Pin-Chun Shen M.S. Photonics and Optoelectronics National Taiwan University, 2014

Transcription

1 Large-area CVD Growth of Two-dimensional Transition Metal Dichalcogenides and Monolayer MoS 2 and WS 2 Metal oxide semiconductor Field-effect Transistors by Pin-Chun Shen M.S. Photonics and Optoelectronics National Taiwan University, 2014 Submitted to the Department of Electrical Engineering and Computer Science in Partial Fulfillment of the Requirements for the Degree of Master of Science in Electrical Engineering at the Massachusetts Institute of Technology June Massachusetts Institute of Technology. All rights reserved. Signature of author Department of Electrical Engineering and Computer Science May 12, 2017 Certified by Jing Kong Professor of Electrical Engineering Thesis Supervisor Accepted by Leslie A. Kolodziejski Professor of Electrical Engineering Chair, Department Committee on Graduate Students 1

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3 Large-area CVD Growth of Two-dimensional Transition Metal Dichalcogenides and Monolayer MoS 2 and WS 2 Metal oxide semiconductor Field-effect Transistors by Pin-Chun Shen Submitted to the Department of Electrical Engineering and Computer Science on May 12, 2017 in Partial Fulfillment of the Requirements for the Degree of Master of Science in Electrical Engineering Abstract Two-dimensional semiconducting materials such as MoS2 and WS2 have been attractive for use in ultra-scaled electronic and optoelectronic devices because of their atomically-thin thickness, direct band gap, and lack of dangling bonds. Methods for large-area growth of 2D semiconducting materials are needed to bring them to practical applications. This thesis aims to develop reliable methods for growing high-quality monolayer MoS2 and WS2 by CVD and explore their intrinsic electrical transport properties for electronic and optoelectronic device applications. The as-grown monolayer MoS2 and WS2 exhibit n-type semiconducting behavior with excellent optical properties. Various techniques are employed to characterize the CVD-grown materials, including photoluminescence, UV-visible absorption, Raman spectroscopy, X-ray photoelectron spectroscopy, and atomic force microscopy. Moreover, the electronic transport characteristics of single-layer CVD-grown MoS2 and WS2 field-effect transistors with a back-gated configuration are demonstrated. Thesis supervisor: Jing Kong Title: Professor of Electrical Engineering 3

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5 Table of Contents Abstract... 3 Table of Contents... 5 List of Figures... 8 Chapter 1. Introduction Beyond Silicon: Opportunities in Two-dimensional Materials Significance of This Work Chapter 2. Physics of Transition Metal Dichalcogenides Atomically Thin Structure Electronic Structure Electron Transport and Scattering Phonon Scattering Coulomb Scattering Surface Phonon Scattering and Roughness Scattering Structural Defect-dependent Properties Chapter 3. CVD Growth of Monolayer MoS Monolayer MoS 2 Growth Characterizations Optical Microscopy Characterizations Raman Spectroscopy Characterizations Atomic Force Microscopy Characterizations Photoluminescence and Optical Absorption of MoS Photoluminescence UV-visible Absorption Defect Characterization of MoS PL and Raman Intensity Mappings of MoS 2 Triangular Domains PL Mappings of MoS 2 Grain Boundaries Graphene/MoS 2 In-plane Heterostructures Lateral Heterostructure Growth Optical and AFM Images PL and Raman Intensity Mappings Chapter 4. CVD Growth of Monolayer WS Monolayer WS 2 Growth WS 2 Growth Evolutions Raman Spectroscopy Characterizations Photoluminescence of WS X-ray photoelectron spectroscopy of WS Chapter 5. Monolayer MoS2 and WS2 Field-effect Transistors MoS 2 Device Fabrication TMD Transfer E-beam Lithography, Electrode Deposition, and Lift Off Electrical Transport Properties of MoS 2 FETs with Ni contacts

6 5.3 Enhancement of Electron Mobility in MoS 2 via a High-k Gate Dielectric Electrical Transport Properties of WS 2 FETs with Ni contacts Chapter 6. Conclusions and Future Work Conclusions Future Work Defect Characterization and Engineering Strategies to improve the electrical performance of CVD-grown 2D TMD transistors References

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8 List of Figures Figure 2-1. (a) The lattice structure of monolayer MX2 and (b) the hexagonal planes of M and X atoms [3, 7] Figure 2-2. Summary of electronic properties of TMD materials [8] Figure 2-3. (a) Comparison for PL spectra of monolayer and bilayer MoS2. (b) The dependence of MoS2 bandgap energy on its thickness. (c) Absorption spectra for monolayer and bilayer MoS2 [9] Figure 2-4. Band structures for bulk, bilayer, and monolayer MoS2 and WS2 [10] Figure 2-5. The dependence of carrier mobility in monolayer MoS2 on (a) temperature and (b) carrier density (calculated from first-principles density functional theory). (c) The scattering contributions from charged impurities and the total mobility of MoS2 due to the combined effect [12, 13] Figure 2-6. Typical defects in 2D TMD materials [24] Figure 2-7. (a) Formation energies of various point defects as functions of sulfur chemical potential. Sulfur vacancy (Vs) has the lowest formation energy (~ 2 ev). (b) Schematic depiction of the defect levels within MoS2 band gap [25] Figure 2-8. (a) Optical microscope image of a WSe2/MoS2 hetero-bilayer. (b) Normalized PL and absorbance spectra for monolayer MoS2, WSe2, and their hetero-bilayers. (c) SEM images of MoSe2/WSe2 heterostructures. (d) PL intensity map of a MoSe2/WSe2 lateral heterostructure (Scale bars, 2 μm) [21-23] Figure 3-1. (a) Molecular structure of PTAS and (b) PTAS aqueous solution Figure 3-2. (a) A schematically illustration of the CVD setup for growth of MoS2 films and (b) the temperature profile used for monolayer MoS2 growth Figure 3-3. (a) Photograph of centimeter-scale monolayer MoS2 grown on SiO2/Si wafer. (b) Optical images of the CVD-grown single-crystal triangular domain of MoS2, (c) the completely continuous film of MoS2, and (d) the transition region between MoS2 triangular flakes and continuous film Figure 3-4. Raman spectra of as-grown monolayer MoS2 (a) triangular flakes and (b) films measured at different regions Figure 3-5. AFM images of monolayer MoS2 (a) films and (b) triangular grains. (c) AFM step height profile of a typical single-layer region Figure 3-6. Room temperature photoluminescence of the as-grown monolayer MoS2 films. Raman spectrum is also displayed in this figure in the unit of photon energy. An excitation wavelength of 532 nm was used for all PL and Raman spectra in this thesis Figure 3-7. Optical absorption spectra of as-grown monolayer MoS2 films on (a) mica, (b) fused silica, and (c) c-plane sapphire Figure 3-8. Intensity mappings of (a) PL emission wavelength and (b) Raman E2g and (c) A1g mode frequencies of a representative as-grown MoS2 single crystal Figure 3-9. (a) PL Intensity mapping for a region containing a linear-shape grain boundary and (b) optical image of as-grown monolayer MoS2 film. The region corresponding to (a) are highlighted by a blue dotted box Figure Schematic illustration of the in-plane heterostructure of Graphene/MoS2 [32]

