二維材料石墨烯與過渡金屬雙硫屬化合物之光譜性質研究

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(1)Optical studies of two dimensional materials: graphene and transition metal dichalcogenides. Presented by Chih-Chiang Shen Advisor: Hsiang-Lin Liu, Ph. D. A Thesis Submitted for the Degree of Doctor of Philosophy. Department of Physics National Taiwan Normal University June 2014.

(2) Acknowledgements Time flies! Unconsciously, I have spent nine years with the Department of Physics, National Taiwan Normal University. This won’t be possible without the help and courtesy of the teachings of many professors, classmates and friends who accompanied me to complete my master's and doctoral studies. Thank You Lord, no matter what course or situation, Your preserving and guiding hand has guided me in this journey. First of all, I would like to thank my advisor Prof. Hsiang-Lin Liu for all his support, kind encouragements and countless consultations. His example, expertise and patience has molded a fresh student to become a qualified scientist. The present work has been a collaborative effort. I would like to greatly thank Dr. Lain-Jong Li of the Institute of Atomic and Molecular Sciences, Academia Sinica, his research group has provide us with high quality samples; S. H. Huang and F. C. Chou of the Center for Condensed Matter Sciences for providing MoS2 single crystals; and Dr. K. H. Chen, Dr. M. C. Chang, and Dr. Y. R. Chen for useful discussion. I would like to also thank all the lab members in the past and present for their contribution to the nice work atmosphere. I am grateful and proud of being part of this wonderful team. I would also like to acknowledge all my friends for encouraging me and sharing their life with me beyond the sphere of science. Finally, I would like to give my appreciation to my dearest family. Their patience, encouragement, and unconditional love have continuously given me the courage to face the many challenges in life. i.

(3) Abstract We present the results of THz absorption and spectroscopic ellipsometric measurements of triazine-doped graphene and monolayer transition metal dichalcogenides (MoS2 and MoSxSey). The triazine-doped graphene thin films were deposited on oxidized silicon substrate (SiO2/Si), using either chemical vapor deposition (CVD) or electrochemical exfoliation (ECE). Monolayer MoS2 and MoSxSey thin films were deposited onto sapphire substrates by CVD. Our aim is to investigate the charge dynamics and electronics structures of these novel materials. THz conductivity of all samples displays a coherent response of itinerant charge carriers at zero frequency. Notably, the CVD-grown graphene thin films with doping show an additional finite frequency peak at about 155 cm-1. A finite-frequency peak, which coexists with a Drude contribution, is likely associated with the significant disorder induced by triazine doping. Furthermore, as the temperature is lowered, the Drude plasma frequency (~ 21 and 7 THz for CVD-grown graphene with doping and MoS2 thin films) decreases, whereas the carrier relaxation time (~ 13 and 26 fs) does not show much temperature variation. These results suggest the semiconducting behavior of the CVD-grown graphene with doping and monolayer MoS2 thin films. Additionally, the Drude plasma frequency of the ECE-grown graphene thin films is three times larger than that of CVD-grown ones. In contrast, the carrier relaxation time of the ECEgrown graphene thin films (~ 10 fs) is shorter than that of the CVD-grown samples (~ 84 fs). Interestingly, the Drude plasma frequency of monolayer MoSxSey thin films is in the range from 6.5 to 8 THz. Carrier relaxation time is in the range from 19 to 26 fs. ii.

(4) The optical properties of all samples were also determined by spectroscopic ellipsometry. The absorption spectrum of the CVD-grown graphene thin films exhibits an asymmetric Fano resonance in the ultraviolet frequency region. This excitonic-dominated charge transfer band in the triazine-doped graphene thin films shows a blueshift in comparison with that of undoped analog. The line shape of the ECE-grown graphene thin films displays less asymmetric. Such behavior could be attributed to the changes of the charge distributions in the graphene thin films prepared by different growth methods. Additionally, monolayer MoSxSey films show a direct gap (~ 1.95 eV for MoS2 and ~ 1.62 eV for MoSe2). The ground-state exciton binding energy is found to be about 0.28 eV for MoS2 and 0.24 eV for MoSe2. These findings bring additional understanding of two-dimensional materials with respect to their charge dynamics and electronic structures and provide the foundation for future technological applications of these materials.. Keyword: Graphene, Transition metal dichalcogenides, Optical properties.. iii.

(5) Contents Acknowledgements……………………………………………………………………………….i Abstract…………………………………………………………………………………………..ii Contents……………………………………………………………………………………….…iv List of Figures……………………………………………………………………………………vi List of Tables…………………………………………………………………………………...xiv Chapter 1 Introduction………………………………………………………………………….1 Chapter 2 Brief survey of graphene and transition metal dichalcogenides………………….5 2-1 Graphene……………………………………………………………………………………5 2-1-1 Physical properties……………………………………………………………………..5 2-1-2 Optical response………………………………………………………………………..8 2-2 Transition metal dichalcogenides…………………………………………………………11 2-2-1 Physical properties........................................................................................................11 2-2-2 Optical response………………………………………………………………………13 Chapter 3 Theoretical background……………………………………………………………40 3-1 Optical theory……………………………………………………………………………..40 3-2 The Spectroscopic ellipsometry…………………………………………………………...48 3-3 Drude and Lorentz model…………………………………………………………………53 Chapter 4 Experimental techniques…………………………………………………………...60 4-1 Raman scattering spectroscopy…………………………………………………………...60 4-2 Optical spectrometers……………………………………………………………………..60 4-3 Spectroscopic ellipsometry………………………………………………………………..64 iv.

(6) Chapter 5 Optical properties of graphene with molecular doping………………………….72 5-1 THz spectra………………………………………………………………………………..74 5-2 Spectroscopic ellipsometric spectra……………………………………………………….79 5-3 Summary………………………………………………………………………………......82 Chapter 6 Optical properties of monolayer MoS2 films……………………………………...98 6-1 THz spectra………………………………………………………………………………..99 6-2 Spectroscopic ellipsometric spectra……………………………………………...............101 6-3 Summary…………………………………………………………………………………105 Chapter 7 Optical properties of monolayer MoSxSey films………………………………..115 7-1 THz spectra……………………………………………………………………………....117 7-2 Spectroscopic ellipsometric spectra……………………………………………...............118 7-3 Summary………………………………………………………………………………....121 Chapter 8 Thesis summary…………………………………………………………………...134 References……………………………………………………………………………………...137. v.

