (vide infra). The large concentration of surface oxygen and carbon, in the case of WS 2, can be attributed to the accumulation of organic residues on the grain boundaries of the WS 2 film. The depth profile obtained from the SIMS analysis of the WS 2 film, is shown in Fig. 5. These results indicate that the molten Ni compound outdiffuse and is distributed quite evenly throughout the film thickness. Contrarily the Cr remains mostly at the film/quartz interface. The constant W / S ratio provides a clear evidence that the Ni is not dissolved in the film itself but concentrates at the grain boundaries. The spatial distribution of the N i / C r in the WSe 2 film was analyzed with SEM equipped with an EDS analyzer. Fig. 6a-6c show SEM images of the surface of a typical WSe 2 film, while Fig. 6d shows SEM image of a WS 2 film. Fig. 6a shows a low

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magnification image, which gives a general view of the film. In this magnification the W S e 2 film appears quite smooth and featureless. Cracks divide the film into uniform zones with a typical size of 50 × 50 Ixm 2. The cracks are decorated with bright particles,

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time/s Fig. 5. Depth profile of the W S 2 film obtained by secondaryion mass spectrometry(SIMS). Rate of sputtering ca. 5 nm/s. which upon closer examination by local EDS analysis are identified as NigSe 4. This composition coincides with the stable phase of the N i - S e phase diagram at 950°C [15]c. The distribution of these particles is not limited to the cracks only, as shown by the magnified image in Fig. 6b. Particles with a cauliflower morphology and the composition Ni3Se 4 are observed (arrowed). This morphology indicates that the molten N i - S e alloy pierced through the film and solidified. The superficial Ni 3Se4 particles could be fully removed by a 1 min etch in nitric acid solution (10%). The Ni free film is shown in Fig. 6c. The film is perforated at sites which were previously occupied by the Ni3Se 4 particles. The remaining elongated bright particles on the film surface are WSe 2 platelets, no evidence for Ni was left after this treatment, and the adhesion of the film to the substrate did not show any indication of deterioration. Since the typical crystallite size in the WSe 2 film ( > 50 txm) largely exceeds the typical crystallite size in the case of the WS 2 film (3 Ixm), it is likely that the spatial distribution of Ni on the surface of the latter film, is much more uniform than for the WSe 2 film. Lately, better control of the thermal cycle resulted in crack-free WSe 2 films which do not have any Ni3Se 4 particles on the surface. Microwave photoreflectivity measurements were carried-out on both kinds of films. First order (1054 nm) and second order (532 nm) lines of the Nd-YAG laser were used for the excitation of the bulk and near surface region, respectively. Analysis according to a previously published model [20], was undertaken for the calculation of minority

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Fig. 6. Scanning electron microscopy of WSe 2 films: (a) Low magnification, in which the cracks are clearly exhibited; (b) Large magnification shows the Ni3Se ~ crystallites with a cauliflower shape in bright contrast; (c) After etching of the film with nitric acid, perforations are easily observed on the film surface; (d) WS, film.

carrier's lifetime and mobility. For the WS 2 film, the value of the mobility varied between 9-15 c m 2 / s V , while for the WSe 2 film, the value varied between 80-100 c m 2 / s V . The value obtained for the WSe 2 film compares favorably with single crystal data [21], and reflects the high quality of the film. Lifetimes on the order of a few hundred ~s, comparable to a single crystal, were measured for these films. Further improvements are expected through careful fine-tuning of the reaction conditions. Fig. 7 shows the the photoconductivity spectrum of WS 2 (a) and WSe~ (b). The inset shows a model calculation of the indirect bandgap - - 1.3 eV (WS 2) and 1.26 eV (WSe2), which is in good agreement with the literature [21]. It can be noticed that the position of the A and B excitons in the photoconductivity spectrum of WSe 2 [22] coincide with the dips (756 and 563 nm, respectively) in the spectrum. Photoconductivity is obtained through flow of photoexcited electrons and holes in opposite direction and their collection by the positive and negative electrodes, respectively. The excitonic interaction reduces the probability for charge separation, and consequently the photoresponse of the sample in this spectral range declines. The present results show that the Ni distribution on the front surface of the film is not even. Most of the Ni is concentrated in the cracks and in cauliflower-like crystallites. No