9 Figure (a)-(c) Typical optical images of as-grown in-plane heterostructure of Graphene/MoS2. (d) The AFM image of the region highlighted by yellow circle in (a). (d) PL and (e) Raman spectra of the grown MoS2 monolayer wrapping graphene Figure PL and Raman intensity mappings of the graphene/mos2 junction corresponding to the region highlighted by the yellow circle Figure 4-1. Schematic illustrations of (a) the experimental setup and (b) the temperature profile used for monolayer WS2 growth Figure 4-2. (a)-(d) Optical images of WS2 grains grown by different recipes. The WS2 monolayer grown by recipe A is the focus of this chapter. (e) The AFM image of WS2 monolayer grown by recipe A and (f) the thickness profile along the yellow line in the AFM image Figure 4-3. A schematic illustration of the CVD-grown WS2 evolution with growth time at 800 C Figure 4-4. Room-temperature Raman spectrum of as-grown WS2 flakes using a 532 nm laser excitation Figure 4-5. Room-temperature photoluminescence spectrum of as-grown monolayer WS2 (under 532 nm excitation) Figure 4-6. Chemical composition analysis of as-grown WS2 using X-ray Photoelectron Spectroscopy Figure 5-1. Schematic process flow for fabrication of back-gate single-layer MoS2 field-effect devices. The WS2 devices are fabricated using the same process flow Figure 5-2. Schematic and optical image of the single-layer back-gate MoS2 EFT. Symmetrical 50 nm Ni contacts are defined on the CVD-grown monolayer MoS2 by e-beam lithography, evaporation, and lift-off Figure 5-3. Room temperature transfer characteristic for the single-layer MoS2 FET with Vds = 500 mv (L ~ 1 m and W ~ 10 m). Back-gate voltage is applied to the substrate Figure 5-4. Ids-Vds characteristic for the MoS2 device acquired for different values of Vbg Figure 5-5. Low-temperature Ids-Vbg transfer curve of the MoS2 FET acquired at 77 K Figure 5-6. The Schematic of a back-gated MoS2 FET using a 25 nm HfO2/ 300 nm SiO2 hybrid gate dielectric Figure 5-7. Transport properties of single-layer MoS2 FETs (L ~ 1 m and W ~ 6 m) on a HfO2/SiO2 substrate Figure 5-8. Output characteristics (Ids-Vds) of the single-layer MoS2 FETs at different back-gate voltages Figure 5-9. Schematic and optical image of the single-layer back-gate WS2 FET. Symmetrical 50 nm Ni contacts are defined on the CVD-grown monolayer WS2 by e-beam lithography, evaporation, and lift-off Figure (a) Saturation Ids-Vds behavior of a representative monolayer WS2 FET with Ni contacts (L ~ 1 m and W ~ 15 m). (b) Output performance of the devices at small fields Figure Saturation Ids-Vds behavior of a representative monolayer WS2 FET with Ni contacts (L ~ 1 m and W ~ 15 m). On/Off ratio of the device is larger than 10 6 at room temperature. Inset: The Ids-Vds curve plotted in the logarithmic scale

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11 Chapter 1. Introduction 1.1 Beyond Silicon: Opportunities in Two-dimensional Materials Two-dimensional (2D) materials, a class of materials possessing ultimate limit of thinness in vertical dimension, and representing the thinnest artificial materials in the universe, have demonstrated themselves as a fertile ground for discovering exotic phenomena in condensed matters and as a promising platform to push the frontier of semiconductor technology beyond the Moore s law. Semiconducting materials, in particular silicon (Si), establish the foundation of modern electronics. The configuration of conventional field-effect transistors (FETs) is composed of a semiconducting channel contacted with source and drain electrodes and a gate electrode that can create a vertical electric field coupling with the carriers on the channel material surface to control the channel conductivity. Subsequent decrease in dimension of Si-based FETs, however, is approaching limits of 5 nm gate lengths due to severe short channel effects such as increased direct source-to-drain tunneling current, loss of gate control, and increased subthreshold swing (SS). New device architectures and novel channel materials for next-generation FETs have therefore been intensively searched. In this regard, 2D monolayer transition metal dichalcogenides (TMDs) come into play because of their favorable electrostatic properties. Due to the atomically ultra-thin nature, 2D TMDs can approach the ideal effective screening length and can be operated beyond the quantum capacitance limit, offering extremely high degree of gate electrostatic control for ideal subthreshold swing (60 mv/decade) and low-power consumption. Moreover, most 2D TMDs possess a higher band gap and heavier carrier effective masses than those of Si. These features are advantageous to reduce direct source-to-drain leakage current and reach a high on/off current ratio for ultra-scaled transistors [1]. As a member of 2D TMD material family, molybdenum disulfide (MoS2) has been theoretically demonstrated to be superior to Si with respect to the sub-5 nm scaling limit [2]. The first single-layer MoS2 transistor with a top-gated configuration was implemented in The device showed n-type transfer characteristics with an excellent on/off ratio of ~10 8, a room- 11

12 temperature field-effect mobility of > 200 cm 2 V -1 s -1, and SS of 74 mv/decade [3]. A near ideal SS of ~ 65 mv/decade has also been reached by MoS2 transistors using single-walled carbon nanotube as the gate electrode [1], which confirms the great potential of MoS2 for low-power electronic applications. WS2, another representative of semiconducting TMDs, has also been a focus as nextgeneration nanoelectronic and photonic materials. The large valence band splitting (~ 426 mev) originating from strong spin-orbit coupling enables WS2 to be a perfect platform for realization of spintronic and valleytronic devices. Theoretical simulations predict that the room-temperature phonon-limited electron mobility in monolayer WS2 is over 1000 cm 2 V -1 s -1, the highest value among the 2D semiconducting TMDs. Recently, back-gated single-layer WS2 FETs sandwiched between hexagonal boron nitride (hbn) films have been demonstrated, showing a high field-effect mobility (~ 200 cm 2 V -1 s -1 ) at room temperature with a high on/off ratio (~ 10 7 ) [4]. Despite the great progress in the electrical performance of TMD transistors as mentioned above, the TMD materials used in those high-performance devices were obtained by mechanical exfoliation methods. Such the mechanically exfoliated TMD flakes are typically few micrometers in size with a random number of layers, limiting their application in commercial viable devices. For practical applications, such TMD layered materials must be grown over large areas with good electrical and optical properties. The development of synthesis methods for obtaining large-scale, high-quality TMD materials is therefore of central importance. Chemical vapor deposition (CVD) has been one of the most practical methods for largearea growth of 2D materials because the precursors used for growth are human-friendly and the process is cost-effective. The highest reported mobility of CVD-grown single-layer MoS2 transistors, however, has been below ~ 30 cm 2 V -1 s -1 to date, which is substantially lower than the theoretically predicted value (~ 410 cm 2 V -1 s -1 at room temperature). Similarly, the field-effect mobility of CVD-grown WS2 devices ranges from 1 to 10 cm 2 V -1 s -1 [5, 6], showing significant discrepancies between experiment and theory. Although the carrier transport properties in TMD devices are strongly influenced by various extrinsic effects such as metal contacts, interface traps, charged impurities, and dielectric environment, it is essential to improve the intrinsic properties of 12