(7) List of Figures Fig. 1.1: Graphene is a two-dimensional building material for carbon materials of all other dimensionalities. It can be wrapped up into zero-dimensional buckyballs, rolled into one-dimensional nanotubes, or stacked into three-dimensional graphite [3]……….…..4 Fig. 1.2: Three-dimensional schematic representation of a typical MX2 structure, with the chalcogen atoms (X) inyellow and the metal atoms (M) in grey [36]………………….4 Fig. 2.1: Ambipolar electric field effect in single-layer grapheme. The insets show its low-energy dispersion, indicating the changes in the position of the Fermi energy EF with varying the gate voltage [3]………………………………….………………………………….18 Fig. 2.2: Optical image of an array of ion gel gated graphene FETs fabricated on a plastic substrate [52]…………………………………………………………………………..18 Fig. 2.3: Schematic illustration of bottom-gated graphene/GO transistor. The graphene channel which was monolithically patterned with source and drain electrodes is above the GO dielectric [53]…………………………………………….…………………………….19 Fig. 2.4: Schematic illustration and photo for electrochemical exfoliation of graphite [60]…….19 Fig. 2.5: Schematic illustration of CVD growth of graphene and its transfer process [66]……...20 Fig. 2.6: Raman-scattering spectra of graphite and graphene [89]………………………………20 Fig. 2.7: Evolution of 2D band as a function of number of layers for 514 nm excitation [89]….21 Fig. 2.8: Evolution of G band as a function of number of layers for 514 nm excitation [92]…..21 Fig. 2.9: The real part of the optical sheet conductance of graphite per layer [(a) experiment, (b) calculation)] as well as the calculated conductance of isolated undoped graphene (c). The inset of (c) depicts the optical transitions between hole and electron bands in monolayer graphene [93]……………………………………………...……………….22 Fig. 2.10: Absorption spectra for three different samples of graphene over the range of photon energies between 0.5 and 1.2 eV. The left scale gives the absorbance in units of πα, while the right scale gives the corresponding optical sheet conductivity in units of πG0/4.The black horizontal line corresponds to the universal result of an absorbance of πα = 2.293% with a range indicated of ± 0.1 πα or approximately ± 0.2% [94]……………………………...……………….…………………………………...23 Fig. 2.11: Measured graphene absorption spectra of samples 1 and 2 over a range of photon energies between 0.25 and 0.8 eV. Theory fits lines are based on a model of noninteracting massless fermions [94]……………...……………..………………….23 vi.

(8) Fig. 2.12: Absorption spectra of single-layer graphene. Dashed curves 1 and 2 are calculations from Yang et al. [96]. Solid curves 3 and 4 are experimental data. The symmetric peak at 5.2 eV (curve 1) is expected by noninteracting theory, whereas interaction effects should result in the asymmetric peak downshifted to 4.6 eV (curve 2). The predicted shift and asymmetry are in qualitative agreement with experimental data (curves 3 and 4) [95]…………………………………………………………………24 Fig. 2.13: Optical constants of graphene n (solid line) and k (dashed line) [97]………………...24 Fig. 2.14: Optical functions n and k shown for both a point-by-point fit and an optical dispersion model [98]……………………………………………………………………………..25 Fig. 2.15: Comparison of n and k values between CVD graphene [98] and exfoliated graphene by Kravets et al. [95] and Weber et al. [97]……………………………………………...25 Fig. 2.16: Schematic illustration of HfO2-top-gated monolayer MoS2 FET device [36]………...26 Fig. 2.17: A thin-film MoS2 electric double-layer transistors constructed with an ion gel on a plastic substrate [102]………………………………………………………………..26 Fig. 2.18: The dependence of the drain current at a gate voltage, VG, of 1.5 V (red) and the carrier mobility on the curvature radius. The carrier mobility is normalized by the results without bending (blue). The inset schematically illustrates the bending measurements [102]………………………………………………………………….27 Fig. 2.19: Schematic illustration and optical images of the MoS 2 thin film transistor under stretching [103]………………………………………………………………………27 Fig. 2.20: Top: the strain dependence of the drain current at a reference voltage, VR, of 1.6 V (red) and 0.3 V (blue). Bottom: the strain dependence of the on/off ratio (black) [103]………………………………………………………………………………….28 Fig. 2.21: Top: the electron mobility at various strains (red). Bottom: the specific capacitance of the ion-gel/MoS2 interface at 15 Hz (bottom, black) at various strains. The mobility is normalized by the results obtained in the absence of an applied tension. The blue square in the top panel corresponds to the normalized mobility after stretching at a 5% strain [103]……………………………………………………………………….28 Fig. 2.22: G vs.VG–VT at different temperatures, ranging from 298 to 78 K. The data show a noticeable increase of the dG/d(VG–VT) slope with decreasing T. Inset: μ vs. T for three different MoSe2 devices [50]…………………………………………………..29 Fig. 2.23: Schematic illustration of monolayer to few layer MoS2 by sulfurization of Mo thin film [112]…………………………………………………………………………….29 Fig. 2.24: Schematic of MoS2 layer deposited by two-step thermolysis, and the films obtained on a sapphire and silica substrate [113]…………………………………………………30. vii.

(9) Fig. 2.25: Schematic of CVD of MoS2 from solid S and MoO3 precursors. The red dots indicate the heating elements in the furnace [114]……………………………………………..30 Fig. 2.26: MoS2 shows great flexibility to surface corrugations [115]…………………………..30 Fig. 2.27: Thickness-dependence of Raman-scattering spectra for MoS2 [116]………………...31 Fig. 2.28: Schematic illustration of in-plane phonon modes E2g and the out-of-plane phonon mode A1g for the MoS2 [117]………………………………………………………...31 Fig. 2.29: Peak position shifts for the E2g and A1g modes as a function of MoS2 layer thickness for the spectra in Fig. 2.27 [116]………………………………………………………32 Fig. 2.30: Raman spectra acquired on a MoSe2 flake using excitation wavelengths of 442 nm [50]…………………………………………………………………………………...32 Fig. 2.31: Raman spectra acquired on a MoSe2 flake using excitation wavelengths of 532 nm [50]…………………………………………………………………………………...33 Fig. 2.32: Raman spectrum of single (solid red line) and more than 10 layers (dashed blue line) MoX2 (X = S, Se) [119]……………………………………………………………….33 Fig. 2.33: Raman spectra of MoSe 2 with various numbers of layers; the bulk spectrum is displayed at a ten fold scale [120]……………………………………………………34 Fig. 2.34: Band structures calculated from first principles density functional theory (DFT) for bulk and monolayer MoS2 [121]……………………………………………………..34 Fig. 2.35: Calculated band structure of single-layer and bulk MoSe2 [119]……………………..35 Fig. 2.36: Normalized PL spectra by the intensity of peak A of thin layers of MoS2 for N = 1-6. Feature I for N = 4-6 is magnified and the spectra are displaced for clarity [33]……..35 Fig. 2.37: PL spectra of MoS2 for monolayer, bilayer, hexalauer, and bulk [45]………………..36 Fig. 2.38: Photoluminescence spectra of MoS2 thin films with average thickness ranging from 1.3 to 7.6 nm. Inset of Fig. 2.38 shows energy of the A exciton peak as a function of average film thickness [106]…………………………………………………………..36 Fig. 2.39: Measured room-temperature photoluminescence on a single-layer (red), three-layer (blue dashed), and bulk (green dotted dashed) MoSe 2 . Here the measurement parameters including laser excitation intensity are the same [119]………………….37 Fig. 2.40: Temperature dependence of photoluminescence on (a−b) single-layer MoSe2 and MoS2 and (c−d) few-layer MoSe2 and MoS2 [119]………………………………….37 Fig. 2.41: Photoluminescence spectrum of MoSe2 up to three layers [120]……………………..38 Fig. 2.42: Absorption and photoluminescence spectra (red lines) in the range of 1.8–2.0 eV for the indicated back-gate voltages. The exciton (A) and trion (A－) resonances behave viii.