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evidence for a mixed phase between the elements W - N i - S e , has been obtained. On the contrary, the 1:2 ratio between W and Se (S), found by XPS analysis, strongly suggests that the dissolution of Ni within the film is inferior to a fraction of a percent, which is the detectivity limit of the XPS. The high mobility values for minority carriers in both films and their long lifetimes, further indicate that the solubility of Ni in the WS 2 and WSe e crystallites is exceedingly low (on the ppm level), and has no deleterious influence on the electronic properties of the film. The aforementioned growth mechanism (VdWR) is not the only possible way to explain the favorable role of Ni in the process. Alternatively, the Ni3Se 2 melt could serve as a flux which induces a lateral (2D) growth of the incipient nuclei. Thus, the Ni is expected to remain at the edges of the crystallites until they coalesce and form grain boundaries (VLS growth). A combination of the two growth mechanisms can not be fully excluded. Further research is necessary to resolve that issue.

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Together with the high degree of crystallinity and the excellent texture of the film, a rough morphology, on the submicron scale, is revealed for both kinds of films. One possible explanation for this, is the demixing and crystallization of the underlying solid Cr2S % (Cr2S3). However the total amount of chromium in the film is too small to account for a surface roughness of this magnitude. Another possible explanation for the rough morphology, is that it emerges from the differences between the thermal expansion coefficients of the substrate and the film. The scaling of the average roughness with the sample dimensions suggests that this is a likely explanation, and a more complete series of experiments are underway to resolve this issue. The favorable influence of the chromium compound on the texture and particularly on the adhesion of the film, is quite remarkable. In fact, when the ultra thin layer consisted of Ni only, WSe 2 films thicker than 1 ~m could not be obtained. The films either did not stick to the substrate or they disintegrated into a fine powder. Using a N i - C r interlayer, films as thick as 10 p~m have been successfully synthesized, reproducibly. The remaining major obstacle facing the developers of a thin film PV cell, which is based on layered compounds, is the back contact. In the absence of a reliable, low resistivity back contact, no current can be drawn from the cell, due largely to the high resistivity of the quartz substrate. However, low resistivity contacts, which can tolerate the harsh environment of the reactor, cannot be easily found or developed. Alternatively, one may attempt eutectics of lower melting point, such as SnS 2 or InSe [23], an idea that has only been briefly pursued, so far.

Acknowledgements We are grateful to Dr. E. Galun for his help with the work. This work was supported by the Israeli Ministry of Energy and Infrastructure, contract No. 9 5 / 1 1 / 0 0 8 ; by a CEE contract JOU2-CT93-0352; and by the BMFT contract No. 0329639A.

References [1] H. Tributsch, J. Electrochem. Soc. 125 (1978) 1086. [2] G. Kline, K. Kam, D. Canfield and B.A. Parkinson, Solar Energy Mater. 4 (1981) 301. [3] (a) R. Tenne and A. Wold, Appl. Phys. Lett. 47 (1985) 707; (b) D. Mahalu, L. Margulis, A. Wold and R. Tenne, Phys. Rev. B45 (1992) 1943. [4] M. V6gt, M.Ch. Lux-Steiner, H.-P. Schweikardt, P. Dolatzoglou, M. Keil, W. Reetz and E. Bucher, Proc. 4th Int. Nat. Photovoltaic Science and Engineering Conf., IREE, 1989, pp. 493-498. [5l A. Segura, M.Cm. Martinez-Tomas, B. Mari, A. Casanovas and A. Chevy, Appl. Phys. A44 (1987) 249. [6] (a) A. Aruchamy and M.K. Agarwal, in: A. Aruchamy (Ed.), Photoelectrochemistry and Photovoltaics of Layered Semiconductors (Kluwer, Dordrecht. 1992) pp. 319-347; (b) E. Schmidt, F. Weft, G. Meunier and A. Levasseur, Thin Solid Films 245 (1994) 34; (c) J. Moser and F. Levy, J. Mater. Res. 7 (1992) 734. [7] E.I. Givargizov (Ed.), Oriented Crystallization on Amorphous Substrates (Plenum Press, New York. 1991). [8] (a) R. Tenne, L. Margulis, M. Genut and G. Hodes, Nature 360 (1992) 444; (b) L. Margulis, G. Salitra, R. Tenne and M. Talianker, Nature 365 (1993) 113; (c) Y. Feldman, E. Wasserman. D.J. Srolovitz and R. Tenne, Science 267 (1995) 222.

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