13 TMD materials from the synthesis side. Furthermore, the intrinsic defects and structural disorder in TMD layered materials could also be a fertile ground to explore their undiscovered properties and physical tunability. Thus, methods for large-scale growth of high-quality TMD monolayers, a deep understanding of TMD structural defects, and strategies to liberate their remarkable carrier transport properties are of vital importance to move the TMD field forward beyond Si technology. 1.2 Significance of This Work This thesis aims to develop reliable methods for growing high-quality monolayer MoS2 and WS2, explore their intrinsic electrical transport properties for electronic and circuit applications, and characterize the structural defects and homogeneity in the CVD-grown 2D TMD materials. This thesis is organized as follows: Chapter 2 offers an overview of 2D TMD materials, in particular, their intrinsic electronic and optical properties and structural defects. Chapter 3 illustrates a CVD synthesis method for large-area growth of high-quality monolayer MoS2 films and graphene/mos2 in-plane heterostructures. Various techniques are employed to characterize the CVD-grown monolayer MoS2, including Raman spectroscopy, atomic force microscopy, photoluminescence, and absorption spectra. Chapter 4 demonstrates a simple CVD method for monolayer WS2 growth and discusses the growth mechanism. X-ray photoelectron spectroscopy, Raman spectroscopy, atomic force microscopy, and photoluminescence are performed on the as-grown WS2 monolayer. Chapter 5 explores the electrical transfer characteristics of the single-layer MoS2 and WS2 triangular grains grown by the proposed CVD methods. Single-layer MoS2 and WS2 FETs with a back-gated configuration using nickel (Ni) as source and drain contacts are demonstrated. The enhancement of MoS2 electron mobility enabled by a high-k dielectric substrate is also presented. Chapter 6 provides a summary of our findings and discusses the future work for improving the performance of TMD transistors and for further study of 2D TMD defect characterizations and engineering. 13

14 Chapter 2. Physics of Transition Metal Dichalcogenides 2.1 Atomically Thin Structure Two-dimensional transition metal dichalcogenides (2D TMDs) which are only three-atom-thick are a class of materials with layered structures of the form X-M-X (MX2), where M is a transition metal element from group IV, V, and VI and X is a chalcogen (S, Se, and Te). The layered structure is constructed by a hexagonal plane of metal atoms which separates two hexagonal planes of chalcogen atoms with covalent bonds, as shown in Figure 2-1. (a) (b) Figure 2-1. (a) The lattice structure of monolayer MX2 and (b) the hexagonal planes of M and X atoms [3, 7]. These layered materials may exhibit semiconducting, metallic, or superconducting properties, depending on the selection of the transition metal elements. Figure 2-2 summarizes the electronic properties of some representative TMD materials. 2.2 Electronic Structure The electronic structures and properties of TMD materials vary with their thickness. For instance, semiconducting TMDs such as molybdenum disulfide (MoS2), tungsten disulfide (WS2), and tungsten diselenide (WSe2) exhibit a transition from direct bandgap to indirect bandgap as the thickness varies from single layer to multilayer. As an example, monolayer MoS2 has a direct bandgap of 1.8 ev, while bulk MoS2 possesses an indirect bandgap at 1.2 ev. The direct band gap results in a sharp photoluminescence emission at ~ 1.8 ev from monolayer MoS2. Figure 2-3 clearly shows that the absorption, photoluminescence (PL), and bandgap energy of MoS2 highly depend on its thickness. Single-layer MoS2 exhibits a much stronger PL emission than that of bilayer MoS2 and the MoS2 bandgap decreases as its thickness increases. 14

15 Figure 2-2. Summary of electronic properties of TMD materials [8]. Figure 2-3. (a) Comparison for PL spectra of monolayer and bilayer MoS2. (b) The dependence of MoS2 bandgap energy on its thickness. (c) Absorption spectra for monolayer and bilayer MoS2 [9]. The electronic band structures of monolayer, bilayer, and bulk MoS2 and WS2 calculated by first principles are presented in Figure 2-4. The influence of layer number on band structure is due to the quantum confinement effect and the change in orbital hybridization between d orbitals of Mo atoms and pz orbitals of S atoms. Unlike graphene that the lattice is all occupied by carbon atoms, the A and B sublattices of in MoS2 (WS2) lattice structure are occupied by Mo (W) atoms and a pair of S atoms (Figure 2-1 (b)). The difference between A and B sublattices results in the lift of the decency at K (K) points in the Brillouin zone and creates a desirable bandgap in MoS2 (WS2). 15

16 Figure 2-4. Band structures for bulk, bilayer, and monolayer MoS2 and WS2 [10]. 2.3 Electron Transport and Scattering In 2D TMD materials, transport and scattering of carries are confined in the plane of the materials. There are four mechanisms that affect the mobility of carriers [11]: (i) acoustic and optical phonon scattering, (ii) Coulomb scattering caused by charged impurities, (iii) surface interface phonon scattering, and (iv) roughness scattering Phonon Scattering Crystal deformation in 2D TMDs results in polarization fields that interact and scatter electrons. As temperature increases, carrier mobility is increasingly influenced by phonon scattering. Figure 2-5 (a) and (b) illustrate the dependence of MoS2 carrier mobility on temperature and carrier density. At low temperature (< 100 K), the acoustic phonon scattering dominates, while the optical phonon scattering dominates at higher temperature. The maximum mobility of MoS2 that can be 16

17 reached at room temperature is limited to ~ 410 cm 2 V -1 s -1, mainly due to the optical phonons. Similar values are expected for other monolayer semiconducting TMDs Coulomb Scattering Coulomb scattering in 2D TMDs originates from the charged impurities that randomly distribute within the layer or on the surface. The Coulomb scattering plays a dominant role for lowtemperature carrier transport properties, as shown in Figure 2-5 (c) [12]. In device engineering, the carrier concentration can be controlled through adding ionic impurities, while the mobility is also decreased due to the Coulomb scattering effect. The performance of devices is therefore significantly affected by the doping level of the materials. Theoretical calculations shows that the impurity scattering dominates over phonon scattering when the impurity concentration reaches ~ 5 x cm -2 [13], which is considered as heavily doping Surface Phonon Scattering and Roughness Scattering In metal-oxide-semiconductor field-effect transistors, 2D TMD materials are placed on a dielectric material such as SiO2. The surface roughness of SiO2 would cause ripples or wrinkles on the 2D TMD materials, which may also contribute to scattering and thus reduce carrier mobility. The effects of interface phonon scattering and roughness scattering have been experimentally observed in GaAs-based quantum wells [14], graphene placed on SiO2, and freely suspended graphene [15, 16]. 17

18 (a) (b) (c) Figure 2-5. The dependence of carrier mobility in monolayer MoS2 on (a) temperature and (b) carrier density (calculated from first-principles density functional theory). (c) The scattering contributions from charged impurities and the total mobility of MoS2 due to the combined effect [12, 13]. 2.4 Structural Defect-dependent Properties Defects in TMD layered materials can be classified as zero-dimensional, one-dimensional, and two dimensional defects, as shown in Figure 2-6. Zero-dimensional defects are the most abundant defects in TMDs, including point defects, dopants, or non-hexagonal rings. One-dimensional defects contain grain boundaries, edges, and in-plane heterostructures. Layer stacking of different TMDs, wrinkling, folding, and scrolling are assigned to two dimensional defects. Structural defects in the crystal lattices of TMDs can significantly change their physical and chemical properties. For example, sulfur vacancies, the most common defects in chemically synthetic and mechanically exfoliated MoS2 monolayers due to the lowest formation energy of these defects, introduce unpaired electrons into the lattice, resulting in a n-doping effect on the material. These sulfur vacancies create additional density of states within the band gap (Figure 2-7), and further alter the electrical transport properties of MoS2. Most MoS2 devices show a n-type transfer characteristic due to the plenty of sulfur vacancies in the material. A p-type transfer behavior has also been reported from sulfur-rich (or molybdenum-deficient) MoS2 and Nb-doped MoS2. In addition to electronic properties, optical properties of TMDs are also strongly affected by structural defects. Tunable photoluminescence emissions can be achieved by doped TMDs such as MoxW1- xs2 and MoSxSe2-x [17], which is promising for LED and display applications. Bi-sulfur vacancies 18