(10) differently with gate voltage. Left: Absorption spectra, with the dashed blue lines as a guide to the eye for the threshold energies of A and A－features. The green lines are power-law fits to the experimental results, as described in the main text, with the A and A components shown as the blue lines. Right: The photoluminescence spectra of the A and A－features are fit to Lorentzians (green lines). The dashed blue line indicates the absorption peak of the A－resonance and the arrows show the doping-dependent Stokes shift of the trion photoluminescence [123]……………………………………………39 Fig. 3.1: Light reflects and transmits in the interface……………………………………………58 Fig. 3.2: A beam of light incidents on an interface at angle ψ1, and then the reflectance beam from the substrate surface which has reflectance index N2…………………………….58 Fig. 3.3: The fitting model of films………………………………………………………………59 Fig. 4.1: A sketch of the setup of the micro-Raman scattering…………………………………..67 Fig. 4.2: The sketch map of the FTIR spectrometer. (S: light sources, A: aperture, and D: detectors)………………………………………………………………………………68 Fig. 4.3: The sketch map of the grating-type monochromatic spectrometer…………………….68 Fig. 4.4: The schematic diagram of a null ellipsometer………………………………………….69 Fig. 4.5: The schematic diagram of a polarization modulation ellipsometer…………………….69 Fig. 4.6: The schematic diagram of a rotating polarizer ellipsometer…………………………...70 Fig. 4.7: The schematic diagram of a rotating analyzer ellipsometer……………………………70 Fig. 4.8: The schematic diagram of a rotating compensator ellipsometer……………………….71 Fig. 5.1: AFM images of (a) CVD-undoped thin film, (b) CVD-doped thin film, (c) ECEundoped thin film, and (d) ECE-doped thin film……………………………………...84 Fig. 5.2: Raman-scattering spectra of the CVD-undoped, CVD-doped, ECE-undoped, and ECEdoped graphene thin films………………………………………………………………85 Fig. 5.3: Room-temperature optical transmission spectra of a bare silicon substrate, CVDundoped, and CVD-doped graphene thin films……………………………………….86 Fig. 5.4: Real part of the room-temperature optical conductivity spectrum of the CVD-doped graphene thin film (solid line). The various terms in the fits are also shown (dashed line): the Drude band and two Lorentz oscillators…………………………………….86 Fig. 5.5: Temperature dependence of the Drude plasma frequency and the carrier relaxation time……………………………………………………………………………………87 Fig. 5.6: Temperature dependence of the Drude conductivity…………………………………...87 ix.

(11) Fig. 5.7: Measured optical sheet conductance spectra (solid line) of the CVD-undoped and CVDdoped graphene thin films at 10 K along with the theoretical spectrum (dashed line)…88 Fig. 5.8: Room-temperature optical transmission spectra of a bare silicon substrate, CVDundoped, CVD-doped, ECE-undoped, and ECE-doped graphene thin films…………89 Fig. 5.9: Real part of the room-temperature optical conductivity spectrum of the CVD-undoped, CVD-doped, ECE-undoped, and ECE-doepd graphene thin film (solid line). The fits term with an Drude band are also shown (dashed line)………………………………...89 Fig. 5.10: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the undoped graphene thin films prepared by CVD…………………………………………………………...90 Fig. 5.11: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the triazine-doped graphene thin films prepared by CVD………………………………………………...90 Fig. 5.12: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the undoped graphene thin films prepared by ECE……………………………………………………………91 Fig. 5.13: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the triazine-doped graphene thin films prepared by ECE…………………………………………………91 Fig. 5.14: Room temperature refraction index for the triazine-doped graphene thin films compared with those of the undoped thin films prepared by CVD and ECE………..92 Fig. 5.15: Room temperature extinction coefficient for the triazine-doped graphene thin films compared with those of the undoped thin films prepared by CVD and ECE………..92 Fig. 5.16: Optical absorption coefficient (symbols) of the undoped and triazine-doped graphene thin films prepared by CVD and ECE. The dashed lines are the results of the fitting using the Fano model………………………………………………………………...93 Fig. 5.17: Temperature dependence of frequency, linewidth, normalized intensity, and asymmetric parameter of the excitonic resonance of CVD-undoped graphene thin films………………………………………………………………………………...94 Fig. 5.18: Temperature dependence of frequency, linewidth, normalized intensity, and asymmetric parameter of the excitonic resonance of CVD-doped graphene thin films………………………………………………………………………………...95 Fig. 5.19: Temperature dependence of frequency, linewidth, normalized intensity, and asymmetric parameter of the excitonic resonance of ECE-undoped graphene thin films………………………………………………………………………………...96. x.

(12) Fig. 5.20: Temperature dependence of frequency, linewidth, normalized intensity, and asymmetric parameter of the excitonic resonance of ECE-doped graphene thin films………………………………………………………………………………...97 Fig. 6.1: Raman-scattering spectra of a monolayer-CVD-MoS 2 film and a MoS 2 single crystal………………………………………………………………………….........107 Fig. 6.2: Temperature-dependence Raman-scattering spectra of a monolayer-CVD-MoS 2 film………………………………………………………………………………….107 Fig. 6.3(a): Temperature dependence of frequency, damping, oscillator strength, and the asymmetry factor of E2g phonon mode of a monolayer-CVD-MoS2 film. The thin solid lines are results of the ﬁtting taking into account the temperature-induced anharmonicity……………………………………………………………………108 Fig. 6.3(b): Temperature dependence of frequency, damping, oscillator strength, and the asymmetry factor of A1g phonon mode of a monolayer-CVD-MoS2 film. The thin solid lines are results of the ﬁtting taking into account the temperature-induced anharmonicity……………………………………………………………………109 Fig. 6.4: Room-temperature optical transmission spectra of a bare silicon substrate and monolayer-CVD-MoS2 film………………………………………………………...110 Fig. 6.5: Real part of room-temperature optical conductivity spectrum of a monolayer-CVDMoS2 film (solid line). The various terms in the fits are also shown (dashed line): the Drude band and one Lorentz oscillator………………………………………………110 Fig. 6.6: Temperature dependences of Drude plasma frequency and carrier relaxation time….111 Fig. 6.7: Temperature dependence of the Drude conductivity………………………………….111 Fig. 6.8: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of a monolayer-CVDMoS2 film…………………………………………………………………………….112 Fig. 6.9: Refractive index n and extinction coefficient k of a monolayer-CVD-MoS2 film……112 Fig. 6.10: Room-temperature optical absorption coefficient of a monolayer-CVD-MoS2 film compared with that of a single crystal. The dashed line is the best fit using a broadened Lorentzian line shape…………………………………………………...113 Fig. 6.11: Temperature dependence of band-gap energy and Bose-Einstein model fitting results (thin solid line)……………………………………………………………………...113 Fig. 6.12: Temperature dependence of A and B excitons in a monolayer-CVD-MoS2 film…...114 Fig. 6.13: Optical absorption coefficient spectra of monolayer-CVD-MoS2 film at 10 K……..114 Fig. 7.1: AFM images of a monolayer-CVD-MoS2 film with selenization temperature of (a) 600 ˚C, (b) 700 ˚C, (c) 800 ˚C, and (d) 900 ˚C…………………………………………….123 xi.