19 generated by plasma irradiation in exfoliated MoS2 monolayer induce a PL peak with an energy lower than the band gap value [18]. As 1D defects in TMDs, visible light emissions from the edges of CVD-grown WS2 single-crystalline domain show similar or higher intensities compared to the interior regions [19, 20]. Vertical and in-plane heterostructures of TMD materials also enable new excitonic transitions. For example, mechanically stacked MoS2/WSe2 hetero-bilayers show a strong PL emission at 1.50 ~ 1.56 ev (Figure 2-8 (a) and (b)) [21], originating from strong interlayer coupling of charge carriers between two single-layer TMDs. Moreover, in parallel stitched heterostructures of MoS2/WS2 and MoSe2/WSe2, their in-plane interfaces can emit visible light with an energy laying between the band gap values of the two materials and the emission intensity from the in-plane junctions is stronger than those from both sides [22, 23], as shown in Figure 2-8 (c) and (d). Figure 2-6. Typical defects in 2D TMD materials [24]. 19

20 (a) (b) Figure 2-7. (a) Formation energies of various point defects as functions of sulfur chemical potential. Sulfur vacancy (Vs) has the lowest formation energy (~ 2 ev). (b) Schematic depiction of the defect levels within MoS2 band gap [25]. (a) (b) (c) (d) Figure 2-8. (a) Optical microscope image of a WSe2/MoS2 hetero-bilayer. (b) Normalized PL and absorbance spectra for monolayer MoS2, WSe2, and their hetero-bilayers. (c) SEM images of MoSe2/WSe2 heterostructures. (d) PL intensity map of a MoSe2/WSe2 lateral heterostructure (Scale bars, 2 μm) [21-23]. 20

21 Chapter 3. CVD Growth of Monolayer MoS2 In this chapter, a synthesis method based on chemical vapor deposition (CVD) for growing highquality continuous films of monolayer MoS2 is demonstrated. Various optical and surface characterizations including photoluminescence (PL), optical absorption, Raman spectroscopy, PL and Raman intensity mappings, and atomic force microscope (AFM) are employed to characterize the quality and homogeneity of the CVD-grown MoS2 films and the MoS2 single-crystal domains. This proposed method offers a high reliability for growth of large-area high-quality MoS2 and can be applied to various substrates such as mica, fused silica, and sapphire. 3.1 Monolayer MoS2 Growth In this work, the monolayer MoS2 films were synthesized by chemical vapor deposition (CVD) method. Here, 300 nm SiO2/Si wafers were used as substrates for MoS2 growth. Perylene-3,4,9,10- tetracarboxylic acid tetrapotassium salt (PTAS) as shown in Figure 3-1 (a) and (b) were used as seeding molecules. Figure 3-2 (a) schematically illustrates our experimental setup. At the center of the CVD furnace, three wafers were faced down and placed directly above a crucible containing 15 mg of molybdenum oxide (MoO3) precursor. The upstream and downstream wafers were coated with PTSA molecules, while the central substrate is a cleaned, bare SiO2/Si wafer for MoS2 film growth. Another crucible that contains 20 mg of sulfur (S) powder was put at the inlet of the CVD furnace. 15 sccm Argon (Ar) was introduced into the CVD system as a carrier gas. The system was then heated to 625 C at a rate of 30 C min 1, and the MoS2 monolayer films were grown on the central substrate at 625 C for 3 min under atmospheric pressure. Finally, the system was naturally cooled down to room temperature. Figure 3-2 (b) schematically shows the temperature profile of the growth process. 21

22 (a) (b) Figure 3-1. (a) Molecular structure of PTAS and (b) PTAS aqueous solution. (a) (b) Figure 3-2. (a) A schematically illustration of the CVD setup for growth of MoS2 films and (b) the temperature profile used for monolayer MoS2 growth. 3.2 Characterizations Optical Microscopy Characterizations Three-atom-thick semiconducting MoS2 films deposited on SiO2/Si wafers are blue-green in color which exhibit apparent contrast with the purple of the substrates as shown in Figure 3-3 (a). The continuous area of the MoS2 thin films grown in this work is typically ~ 1 cm x 1cm (limited by the dimension of the 1-inch quartz tube of our CVD system). As shown in Figure 3-3 (b), triangular-shape domain of MoS2 single crystals can be found at the edges of the continuous region. The largest size of the single triangular grains is ~ 50 μm. Figure 3-3 (c) shows optical microscopy images of the typical as-grown monolayer MoS2 film. The film is completely continuous with high uniformity. With optimized process conditions, regions of multilayer MoS2 or particle clusters are 22

23 not observed in our MoS2 films. The transition from single-layer grains to continuous films is presented in Figure 3-3 (d). (a) (b) (c) (d) Figure 3-3. (a) Photograph of centimeter-scale monolayer MoS2 grown on SiO2/Si wafer. (b) Optical images of the CVD-grown single-crystal triangular domain of MoS2, (c) the completely continuous film of MoS2, and (d) the transition region between MoS2 triangular flakes and continuous film Raman Spectroscopy Characterizations To confirm the layer thickness of the as-grown MoS2 thin films and flakes, Raman spectroscopy is employed. As shown in Figure 3-4 (a), the CVD-grown MoS2 triangular grains exhibit two characteristic Raman modes: the out-of-plane vibration of S atoms (A1g) at ~ cm -1 and the in-plane vibration of S and Mo atoms (E 1 2g) at ~ 385 cm -1. The frequency difference of Raman shifts between A1g and E 1 2g modes ( ) can be related to the thickness of MoS2. Our deposited MoS2 triangular domain shows a value of ~ 20.5 cm -1, which evidences the existence of single-layer MoS2. The ratio between intensities of the A1g and E 1 2g modes can be linked to the doping levels 23

24 in MoS2 [26]. The results indicate that our CVD-grown MoS2 can be analogous to the case of less doping. Additionally, since the widths of the Raman peaks are narrow, it can be inferred that there does not exist high structural disorders in our MoS2 samples [27]. Figure 3-4 (b) displays the Raman spectra performed at different locations in the continuous regions of the deposited MoS2 films. All of the regions measured show a similar Raman profile with the monolayer MoS2 characteristic of ~ 20 cm -1. No apparent variations of Raman peak intensity and are observed. One can notice that both the positions of A1g and E 1 2g modes at point C are slightly shifted toward lower frequencies while maintains a value of ~ 20 cm -1. Since the Raman frequencies are strainsensitive, one can infer that the shifts at point C may originate from the strain of grain boundaries near point C. A typical image of the grain boundaries in the as-grown MoS2 films is shown in the following section Overall, the Raman spectroscopy analysis indicates that the as-grown films of monolayer MoS2 possess highly spatial uniformities. (a) (b) Figure 3-4. Raman spectra of as-grown monolayer MoS2 (a) triangular flakes and (b) films measured at different regions. 24

25 3.2.3 Atomic Force Microscopy Characterizations Figure 3-5 shows the AFM images of the grown MoS2 films. Grain boundaries can be clearly seen in Figure 3-5 (a), indicating our MoS2 films are formed from merging different grain domains. The average grain size is few micrometers large. Compared with the MoS2 films grown by MOCVD [28], the grain size of our CVD-grown MoS2 films is much larger. A thin-film material with a larger grain size is typically expected to have better performance on its intrinsic electrical transport properties, which is critical to device and circuit applications. The electrical properties of the CVDgrown MoS2 are evaluated in chapter 5. Figure 3-5 (b) and (c) illustrate a thickness of 0.75 nm for a typical as-grown MoS2 triangular domain, which corresponds to single-layer thickness plus a van der Waals gap. (a) (b) (c) Figure 3-5. AFM images of monolayer MoS2 (a) films and (b) triangular grains. (c) AFM step height profile of a typical single-layer region. 25

26 3.3 Photoluminescence and Optical Absorption of MoS Photoluminescence In addition to Raman spectroscopy, the photoluminescence (PL) of MoS2 exhibits a high dependence on its thickness. Due the 2D confinement, monolayer MoS2 possesses a direct band gap of at least 1.8 ev and allows strong PL emission, while multilayer or bulk MoS2 has an indirect band gap smaller than 1.8 ev and the emission intensity is thus much weaker. Figure 3-6 shows the photoluminescence spectrum acquired at room temperature of our CVD-grown MoS2 films. The MoS2 films exhibit a band gap at ~1.84 ev, which further confirms that the as-grown MoS2 is single-layer. The PL emission profile of MoS2 depends on the relative contributions of A and B direct excitonic transitions. In our MoS2 films, the intense A excitonic emission at ~1.84 ev (674 nm) can be clearly observed, which is assigned to the direct transition from the conduction band minimum to the uppermost valence band maximum at the K valley in the Brillouin zone. The B exciton, corresponding to the direct transition from the conduction band minimum to the lower valence band maximum at the K valley, however, is not observed. That suggests most of the photogenerated holes relax to the upper valence band before recombination. It has been reported that the B excitonic peak intensity is tunable at higher excitation densities [29]. At a high excitation density, the occupation of hole at the upper valence band increases, so the available states for the photogenerated hole relaxation from lower to upper valence band reduce, which is called statefilling effects. Besides the direct band gap characteristic, since the intensity of PL emission is also significantly associated with the defects in the materials and interface traps between the deposited materials and substrates, the PL intensity in principle can be an indicator of the quality of the asgrown materials. It is worth mentioning that the ratio of PL to Raman peak intensities (IPL/IA1g) of our MoS2 is ~ 100, suggesting an excellent optical quality. The full-width-half-maximum (FWHM) of the PL emission peak is as small as ~ 55 mev (~ 20 nm), which is comparable to or even smaller than that of monolayer MoS2 grown on sapphire substrates [30, 31]. The results evidence that highquality CVD-grown MoS2 films are deposited on SiO2/Si substrates via the proposed growth method. 26

27 Figure 3-6. Room temperature photoluminescence of the as-grown monolayer MoS2 films. Raman spectrum is also displayed in this figure in the unit of photon energy. An excitation wavelength of 532 nm was used for all PL and Raman spectra in this thesis UV-visible Absorption To investigate the pristine behavior of MoS2 optical absorption, we directly grown MoS2 films on various transparent substrates, including mica, sapphire, and fused silica. Thanks to the high transparency of these substrates, we can easily perform optical absorption on those CVD-grown samples without PMMA transfer processes. Therefore, the potential contaminations such as PMMA and KOH residuals can be avoided. The absorption spectrum of the MoS2 grown on mica (Figure 3-7) shows three pronounced peaks in the wavelength range from 350 to 750 nm. MoS2 films deposited on fused silica and sapphire also display the same absorption characteristic. The A and B excitonic absorptions at 670 nm (~ 1.85 ev) and 620 nm (~ 2.00 ev) originate from transitions from the highest-laying spin-split valence bands to lowest conduction bands. The C absorption peak at ~ 445 nm is ascribed to van Hove singularity of monolayer MoS2. The existence of van Hove singularity indicates strong light-matter interactions in MoS2 and could be used to enhance the performance of photonic and photovoltaic devices. 27

28 (a) C (b) C B A B A (c) C B A Figure 3-7. Optical absorption spectra of as-grown monolayer MoS2 films on (a) mica, (b) fused silica, and (c) c-plane sapphire. 28

29 3.4 Defect Characterization of MoS PL and Raman Intensity Mappings of MoS 2 Triangular Domains To more qualitatively characterize the uniformity of our CVD-grown MoS2, the intensity mappings of PL emission wavelength and Raman mode frequencies (A1g and E2g) are performed, as shown in Figure 3-8. The as-grown MoS2 grain does not show obvious local PL variations. The PL quenches associated with micro-scale defects such as adlayers, cracks and other physical damages are not observed in the MoS2 single-crystal domain, indicating a high spatial homogeneity of the MoS2 triangular domain PL Mappings of MoS 2 Grain Boundaries Figure 3-9 (a) shows the PL mapping of the CVD-grown MoS2 film with a focus on a region containing linear-shape grain boundary (highlighted by a blue dotted square in Figure 3-9 (b)). An enhancement of PL emission is observed at the grain boundary (the yellow region in Figure 3-9 (a)). The enhanced PL emission could be caused by the inhomogeneous strain at the boundary which alters the electronic structure of the local MoS2 and leads to an increased exciton binding energy. Additionally, higher PL intensities are also observed at some locations rather than the grain boundaries in the MoS2 film. The enhanced PL at those points could be interpreted as carrier inhomogeneity at those locations. Lattice defects such as vacancies or impurities could contribute to extra carriers that can recombine with the photogenerated carriers, resulting in stronger PL emissions at those locations. To further confirm the charge inhomogeneity in the CVD-grown MoS2, we also tried to perform conductive atomic force microscopy (C-AFM) on our CVD MoS2 samples. However, since the CVD-grown MoS2 generally shows an insulating behavior as no gate voltage is applied, it has been difficult to attain meaningful signals from that measurement. More efforts will be made to characterize the lattice defects or charge inhomogeneity in the CVD-grown TMDs. Other strategies such as introducing extrinsic doping and employing scanning gate microscopy (SGM) will be considered for our future study of TMD structural defect characterizations. 29

30 (a) (b) (c) MoS 2 PL MoS 2 E 2g MoS 2 A 1g Figure 3-8. Intensity mappings of (a) PL emission wavelength and (b) Raman E2g and (c) A1g mode frequencies of a representative as-grown MoS2 single crystal. (a) (b) Figure 3-9. (a) PL Intensity mapping for a region containing a linear-shape grain boundary and (b) optical image of as-grown monolayer MoS2 film. The region corresponding to (a) are highlighted by a blue dotted box. 30

31 3.5 Graphene/MoS2 In-plane Heterostructures Heterostructures formed by 2D materials embrace rich physics and offer opportunities to create new device architectures with multifunctionality and high performance. For example, hybrid structures based on semiconducting monolayer TMDs can serve as building blocks for flexible p- n junction devices, high-speed electronics, and optoelectronics. Vertical van der Waals herterostructures can be achieved by layer-by-layer stacking of 2D materials using mechanical transfer. However, van der Waals herterostructures in horizontal direction (in-plane junction) can only be formed by growth. The growth of atomically clean and sharp interfaces of in-plane herterostructures remains challenging. In this section, a two-step growth method for large-area growth of metal-semiconductor graphene/mos2 herterostructures is demonstrated. The lateral graphene/mos2 heterojunction is schematically shown in Figure Graphene MoS2 Figure Schematic illustration of the in-plane heterostructure of Graphene/MoS2 [32] Lateral Heterostructure Growth First, SiO2/Si substrates are cleaned by piranha solution (a 3:1 mixture of concentrated sulfuric acid with hydrogen peroxide) for 1 h. Next, the mechanically exfoliated graphene flakes are transferred onto the piranha-treated SiO2/Si substrate. After that, the graphene/sio2/si samples are annealed in a H2 (100 sccm)/ar (300 sccm) environment at 350 C for 3 h for removing the residue of scotch tape. The graphene/sio2/si samples then serves as the growth substrate for CVD MoS2 deposition. The MoS2 growth process follows the same steps described in section 3.1. The growth patterns of MoS2 are dominated by the hydrophile and hydrophobe of the local areas on the growth substrates. The mechanism of the PTAS molecule diffusion during the growth allows MoS2 to 31