(13) Fig. 7.2 (a): Step height measurement from substrate to 600 ˚C sample (0.73 nm)……………124 Fig. 7.2 (b): Step height measurement from substrate to 700 ˚C sample (0.65 nm)……………124 Fig. 7.2 (c): Step height measurement from substrate to 800 ˚C sample (0.71 nm)……………125 Fig. 7.2 (d): Step height measurement from substrate to 900 ˚C sample (1.1 nm)……………..125 Fig. 7.3: Room-temperature Raman scattering spectra of the monolayer-CVD-MoS2 films with different selenization temperature……………………………………………………126 Fig. 7.4: Room-temperature optical transmission spectra of a bare silicon substrate, 600 ˚C, 700 ˚C, 800 ˚C, and 900 ˚C sample………………………………………………………..127 Fig. 7.5: Room temperature refraction index for the monolayer-CVD-MoS2 films with different selenization temperature over a frequency range of 0.5 – 3 THz……………………..127 Fig. 7.6: Real part of room-temperature optical conductivity spectrum of a monolayer-CVDMoS2 film with selenization temperature of 600 ˚C (symbols). The fits term with a Drude band is also shown (dashed line)……………………………………………..128 Fig. 7.7: Real part of room-temperature optical conductivity spectrum of a monolayer-CVDMoS2 film with selenization temperature of 700 ˚C (symbols). The fits term with a Drude band is also shown (dashed line)……………………………………………..128 Fig. 7.8: Real part of room-temperature optical conductivity spectrum of a monolayer-CVDMoS2 film with selenization temperature of 800 ˚C (symbols). The fits term with a Drude band is also shown (dashed line)……………………………………………..129 Fig. 7.9: Real part of room-temperature optical conductivity spectrum of a monolayer-CVDMoS2 film with selenization temperature of 900 ˚C (symbols). The fits term with a Drude band is also shown (dashed line)……………………………………………..129 Fig. 7.10: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the monolayer-CVDMoS2 film with selenization temperature of 600 ˚C…………………………………130 Fig. 7.11: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the monolayer-CVDMoS2 film with selenization temperature of 700 ˚C…………………………………130 Fig. 7.12: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the monolayer-CVDMoS2 film with selenization temperature of 800 ˚C…………………………………131 Fig. 7.13: Room temperature experimental (symbols) at 60º and 70º incidence angles and fitted (dashed lines) values of ellipsometric parameters of Ψ and ∆ of the monolayer-CVDMoS2 film with selenization temperature of 900 ˚C…………………………………131 xii.

(14) Fig. 7.14: Room temperature refractive index for the monolayer-CVD-MoS2 films with different selenization temperature in the visible frequency region………………………........132 Fig. 7.15: Room temperature extinction coefficient for the monolayer-CVD-MoS2 films with different selenization temperature in the visible frequency region…………………..132 Fig. 7.16: Room-temperature optical absorption coefficient of monolayer-CVD-MoS2 films with different selenization temperature. The dashed line is the best fit using a broadened Lorentzian line shape………………………………………………………………...133 Fig. 7.17: A and B excitons in different Mo-Se percentage. The Mo-Se percentage for selenized at 600 ˚C, 700 ˚C, 800 ˚C, and 900 ˚C is 14.5 %, 73.8 %, 95 %, and 100 % respectively [161]………………………………………………………………………………….133. xiii.

(15) List of Tables Table 5.1: Parameters of a Lorentz fit for the Raman-scattering spectra of four thin films. All units are in cm-1………………………………………………………………………73 Table 5.2: Parameters of a Drude-Lorentz fit for the room-temperature transmission data. All units are in cm-1………………………………………………………………………75 Table 5.3: Parameters of a Drude fit for the room-temperature transmission data. All units are in cm-1…………………………………………………………………………………..79 Table 5.4: Parameters of a stacked layer model fit for the undoped and nitrogen-doped graphene thin films. All units are in nm………………………………………………………..80 Table 5.5: Parameters of a Fano fit for the room-temperature optical absorption data………….81 Table 6.1: Parameters of a Drude-Lorentz fit for room-temperature transmission data. All units are in cm-1. Drude plasma frequency = ωpD /2π ~ 7 THz; carrier scattering time = τD ~ 26 fs…………………………………………………………………………………100 Table 6.2: The exciton band-gap energies, exciton binding energies, and exciton broadening parameters of a monolayer MoS2 film……………………………………………...103 Table 7.1: Parameters of a Lorentzian fit for Raman-scattering spectra data. All units are cm1 ……………………………………………………………………………………..116 Table 7.2: Parameters of a Drude model fit for the room-temperature transmission data. All units are in cm-1…………………………………………………………………………..117 Table 7.3: Properties of carriers obtained from calculation fitting parameters………………...118 Table 7.4: Parameters of a stacked layer model fit for four different selenization temperature samples……………………………………………………………………………...120 Table 7.5: Exciton peak position and spin-orbit splitting value of all samples………………...120 Table 7.6: The exciton band-gap energies, exciton binding energies, and exciton broadening parameters of all samples…………………………………………………………...121. xiv.