32 grow only on the hydrophilic SiO2 surface, while MoS2 deposition on the hydrophobic graphene surface is significantly suppressed. As a result, a heterostructure with graphene wrapped by monolayer MoS2 can be obtained Optical and AFM Images The typical configuration of the CVD-grown lateral graphene/mos2 heterostructures is shown in Figure 3-11 (a)-(c). The in-plane interfaces between graphene and MoS2 are distinguishable by the contrast difference. The graphene flakes are encircled by CVD-grown single-layer MoS2 (identified by PL and Raman spectra as shown in Figure 3-11 (e) and (f)). The AFM image (Figure 3-11 (d)) shows that a continuous MoS2 film is grown on the hydrophilic SiO2 region and are well stitched with the graphene flake. The boundaries between MoS2 and graphene are sharp with an overlap height ~ 2 nm, as shown in the AFM step profile. Moreover, there are no MoS2 clusters or adlayers observed on the graphene surface, revealing the selective growth nature of MoS2. 32

33 (a) (b) (c) Monolayer MoS2 Monolayer MoS2 Film Graphene Monolayer MoS2 Film Graphene Graphene (d) Graphene MoS2 (e) (f) Figure (a)-(c) Typical optical images of as-grown in-plane heterostructure of Graphene/MoS2. (d) The AFM image of the region highlighted by yellow circle in (a). (d) PL and (e) Raman spectra of the grown MoS2 monolayer wrapping graphene PL and Raman Intensity Mappings The spatial homogeneity of the graphene/cvd-grown MoS2 in-plane heterosturctures can be further confirmed by the mappings of the PL and Raman intensities. Figure 3-12 shows that there is no MoS2 PL signal on the graphene surface, which evidences that the monolayer MoS2 is grown only outside the graphene flakes. The mapping images suggest that sharp and well-stitched 33

34 boundaries are formed at the junctions of MoS2 and graphene, with no cracks or tears. Interestingly, the PL emission at the MoS2/graphene junction exhibits much a stronger intensity than that at other MoS2 areas, indicating an inhomogeneous carrier distribution or a higher photogenerated carrier recombination rate at the edges. The unique enhancement of light-matter interactions at the inplane junction of MoS2/graphene heterostructures could play an important role in designing novel solar cells, light-emitting diodes, and broad-spectrum photosensors. (a) Monolayer MoS2 (b) Graphene G band (c) Graphene G band Graphene flake (d) MoS 2 PL (e) MoS 2 E 2g (f) MoS 2 A 1g Figure PL and Raman intensity mappings of the graphene/mos2 junction corresponding to the region highlighted by the yellow circle. 34

35 Chapter 4. CVD Growth of Monolayer WS2 4.1 Monolayer WS2 Growth In this section, a simple CVD growth method is demonstrated. The synthesis of monolayer WS2 is performed in a quartz tube furnace at atmospheric pressure. Figure 4-1 schematically illustrates the CVD experimental setup. At the center of the furnace, ~10 mg of tungsten trioxide (WO3) powder is directly sprayed onto a piece of bare SiO2/Si wafer (1.6 cm x 1.6 cm), which serves as a plate for carrying the WO3 precursor. A pre-cleaned Si/SiO2 (300 nm) substrate for WS2 growth is positioned face-up 1 cm away the upstream wafer containing WO3. A crucible containing Sulfur (S) powder is placed upstream, 1 cm away from the edge of the furnace heating zone. The optimized distance between the S crucible and the WO3-contained wafer is ~ 18 cm. Prior to synthesis, 1000 sccm of Argon (Ar) is employed to purge the quartz tube for 5 min. To grow monolayer WS2, the furnace temperature is quickly ramped to 800 C at a rate of 39 C/min and then hold the temperature at 800 C for 5 min. A continuous 50 sccm Ar flow is used during the growth process. Finally, the furnace temperature is naturally cooled down to room temperature. (a) (b) Figure 4-1. Schematic illustrations of (a) the experimental setup and (b) the temperature profile used for monolayer WS2 growth. 35

36 4.2 WS2 Growth Evolutions In our experimental setup, the optimized condition for monolayer WS2 growth is at 800 C for 5 min as described above. To further reveal the evolution of WS2 growth with growth time and temperature, we compare the WS2 grains grown by various recipes, namely, recipe A: 800 C for 5 min, recipe B: 775 C for 15 min, and recipe C: 800 C for 15 min. First, it is instructive to investigate the effect of temperature. By comparing the optical images of WS2 grown by recipe A and B, one can clearly see that a lower growth temperature generally results in a smaller grain domain. The average size of WS2 single domain grown at 800 C using recipe A is ~ 40 m and the largest size found is ~ 65 m, as shown in Figure 4-2 (a). On the other hand, recipe B (775 C) shows a relatively small domain size ~ 5 m. Furthermore, a longer growth time (15 min for both recipe C and B) is favorable for multilayer WS2 growth. Compared to recipe A, both recipe B and C lead to thicker WS2 grains. Note that the domain size grown by recipe C is larger than that of recipe B, which again suggests that higher temperatures would be beneficial to large-domain WS2 growth. Next, we investigate the effect of growth time. A recipe D using the same temperature of 800 C instead for 1 min is used for making a comparison with recipe A. The recipe D leads to monolayer WS2 but with a small average grain size of ~ 3 m, revealing the existence of monolayer WS2 deposited on SiO2/Si substrates at the very beginning of growth process. Based on the results, we conclude the evolution of WS2 growth: at the beginning of 800 C, nucleation sites (probably WO3) are first absorbed onto the substrates. The nucleation sites continue to grow and then become monolayer triangular domain within 1 min. After that, the lateral growth mechanism dominates in the following 5 min, meaning that the small WS2 monolayers start to grow laterally and eventually enlarge to ~ 40 m. After the 5th minute of the growth process at 800 C, the vertical growth mechanism starts to take over the WS2 growth, meaning that the rate of vertical growth become higher than that of lateral growth. As a result, the WS2 grains grown at either 775 C or 800 C for 15 min shows a nature of multilayer. Figure 4-2 (e) and (f) shows the AFM image for the WS2 tringle grown at 800 C, where the thickness of ~ 1 nm is consistent with the monolayer WS2 thickness plus a van der Waals gap. The sharp height profile at the edge of the WS2 tringle indicates that there are active sites at the WS2 edges where the precursors are absorbed and then undergo chemical reactions. The active sites at the edges are responsible for the lateral enlargement of WS2 grains. There are also particles absorbed on the WS2 surface, which may suggest that the adlayer 36

37 or multilayer WS2 start to grow around the moment of the 5th minute. Figure 4-3 schematically summaries the growth dynamics of WS2 grown at 800 C. (a) Recipe A: 800 C for 5min The Largest WS 2 grain grown by recipe A (b) Recipe B: 775 C for 15min (c) Recipe C: 800 C for 15min (d) Recipe D: 800 C for 1min (e) (f) Figure 4-2. (a)-(d) Optical images of WS2 grains grown by different recipes. The WS2 monolayer grown by recipe A is the focus of this chapter. (e) The AFM image of WS2 monolayer grown by recipe A and (f) the thickness profile along the yellow line in the AFM image. 37

38 Figure 4-3. A schematic illustration of the CVD-grown WS2 evolution with growth time at 800 C. 4.3 Raman Spectroscopy Characterizations Raman spectroscopy using an excitation wavelength of 532 nm is employed to characterize the thickness of the CVD-grown WS2. As shown in Figure 4-4, two Raman peaks located at ~ 350 and ~ 417 cm -1 are observed. The 350 cm -1 peak is attributed to the 2LA (longitudinal acoustic) mode merged with the E 1 2g modes. The LA phonon vibrational mode, as a function of crystalline disorder, arises from in-plane collective movements of atoms in the lattice, while the E 1 2g is optical mode and originates from the in-plane vibration of S and W atoms. On the other hand, the 417 cm -1 Raman peak is the out-of-plane vibration A1g characteristic of WS2. It has been reported that not only the frequency difference ( ) of E 1 2g and A1g peaks, but also the peak intensity ratio of 2LA to A1g of WS2 is highly sensitive to its thickness. For single-layer WS2 grown on SiO2 at an excitation wavelength of 514 nm, the height of the 2LA peak is roughly 2 times that of the A1g peak (I2LA/IA 1g ~ 1 for bilayer and smaller than 1 for three or more layers). Our-grown WS2 shows I2LA/IA 1g ~ 2.5 with a smaller than 67 cm -1 under 532 nm excitation, which evidences that our CVD-grown material is monolayer WS2. 38

39 Figure 4-4. Room-temperature Raman spectrum of as-grown WS2 flakes using a 532 nm laser excitation. 4.4 Photoluminescence of WS2 For bulk WS2, there are two direct transitions at the K point in the Brillouin zones due to the splitting of the valence band. These two transitions are assigned to A (1.95 ev) and B (2.36 ev) excitons, respectively, and have been experimentally detected by absorption spectroscopy. On the other hand, this splitting of the valence band for a monolayer WS2 is absent, which means only one direct electronic transition is expected to be observed from optical spectroscopy. Figure 4-5 represents the PL spectrum of as-grown WS2 samples (with 532 nm laser excitation). The WS2 samples exhibits a single strong PL emission at ~ 1.97 ev (~ 629 nm), which is consistent with the direct band gap property of monolayer WS2. Also, PL FWHM can be an indicator of sample quality. A smaller FWHM in principle suggests a higher quality. Our CVD-grown WS2 shows a PL FWHM of ~ 51 mev, which is comparable or even narrower compared to those previously reported CVDgrown and exfoliated monolayer WS2 on SiO2/Si substrates. This result demonstrates a high optical quality in our CVD-grown monolayer WS2. The electrical transport quality of as-grown WS2 is performed in chapter 5. 39

40 Figure 4-5. Room-temperature photoluminescence spectrum of as-grown monolayer WS2 (under 532 nm excitation). 4.5 X-ray photoelectron spectroscopy of WS2 To characterize the chemical composition and further estimate the potential structural defects of our CVD-grown WS2, X-ray photoelectron spectroscopy (XPS) is performed. Figure 4-6 reveals the XPS spectra of the as-grown WS2 samples. Three characteristic XPS peaks of WS2 at binding energies 33.0 ev, 35.2 ev, and 37.8 ev corresponding to W4f7/2, W4f5/2, and W5p3/2 core energy levels, respectively, are observed for tungsten (W) atom. The W4f7/2, which represents the 4+ valence state, shows a dominant contribution and it indicates the WO3 (6+) precursor is sufficiently sulfurized even without employing H2 in our experimental setup. A S2p doublet is also observed, confirming the grown materials are WS2. The stoichiometry of the as-grown WS2 monolayer can be calculated by [W] [S] = λ S2p λ W4f σ S2p(hυ) σ W4f (hυ) I W4f I S2p (4-1) where σs2p(hν) and σw4f(hν) are photo-ionization cross sections of the 2p and 4f core level of S and W, respectively, and λs2p and λw4f are inelastic mean free paths of the photoelectrons with 40

41 kinetic energies corresponding to the S and W core levels, respectively. The values of these abovementioned parameters can be obtained from literatures. Accordingly, the [W]/[S] ratio is estimated to be ~ 0.6, suggesting ~ 20% sulfur vacancies in the CVD-grown monolayer of WS2. Since the existence of sulfur vacancies, we expect that our grown WS2 monolayer is a n-type semiconducting material. The n-type behavior of as-grown monolayer WS2 is demonstrated by field-effect devices in chapter 5. Figure 4-6. Chemical composition analysis of as-grown WS2 using X-ray Photoelectron Spectroscopy 41

42 Chapter 5. Monolayer MoS2 and WS2 Field-effect Transistors For field-effect transistors (FETs) based on TMDs such as MoS2 and WS2, the electrical contacts can significantly influence the device performance. In general, the electrical performance of TMD devices is limited by Schottky barriers at the metal/tmd interface. Therefore, the realization of ohmic contacts on TMD materials plays an important role for improving the performance of TMD devices. In this chapter, CVD-grown MoS2 and WS2 FETs with Nickel (Ni) contacts are explored. Their electrical transport characteristics are discussed. 5.1 MoS2 Device Fabrication TMD Transfer The CVD-grown monolayer MoS2 grown on a SiO2/Si substrate is transferred using a KOH wet transfer method to SiO2/Si(p ++ ) substrates, which also serve as back gates for field-effect transistors. Poly-methylmethacrylate (PMMA, 950k 4.5% dissolved in Anisole) is spin-coated onto the CVD-grown monolayer MoS2 samples. The stack is then placed in a KOH solution and the solution is heated up to 85 C. The PMMA/MoS2 stack is able to be separated from the substrate and remains floating once the SiO2 layer is etched away. The PMMA/MoS2 film is then transferred into distilled water using a glass slide for 20 min to remove the KOH residues. It is worth mentioning that because the KOH residues would damage the surface of the SiO2/Si(p ++ ) substrates, leading to current leakages and the potassium ion (K + ) left between SiO2 and MoS2 may cause the threshold voltage shifts and significant hysteresis of electrical transport, the above-mentioned distilled water rinse step is repeated for at least three times. After that, the distilled water rinsed PMMA/MoS2 film is transferred onto the SiO2/Si(p ++ ) substrates and is then baked at 80 C for 10 min and 130 C for another 10 min. The bake steps can remove moisture and improve the adhesion between MoS2 and the substrates. Finally, the MoS2/SiO2/Si(p ++ ) sample is immersed in acetone for 1 min and then the sample is annealed at 350 C for 3 h in a mixture of argon (300 sccm) and hydrogen (100 sccm) environment. 42

43 5.1.2 E-beam Lithography, Electrode Deposition, and Lift Off In this section, the detailed process flow for fabrication of bottom-gate monolayer MoS2 fieldeffect transistors is described. Overall process flow is depicted in Figure 5-1. Figure 5-1. Schematic process flow for fabrication of back-gate single-layer MoS2 field-effect devices. The WS2 devices are fabricated using the same process flow. 5.2 Electrical Transport Properties of MoS2 FETs with Ni contacts A schematic depiction and optical microscope image of a representative single-layer MoS2 fieldeffect transistor (FET) are shown in Figure nm pure Nickel (Ni) contacts for source and drain were deposited directly on MoS2 by e-beam evaporation. All electrical measurements were carried out in a vacuum probe station (~ 2.5 x 10-4 torr). 43