(16) Chapter 1 Introduction Two-dimensional (2D) materials have recently attracted enormous attention due to both scientific and technologic aspects. Many of them exist in bulk form as stacks of strongly bonded layers with weak interlayer attraction that can be exfoliated as stable individual atomically thin layers [1]. In 2004, Novoselov et al. [2] have demonstrated that it is not only possible to exfoliate stable monolayer graphene from van der Waals solids, but that graphene also can exhibit unique and fascinating physical properties. Graphene is a monolayer of sp2 bonded carbon atoms packed into a honeycomb crystal structure and can be viewed either as an individual atomic plane extracted from graphite or unrolled single-wall carbon nanotubes (Fig.1.1) [3]. The fascinating properties of graphene, such as high thermal conductivity (~ 3000 W/mK) [4], high roomtemperature carrier mobility (~ 2.5 × 105 cm2 V−1 s−1) [5], strong Young’s modulus (~ 1 TPa) [6], and high surface area (~ 2630 m2/g) [7] have been well documented. The electronic band structure of graphene has a linear dispersion near the K point, giving rise to novel phenomena such as the anomalous room-temperature quantum Hall effect (QHE). This is due to the charge carriers in graphene, which behave as massless relativistic particles (Dirac fermions). It has opened up a new category of “Fermi-Dirac” physics [8]. Furthermore, graphene is known as a gapless semiconductor. This lack of an energy gap directly leads to its low on-off current ratio [9]; hence, opening of an electrical band gap in graphene is crucial for its application in logic circuits and photonic devices [10,11]. Various graphene superstructures with 1.

(17) quantum confinement effect such as graphene nanoribben and graphene dots (to break the energy degeneracy of the electronic bands) have been recently proposed [12-16]. Doping graphene with other elements is an alternative method for achieving this goal [17-19]. In this method, doping atoms change graphene’s lattice structure by forming covalent bonds with C atoms. This modifies the electronic structure of graphene and suppresses graphene’s density of states near the Fermi energy level, thus opening up a gap between valence and conduction bands [20]. Modulating the bands structure of graphene further expands its scope of applications. There has been particular interest recently in other 2D materials such as the transition metal dichalcogenides (TMDs) [21,22], transition metal oxides including titania- and perovskite- based oxides [23,24], and boron nitride (BN) [25,26]. In particular, TMDs show a wide range of electronic, optical, mechanical, chemical, and thermal properties [22,27,28]. They possess potential applications such as catalysis [29], nanotribology [30], energy conversion [31,32], optoelectronics [33,34], and superlubricity in nanomachines [35]. TMDs including MX2, where M is a transition metal element from group IV (Ti, Zr, Hf ), group V (V, Nb, Ta) or group VI (Mo, W), and X is a chalcogen (S, Se, Te) [36]. These materials form layered structures of the form X-M-X, with the chalcogen atoms in two hexagonal planes separated by a plane of metal atoms (Fig. 1.2) [36]. Depending on the selection of the metal elements, these layered materials exhibit superconducting (ex. NbSe2, TaS2), metallic (NbS2, VSe2), or semiconducting properties (ex. MoS2, MoSe2) [37-43]. Recently, transistors based on exfoliated monolayer MoS2 have shown field-effect mobility of up to 200 cm2 V-1 s-1 and a high on-off current ratio of 108, which are both necessary qualities in electronic circuits requiring low standby power [36,44]. The 1.9 eV large direct band gap of monolayer MoS2 results in strong photoluminescence, making it attractive for light-emitting 2.

(18) devices [33,45]. Moreover, monolayer MoS2 has been considered an electronically confined system ideal for fundamental studies of broken inversion symmetry and strong spin-valley coupling [46]. These properties make possible a new class of integration in spintronic and valleytronic applications [47-49]. Transistors using ultra-thin mechanically exfoliated MoSe2 flake on SiO2/Si also possess a good gate control, with on/off ratios larger than 106 [50]. This renders them attractive as an ultra-thin body for aggressively scaled FETs. More recently, Bernardi et al. [51] have demonstrated the exfoliated monolayer MoS2, MoSe2, and WS2 can absorb up to 5-10% incident sunlight in a thickness of less than 1 nm. This have been shown to achieve 1 order of magnitude higher sunlight absorption than most commonly used absorbers in solar cells GaAs and Si. This suggests that the monolayer semiconducting TMDs hold potential as sunlight absorbers, and may enable ultrathin photovoltaic (PV) device [51]. The rest of this thesis is organized as follows. Chapter 2 presents a review of the previous research works in graphene and two-dimensional transition metal dichalcogenides. Chapter 3 discusses the basic theory about the general optical properties of solids and spectroscopy ellipsometry. Chapter 4 shows the Raman-scattering spectroscopy, optical spectrometer, and spectroscopy ellipsometry used in this study. Chapter 5 describes the optical properties of graphene with molecular doping. Chapter 6 describes the optical properties of monolayer MoS2 films. Chapter 7 describes the optical properties of monolayer MoSxSey films. Finally, a thesis summary will be given in Chapter 8.. 3.

(19) Fig. 1.1: Graphene is a two-dimensional building material for carbon materials of all other dimensionalities. It can be wrapped up into zero-dimensional buckyballs, rolled into onedimensional nanotubes, or stacked into three-dimensional graphite [3].. Fig. 1.2: Three-dimensional schematic representation of a typical MX2 structure, with the chalcogen atoms (X) inyellow and the metal atoms (M) in grey [36]. 4.

(20) Chapter 2 Brief survey of graphene and transition metal dichalcogenides 2-1 Graphene 2-1-1 Physical properties One of the most important properties of graphene is its electric field effect. In 2007, Geim et al. [3] presented electric field effect in exfoliated monolayer graphene on 300 nm SiO 2 at low temperature (Fig. 2.1). The curve exhibits a clear ambipolar conduction, indicating graphene displays ambipolar electric field effect. Thus, charge carriers can be tuned continuously between electrons and holes. Positive (negative) voltage Vg corresponds to electron (hole) transport. The insets of Fig. 2.1 show its conical low-energy dispersion, indicating changes in the position of the Fermi energy EF with changing gate voltage Vg. The rapid decrease in resistivity ρ on adding charge carriers indicates their high mobility (in this case, μ ~ 5000 cm2V-1s-1). In addition, the charge carrier mobility weakly depends on temperature. Even at 300 K the mobility is still limited by impurity scattering. Recently, Mayorov et al. [5] have demonstrated the roomtemperature carrier mobility up to ~ 2.5 × 105 cm2V-1s-1. This attracts a lot of attention to graphene as a possible material for high frequency device. In 2010, Kim et al. [52] reported a high-performance flexible graphene field-effect 5.