44 Figure 5-2. Schematic and optical image of the single-layer back-gate MoS2 EFT. Symmetrical 50 nm Ni contacts are defined on the CVD-grown monolayer MoS2 by e-beam lithography, evaporation, and lift-off. Figure 5-3 shows the typical current (Ids) and back-gate voltage (Vbg) measurements at a drain-tosource bias Vds = 0.5 V for the single-layer MoS2 FETs. The MoS2 transistor we fabricated behaviors as a n-type channel device with a narrow hysteresis window at room temperature. A large hysteresis gap of MoS2 FETs has been often observed regardless of the device structures and the number of layers. The larger hysteresis represents the more instability of the devices. The origin of the hysteresis behavior could be associated with interface impurities between MoS2 channel and the oxide layer, the gaseous and/or water molecules absorbed from the environment, and the intrinsic defects in MoS2. Our MoS2 FETs exhibit small hysteresis behavior in vacuum, indicating an intrinsically high-quality MoS2 monolayer is grown. The threshold voltage (VT) of the device can be obtained by the linear extrapolation method, suggesting the VT is ~ -7.5 V. Figure 5-4 shows the Ids-Vds curves at different Vbg values. The linear dependence of the channel current on drainsource bias indicates that the Ni contacts are ohmic. Based on the assumption of linear charge dependence on the gate voltage overdrive, the carrier density in MoS2 channel surface is estimated by n C ox q (V gs V T ) (5-1) where Vgs = Vbg (Vs = 0 V) and Cox is the oxide capacitance of 1.15 x 10-8 F/cm -2 in our case. The carrier densities of MoS2 channel are in the order of ~ cm -2 at various back gate bias from 0 44

45 to 30 V. The field-effect mobility can be extracted from the data as shown in Figure 5.3 using the expression = g m 1 L V ds WC ox (5-2) where L and W are the channel and the width of the device and gm (= dids/dvgs) is the transconductance. The peak room-temperature field-effect mobility of the MoS2 FETs with Ni contacts is ~ 5.2 cm 2 V -1 S -1. Figure 5-3. Room temperature transfer characteristic for the single-layer MoS2 FET with Vds = 500 mv (L ~ 1 m and W ~ 10 m). Back-gate voltage is applied to the substrate. 45

46 Figure 5-4. Ids-Vds characteristic for the MoS2 device acquired for different values of Vbg. Low-temperature of 77 K transport characteristic of the monolayer MoS2 FETs was also performed, as shown in Figure 5.5. A n-type characteristic with Ion/Ioff of at least 10 5 is obtained. The threshold voltage VT is shifted to ~17.5 V at 77 K. Figure 5-5. Low-temperature Ids-Vbg transfer curve of the MoS2 FET acquired at 77 K. 46

47 5.3 Enhancement of Electron Mobility in MoS2 via a High-k Gate Dielectric In this section, we explore the effect of high-k gate dielectrics on the field-effect mobility of MoS2. A high-k gate dielectric Hafnium(IV) oxide (HfO2) of 25 nm is introduced between MoS2 and SiO2 by electron beam physical vapor deposition. The structure of the resultant MoS2 devices is shown in Figure 5.6. The dielectric constant of HfO2 is ~ 25 with a bandgap of 5.8 ev [33], which is much higher than that of SiO2 (εsio2 ~ 3.9). As a result, the gate capacitance Cg of the hybrid HfO2/SiO2 dielectric for the MoS2 FETs is ~ 11.3 nf/cm 2, calculated by considering two capacitances (CSiO2 and CHfO2) are connected in series. Figure 5.7 represents the typical electrical transport properties (Ids-Vbg) of the monolayer MoS2 transistors with the hybrid gate dielectric. The MoS2 devices with HfO2/SiO2 hybrid dielectric show n-type conduction and exhibit one order of magnitude higher current density compared to the MoS2 devices using bare SiO2 dielectric at the same gate biases (~ 0.9 µa/µm for the MoS2/SiO2 device and ~ 7.2 µa/µm for the MoS2/HfO2/SiO2 device at Vbg = 30 V). Since the gate capacitance of the hybrid HfO2/SiO2 substrate is only ~ 1.2% smaller than the bare SiO2 substrate, a similar value of carrier density (n = CgVbg) is expected when those two devices are biased at the same gate voltage. According to the definition of current density (J = nqµe), the increased current density indicates an enhanced carrier mobility in the HfO2/SiO2 dielectric environment. We extract the field-effect mobility from the data presented in Figure 5-7. The HfO2-based MoS2 FETs show an improved carrier mobility reaching ~ 18.2 cm 2 V -1 S -1 at room temperature. The reasons for the enhanced field-effect mobility could be linked to the reduction of Coulomb scattering due to the strong screening effect of high k dielectric. Output characteristics (Ids-Vds) at various back-gate voltages ranging from -20 V to 20 V for the HfO2- based MoS2 devices are presented in Figure 5-8. The linear Ids-Vds characteristics indicates the contacts between Ni and MoS2 are ohmic. Theoretical study has revealed that the Ni/MoS2 interface exhibits a lower vertical Schottky barrier height for electrons, compared to the Au/MoS2 and Pt/MoS2 interfaces [34]. 47

48 Figure 5-6. The Schematic of a back-gated MoS2 FET using a 25 nm HfO2/ 300 nm SiO2 hybrid gate dielectric. Figure 5-7. Transport properties of single-layer MoS2 FETs (L ~ 1 m and W ~ 6 m) on a HfO2/SiO2 substrate. 48

49 Figure 5-8. Output characteristics (Ids-Vds) of the single-layer MoS2 FETs at different backgate voltages. 5.4 Electrical Transport Properties of WS2 FETs with Ni contacts Figure 5-9 displays a schematic depiction and the top view optical microscopy image of a representative single-layer WS2 channel contacting 50 nm thick e-beam evaporated Ni electrodes as source (S) and drain (G) on a SiO2(300 nm)/si(p + ) substrate, which serves as a back gate (G). All the electrical characteristics of the devices were investigated under vacuum (~ 3.6 x 10-4 torr). Output characteristics of Ids-Vds at various back-gate voltages (-35 V to 0 V) for the single-layer WS2 FET are depicted in Figure The linear dependence of the current on source-drain biases at small fields indicates that the contact between Ni and WS2 is ohmic. The symmetry of the current with respect to the origin at positive and negative bias voltages (Figure 5-10(b)) further evidences the ohmic nature of the contacts. As is typical of long channel MOSFETs, the single-layer WS2 FET exhibits current saturation at higher drain biases due to the formation of depletion region on the drain side (pinch-off), as shown in Figure 5-10 (a). The transfer characteristics (Ids-Vbg) of single-layer WS2 FET with Ni contacts are presented in Figure The WS2 device shows n- type behavior with channel current Ids = 60 A (4 A/ m) at back-gate voltage Vbg = 60 V and bias Vds = 2 V. It is worth mentioning that there is no hysteresis behavior observed in the WS2 FET, which indicates the monolayer WS2 channel is intrinsically high-quality and clean Ni/WS2 49

50 and oxide/ws2 interfaces are achieved. The on/off ratio is larger than 10 6 at room temperature. The field-effect mobility can be estimated using Equation 5-2. The peak mobility of the singlelayer WS2 FETs reaches ~ 7 cm 2 V -1 S -1 at room temperature. Figure 5-9. Schematic and optical image of the single-layer back-gate WS2 FET. Symmetrical 50 nm Ni contacts are defined on the CVD-grown monolayer WS2 by e-beam lithography, evaporation, and lift-off. (a) (b) Figure (a) Saturation Ids-Vds behavior of a representative monolayer WS2 FET with Ni contacts (L ~ 1 m and W ~ 15 m). (b) Output performance of the devices at small fields. 50