(21) transistors with ion gel gate dielectrics (Fig. 2.2). The graphene FETs fabricated on the plastic substrates showed a hole and electron mobility of 203 and 91 cm2 V-1s-1, respectively. Moreover, ion gel gated graphene FETs on the plastic substrates exhibited remarkably good mechanical flexibility. Recently, Lee et al. [53] presented all graphene-based thin film transistors. The FETs was fabricated on flexible plastic substrates using a graphene active layer, graphene oxide dielectrics, and graphene electrodes (Fig. 2.3). The graphene-based FETs showed a hole and electron mobility of 300 and 250 cm2 V-1s-1, respectively. The performance is better than previous reported by Kim et al. [52]. These methods explore a significant step for the application of graphene toward flexible and stretchable electronics. Moreover, graphene has been involved in highly simple and desirable geometries for stretchable electrodes and transistors [54-56] due to atomic thickness and very large Young’s modulus [6]. Ripple relaxation allows the graphene films to stretch to values in excess of 20% [54,55], while graphene transistors fabricated on rubber substrates can operate at a mechanical strain of 5% [56]. Mechanically exfoliated [2] and epitaxially grown graphene films [57-59] exhibit high quality but are not suitable for large-scale production. In 2011, Su et al. [60] demonstrated a simple and fast electrochemical method to exfoliate graphite into thin graphene sheets (Fig. 2.4), mainly bilayered graphene with a large lateral size up to 30 μm [60]. Recent works on CVD methods using catalytic metal substrates, nickel (Ni) and copper (Cu), have shown the capability of growing large-area graphene, greatly encouraging their applications in highly transparent and flexible conducting films [61-65]. For example, Srivastava et al. [66] prepared the centimeter size (~ 3.5 cm × 1.5 cm), uniform, and continuous single and few-layer graphene films by using chemical vapor deposition technique on polycrystalline Cu 6.

(22) foils with liquid precursor hexane (Fig. 2.5). In addition, Lu et al. [67] reported that the cooling process in CVD-grown graphene affects the crystallographic orientation of the Cu grains underlying graphene. They demonstrated the crystal structure of the Cu grains under graphene layers is governed by two competing processes: (1) graphene induced Cu surface reconstruction favoring the formation of Cu(100) orientation, and (2) recrystallization from bulk Cu favoring Cu(111) formation. The underlying Cu grains, regardless of reconstruction or recrystallization, induce a large hydrostatic compression to the graphene lattice. In spite of the remarkable progress in the growth of graphene films, their gapless properties limit the applications in logic circuits and photonic devices. Doping graphene with different kinds of absorbents like NH3 as electron donor [68] and NO2 as electron acceptor [69] is an effective way to open the band gap. In addition, nitrogen-doped (N-doped) graphene has been recently proved to be an effective approach to tailor the property of graphene and greatly broaden its applications [70]. For example, nitrogen-doped graphene/CdS heterostructures show a higher photocatalytic activity than pure CdS [71]. Nitrogen-doped graphene quantum dots exhibit strong blue luminescence [72]. These results suggest the possibility to use nitrogen-doped graphene in high performance photocatalytic and biological imaging applications. Currently, the widely used method for preparing nitrogen-doped graphene thin films involves arc charge method [73], segregation of trace amount of carbon and nitrogen in bulk metals [74], and various kinds of chemical vapor deposition [75-81]. Very recently, Lu et al. [82] synthesized few-layer nitrogen-doped (N-doped) graphene sheets by chemical vapor deposition using organic molecule 1, 3, 5-triazine as a solid precursor on Cu metal catalyst. By reducing the growth temperature, the atomic percentage of nitrogen doping is raised. In addition, with increasing doping concentration, N-doped graphene sheet exhibits a crossover from p-type to n7.

(23) type behavior accompanied by a strong enhancement of electron-hole transport asymmetry. Opening the band gap in graphene with molecular doping has been reported by different researchers [83-88]. In earlier studies, Ohta et al. [83] reported angle-resolved photoemission spectroscopy (ARPES) measurements of bilayer graphene grown on a silicon carbide (SiC) substrate. They found that by controlling potassium doping in a bilayer of graphene, the magnitude of the band gap between the valence and conduction bands could be manipulated up to ~ 70 meV. Zhou et al. [84] also performed ARPES measurements of single and bilayer epitaxial graphene with adsorbed NO2. NO2 induces hole doping of graphene over a wide doping range, tuning the charge carriers from electrons to holes results in a rigid shift of the band structure. Coletti et al. [85] examined charge neutrality and band-gap tuning of epitaxial graphene on SiC by molecular doping. They showed that the band structure of epitaxial graphene on SiC (0001) can be precisely tailored by functionalizing graphene’s surface with F4-TCNQ molecules. Yavari et al. [86] found that a tunable band gap of up to ~ 0.206 eV can be engineered in graphene by the controlled adsorption of water molecules to the surface of graphene. Matis et al. [87] presented the results of temperature-dependent electronic transport measurements of hydrogenating graphene. They demonstrated that a band gap of up to 50 meV emerges at the charge neutrality point and the size of the gap is tunable with an electric field effect and/or the hydrogen coverage. Recently, we have discovered that the stable organic molecule triazine can be thermally evaporated to form a very uniform layer on the surface of bilayer grapheme. The process opens up a band gap of up to ~ 111 meV, depending on doping concentration [88].. 2-1-2 Optical response 8.

(24) In 2006, Ferrari et al. [89] compares the Raman-scattering spectra of graphene and bulk graphite measured at 514.5 nm excitation (Fig. 2.6). The two most intense features are the G band at about 1580 cm−1 and 2D band at about 2700 cm−1. The G band of graphite materials is a doubly degenerate (TO and LO) phonon mode at the Brillouin zone center (Γ point) [90]. The 2D band is due to second-order Raman scattering involving TO phonons near the K point [90]. The evolution of the 2D band for different graphene sheets has been used for determining graphene thickness as well as for probing electronic structures through the double resonance process [89,91]. Figure 2.7 shows the evolution of the 2D band in nGLs as a function of the numbers of layers [89]. Bilayer graphene has a much broader and upshifted 2D band with respect to monolayer graphene. The shape of this band is quite different from that of bulk graphite. In contrast, the shape of the G band does not change with the number of graphene, as shown in Fig. 2.8 [92]. Only a slight frequency upshift can be seen in the case of monolayer graphene. In 2008, Kuzmenko et al. [93] reported the optical conductance of graphite due to the transitions between hole and electron bands is very close to the universal value of (π/2)e2/h between 0.1 and 0.6 eV, which is the theoretically expected value of dynamical conductance of isolated monolayer graphene (Fig. 2.9 (a)). The sheet conductance per graphene layer was calculated using the relation G (ω) = dc σ (ω), where dc = 0.34 nm is the interlayer distance. In addition, the conductance at low energies shows a strong depletion with increasing temperature in a fashion very similar to the temperature dependence of G1 (ω) in graphene (Fig. 2.9 (c)). The depletion of low-energy conductance with temperature is due to the gradual equilibration of the electron and hole occupation numbers close to the Fermi level. In contrast to the graphene, the conductance of graphite shows a Drude peak below 10-20 meV, an extra structure at about 50 meV, and two broad peak at about 0.7 and 0.9 eV (Fig. 2.9 (b)). In general, they conclude that 9.

(25) the isolated graphene is also presented in graphite in a broad energy range, in spite of the modification of the band structure by a significant c-axis hopping. In 2008, Mak et al. [94] presented the measured of the optical conductivity of graphene on SiO2 substrate between 0.5 and 1.2 eV. For photon energies above 0.5 eV, they observed graphene yielded a spectrally flat optical absorbance of (2.3 ± 0.2)%, which is in agreement with a constant absorbance of πα (α = 2πe2/hc denotes the fine-structure constant), or a sheet conductivity of (π/2)e2/ h, predicted within a model of noninteracting massless Dirac fermions (Fig. 2.10). In addition, there is a significant deviation from the value of the universal absorbance at lower photon energies (Fig. 2.11). This “nonuniversal” behavior is explained by including the effects of doping and finite temperature, as well as contributions from intraband transitions. In 2010, Kravets et al. [95] presented the first variable angle (45° ~ 70°) spectroscopic ellipsometry measurements of exfoliated graphen flakes on amorphous quartz in the wavelength range 245 ~ 750 nm. They reported an asymmetric shape of absorption peak at 4.6 eV because of a van Hove singularity in graphene’s density of state, which is modified by strong excitonic effect (Fig.2.12). The dashed curves 1 and 2 in Fig.2.12 are theoretical study by Yang et al. [96] for the noninteracting case and with excitonic effects included, respectively. In 2010, Weber et al. [97] reported the optical constants of mechanically exfoliated graphene flake on SiO2/Si measured by spectroscopic ellipsometry at an angle of 55° in the wavelength range 210 ~ 1000 nm. An intense peak in k (extinction coefficient) is observed at 270 nm (4.6 eV), which assigned as the effect of strong resonant excitons (Fig. 2.13). In 2010, Nelson et al. [98] performed spectroscopic ellipsometry measurements of CVD-grown graphene on glass substrate in the wavelength range 245 ~ 1700 nm (Fig. 2.14). They observed a resonant exciton absorption peak at about 4.5 eV. In addition, the optical functions showed similar behavior to previous values 10.

(26) reported by Weber et al. [97] for exfoliated graphene on SiO2/ Si, as well as those obtained by Kravets et al. [95] of exfoliated graphene on amorphous quartz (Fig. 2.15).. 2-2 Transition metal dichalcogenides 2-2-1 Physical properties In 2011, Radisavljevic et al. [36] presented the first top-gated transistor based on monolayer MoS2, as shown in Fig. 2.16. This FET exhibits excellent on/off current ratio (~ 108), n-type conduction, room-temperature mobility of > 200 cm2V-1s-1, and subthreshold swing of 74 meV per decade. Moreover, the mechanical strength of monolayer MoS2 is 30 times higher than that of steel [99-101]. The robustness of MoS2 allows it to endure deformation up to 11% before breaking due to the stiff chemical bonds of the Mo-S networks [100]. With the advantages of these electronic and flexible properties, Pu et al. [102] have demonstrated highly flexible MoS2 TFTs on plastic substrates with ion gel gate dielectrics (Fig. 2.17). The transistor exhibited excellent band transport with low threshold voltage (< 1 V), mobility (12.5 cm 2V-1s-1), and a high on/off current ratio (~ 105). Furthermore, the MoS2 TFTs exhibited remarkably high mechanical flexibility, and no degradation in the electrical characteristics was observed when they were bent to a curvature radius of 0.75 mm (Fig. 2.18). The superior electrical performance and excellent pliability of MoS2 films make them suitable for use in large-area flexible electronics. In 2013, Pu et al. [103] fabricated stretchable MoS2 TFTs on PDMS using ion gels as elastic gate dielectrics (Fig. 2.19). The TFTs exhibited an electron mobility of 1.4 cm2V-1s-1 and an on/off current ratio of 104 with a notably low threshold voltage (~ 1 V). Additionally, MoS2 TFTs operated at a mechanical strain of 5% without significant degradation of their 11.

(27) electrical properties (Figs. 2.20-2.21). These results demonstrate the potential for using MoS2 films for stretchable electronics. In 2012, Larentis et al. [50] reported n-type field-effect transistors on ultra-thin mechanically exfoliated MoSe2 flakes. The MoSe2 FETs possess a high gate modulation, with on/off current ratio larger than 106. In addition, the room temperature mobility is as high as ~ 50 cm2V-1s-1, and increases almost four fold when reducing the temperature to 78 K (Fig. 2.22). This suggests phonon scattering plays a dominant role at room temperature. Currently, the most widely used method for preparing thin-layer MoS2 with the dimensions of several micrometers involves various kinds of exfoliation [33,36,45,104-107], physical vapor deposition [108], hydrothermal synthesis [109,110], and thermolysis of the precursor containing Mo and S atoms [111]. However, synthesis of large-area and uniform layers is important for applications such as wafer-scale fabrication of electronic devices and flexible, transparent optoelectronics. Recently, several CVD methods for growing large-area atomically thin films of MoS2 on insulating substrates have been reported [112-114]. These methods use different solid precursors heated to high temperatures: (i) a thin layer of Mo metal deposited onto a wafer heated with solid sulphur (Fig. 2.23) [112] and (ii) substrates dip-coated in a solution of (NH4)2MoS4 and heated in the presence of sulphur gas (Fig. 2.24) [113]. Notably, Liu et al. [113] prepared highly crystalline and large-area tri-layer MoS2 by two-step thermolysis. The addition of sulphur during the second annealing process improved the crystallinity of MoS2. More recently, Lee et al. [114] developed large-scale synthesis of monolayer-MoS2 films up to several millimeters in dimensions using CVD. In this method, MoO3 and S powers vaporized and co-deposited on SiO2/Si substrate (Fig. 2.25). The growth condition is very sensitive to the 12.

(28) surface treatment, where aromatic molecules such as reduced graphene oxides (rGO), perylene3,4,9,10-tetracarboxylic acid tetrapotassium salt (PTAS) and perylene-3,4,9,10-tetracarboxylic dianhydride (PTCDA), promote layer growth. Subsequent breakthroughs have been made in the levels of purity and quality of monolayer-MoS2 film synthesis on various substrates (Fig. 2.26). This has been achieved by allowing enough flexibility to accommodate surface corrugations through careful optimization of CVD growth conditions [115].. 2-2-2 Optical response Lee et al. [116] presented the Raman-scattering studies of monolayer and few-layer MoS2 films (Fig. 2.27). The spectrum of a bulk MoS2 is composed of two main optical phonon modes at about 382 and 406 cm–1, displaying symmetries of E2g and A1g. The in-plane E2g mode results from opposite vibration of two S atoms with respect to the Mo atom while the A1g mode is associated with the out of plane vibration of only S atoms in opposite directions (Fig. 2.28) [117]. For decreasing layer numbers, the E2g mode increases in frequency whereas the A1g mode decreases (Fig. 2.29) [116]. The origins of the shifts have been identified as the influence of neighbouring layers on the effective restoring forces on atoms and the increase of dielectric screening of long-range Coulomb interactions [117]. Similar evidence has also been observed in the mechanically exfoliated MoS2 flakes [116,118]. In 2012, Larentis et al. [50] shows the Raman-scattering spectra using 442 nm (Fig. 2.30) and 532 nm (Fig. 2.31) excitation wavelength on ultra-thin MoSe2. Fig. 2.30 shows four peaks at 169, 242, 285, and 352 cm-1, corresponding to the E1g, A1g, E2g, and A2u modes, respectively. Notably, the A2u mode is an infrared active mode, not present in Raman-scattering spectra in bulk samples. The emergence of this mode in Raman-scattering spectra acquired on small flakes 13.

(29) suggests a breakdown of inversion symmetry, possibly because of the substrate [50]. The Raman-scattering spectra of Fig. 2.31 exhibits similar peaks as Fig. 2.30, but a higher intensity of the A1g (242 cm-1) with respect to other modes. In 2012, Tongay et al. [119] reported the Ramanscattering spectra of single and few-layer mechanical exfoliated MoS2 and MoSe2 flakes onto 90 nm SiO2/Si substrate. Exfoliated few-layer flakes have shown characteristic A1g and E2g Raman modes located at 243.0 and 283.7 cm−1 for MoSe2 and 408.7 and 383.7 cm−1 for MoS2 (Fig. 2.32). In the single layer limit, the A1g Raman mode softens to 241.2 (406.1) cm−1 as the E2g mode stiffens to 287.3 (384.7) cm−1 for MoSe2 (MoS2). The shifts behavior consistent with earlier studies [116-118]. Very recently, Shaw et al. [120] presented CVD growth of monolayer MoSe2 nanosheets. They observed the most prominent and identifiable peak of the Raman-scattering spectra is the A1g mode, which softening from 243.7 cm-1 for bulk MoSe2 to 241.2 cm-1 for monolayer MoSe2 (Fig. 2.33). In addition, the lower frequency and less noticeably split peaks located at 239 cm-1, 240.1 cm-1 and 240.8 cm-1 for three-, four- and five- layer MoSe2 result from the Davydov splitting [120]. First-principle calculations predict that semiconducting TMDs exhibit indirect to direct band gap transformation with decreasing layer numbers [117,121,122]. The band structures of bulk and monolayer MoS2 calculated from first principles are shown in Fig. 2. 34 [121]. At the Γ point, the band gap transition is indirect for the bulk material, but gradually shifts to be direct for the monolayer. The change in the band structure with layer number is due to quantum confinement and the resulting change in hybridization between pz orbitals on S atoms and d orbitals on Mo atoms [33,45]. The electronic distributions are also spatially correlated to the atomic structure [45]. For MoS2, density functional theory (DFT) calculations show that the conduction band states at the K point are mainly due to localized d orbitals on the Mo atoms, 14.

(30) located in the middle of the S–Mo–S layer sandwiches and relatively unaffected by interlayer coupling. However, the states near the Γ point are due to combinations of the antibonding pz orbitals on the S atoms and the d orbitals on Mo atoms, and have a strong interlayer coupling effect [45]. Therefore, as the layer numbers change, the direct excitonic states near the K point are relatively unchanged, but the transition at the Γ point shift significantly from an indirect one to a larger, direct one. The band structures of bulk and monolayer MoSe2 calculated from first principles are shown in Fig. 2. 35 [119]. Bulk MoSe2 displays 0.84 eV (Γ to Γ−K), 1.1 eV (K to Γ−K) indirect bandgap, and a 1.34 eV (K−K) direct band gap. In contrast, for monolayer MoSe2, Γ to Γ−K and K to Γ−K increases, while the K−K direct gap remains nearly unchanged and monolayer MoSe2 shows a direct 1.34 eV band gap at the K symmetry point. The changes of the electronic band structures in TMDs manifest itself as a strong PL feature. Mak et al. [33] observed an increase of the photoluminescence (PL) quantum yield by more than a factor of 104 for the mechanically exfoliated monolayer MoS2 compared with the bulk crystal. Furthermore, the main peak of the monolayer MoS2 PL spectrum is the direct gap luminescence feature at 1.9 eV (Fig. 2.36), whereas few-layer MoS2 also has additional peaks corresponding to the indirect gap luminescence, and direct gap hot luminescence [33]. Splendiani et al. [45] also performed PL measurements of mechanically exfoliated monolayer MoS2 films. They observed a strong PL at the direct excitonic transitions in a monolayer MoS2 films, whereas luminescence is absent in the indirect bandgap bulk MoS2 sample (Fig. 2.37). Eda et al. [106] study photoluminescence spectra of chemically exfoliated MoS2 thin films with average thicknesses ranging from 1.3 to 7.6 nm. They observed the thinnest samples exhibiting the strongest photoluminescence while the emission intensity gradually decreases with increasing film thickness, as shown in Fig. 2.38. The emission spectra for the thin films consist of one major 15.

(31) peak and one minor peak at around 660 and 610 nm, respectively. These peaks labeled as A and B, agree well with the energy of A and B excitons, suggesting that they are from the direct band gap photoluminescence from the K point. In 2012, Tongay et al. [119] reported the temperature dependence of photoluminescence spectra of single and few-layer mechanical exfoliated MoS2 and MoSe2 flakes onto 90 nm SiO2/Si substrate. At room-temperature, they observed the single-layer MoSe2 exhibits the strongest photoluminescence while the emission intensity gradually decreases with increasing layers, as shown in Fig. 2.39. In addition, Fig. 40a-d show the temperature dependence of PL measured on single- and few-layer samples of MoSe2 and MoS2 [119]. The temperature dependence of PL intensity of the single-layer and few-layer MoSe2 show striking difference. The PL intensity of MoS2 decrease with increasing temperature regardless of the layer thickness. In general, the suppression in PL intensity and peak broadening are typically attributed to the exponential enhancement in nonradiative electron−hole recombination processes, reducing the probability of radiative transition [119]. The distinct difference in the temperature behavior of these two materials is due to their intrinsic difference of band structures [119]. The rate of the indirect-to-direct bandgap crossover differs significantly between MoS2 and MoSe2. During this crossover the direct and indirect gaps in the case of single-layer and few-layer MoSe2 becomes nearly degenerate (the indirect bandgap value lies close to the direct bandgap). In contrast to MoSe2, the indirect and direct bandgaps are far from degenerate for MoS2 sample [119]. An increase in temperature slightly expands the interlayer distance. This tends to decouple neighboring MoSe2 layers, pushing the system further toward the bandgap degeneracy. In this case, the contribution from the hot PL across the direct bandgap to the PL intensity becomes much stronger at high temperatures. On the contrary, since the indirect and direct gaps are well16.