2-2-2 Optical properties

In 2013, H. R. Gutierrez et al. [64] presented Raman scattering spectra of triangular WS2

monolayers. The monolayer WS2 was deposited on SiO2/Si substrates by using thermal evaporation

process. Room-temperature Raman scattering spectra of bulk and monolayer WS₂ excided by 488 nm laser line show E and ¹_{2 g} A phonon modes of monolayer at 356 and 417.5 cm_1g ^-1 and E and ¹_{2 g}

A of bulk at 355.5 and 420.5 cm1g ^-1, as shown in Fig. 2.37. Figure 2.38 shows the thickness dependence of the frequency for the E and _{2 g}¹ A Raman modes. It can be observed that the _1g A _1g

mode redshifted and the E mode blueshifted are with decreasing number of layers. The van der ¹_{2 g}

Waals interaction between layers in bulk transition metal dichalcogenides stiffens the lattice, which is consistent with softening of the A mode in the single-layer. Similarly, anomalous behavior of _1g

the E mode is caused by stronger dielectric screening of the long-range Coulomb interactions ¹_{2 g}

between the effective charges in thicker samples.

In 2014, Thripuranthaka M. et al. [65] performed the temperature dependent phonon shifts of

single-layer WS2 fabricated by mechanically exfoliated. Figure 2.39 shows the typical Raman

scattering spectrum of single-layer WS2 at room-temperature excited by 514.5 nm laser. They observed 9 phonon modes. The first-order E and _{2 g}¹ A are approximately at 355 and 419 cm_1g ^-1. In

particular, the second-order ^2LA

( )

^M phonon mode, which is overlap with E¹_{2 g} mode, is approximately at 349 cm^-1.

Figure 2.40 shows the Raman scattering spectra of single-layer WS2 for 2LA

( )

M , E , and ¹_{2 g} A modes as a function of temperature from 77 to 623 K. They fitted the observed data of the peak 1g

positions versus temperature by using Gruneisen model. Therefore, Raman modes of atomic layer WS₂ ^2LA

( )

^{M ,} E , and ¹2 g A change linearly as a function of temperature variation as shown in _1g Fig. 2.41. The temperature coefficients of the frequencies of the ^2LA

( )

^{M ,} E , and ¹2 g A bands of _1g single-layer WS₂ are observed to be -0.008, -0.006, and -0.006 cm^-1 K^-1, respectively. The variation

in the phonon mode peak position as a function of temperature in single-layer WS2 is due to the

temperature contribution that consequences from anharmonicity and contribution from the thermal

expansion or volume contribution.

In the same year, E. D. Corro et al. [66] presented the Raman scattering spectra of 1L, 2L, 3L,

and bulk WSe₂ excited by different laser lines. The bulk WSe₂ crystals were produced using the

chemical vapor transport method. The few-layered WSe₂ was deposited on SiO₂/Si substrates by

mechanically exfoliated. Figure 2.42 shows selected resonant Raman spectra of WSe2 samples with

one to three layers, as well as bulk, collected with different laser lines. They found the relative

intensities of all Raman features are strongly dependent on the laser excitation as well as on the

number of layers. The decomposition of the registered spectra as a sum of individual peaks reveals

that none of the described peaks shifts with the excitation energy.

First-principle calculations predicted that transition metal dichalcogenides exhibit indirect to

direct band gap transformation with decreasing number of layers[64,67,68]. The electronic band

structures and density of states of bulk and monolayer WS2 and WSe2 calculated from first principles

are shown in Fig. 2.43 and 2.44. For bulk WS₂ and WSe₂, the electronic states involved in the

indirect transition originate from linear combinations of tungsten d-orbitals and sulfur or Selenium

p_z-orbitals. These electronic states exhibit a strong interlayer coupling, and their dispersion extremely

depends on the number of layers. For monolayer WS2 and WSe2, the indirect band gap between these

states is larger than direct transition at K, thus making the material a direct band gap semiconductor.

The conduction and valence band states at K are mainly due to tungsten d-orbitals, and their energies

are not very sensitive to the number of layers; the experimental direct band gap energy of WS2 and

WSe₂ at K symmetry point are approximately 2.01 and 1.56 eV.

The changes of the electronic band structures in WS₂ and WSe₂ manifest itself as a strong

photoluminescence feature. In 2013, H. Zeng et al. [69] performed the photoluminescence spectra of

few-layers and bulk WSe2. The few-layered WSe2 was deposited on SiO2/Si substrates by

mechanically exfoliated. Figure 2.45 illustrates the photoluminescence spectra of WSe2 samples with

various thickness. They found the PL intensity of bulk WSe2 is extremely weak, consistent with an

indirect band gap semiconductor. As WSe2 thin to a few atomic layers, the intensity PL from direct

interband transition dramatically increases and reaches maximum at monolayer, more than 3 orders

of magnitude stronger than that from bulk. A peak originating from the indirect band gap transition

(labeled as "I") gradually shifts toward higher energy and fades to null at monolayer. Moreover, the

peaks from indirect transition and the prominent direct transition peak (A), weak PL peak (B) was

observed at higher energy in WSe₂ at all thickness. Strikingly, the splitting between A and B peaks

are almost identical, approximately 0.4 eV for mono-, bi-, tri-, and quad-layer samples.

In the same year, H. R. Gutierrez et al. [64] presented the photoluminescence spectra of

triangular WS2 monolayers. The monolayer WS2 was deposited on SiO2/Si substrates by using

thermal evaporation process. Figure 2.46 displays the photoluminescence spectra of 1L, 2L, 3L, and

bulk WS2 excited by 488 nm laser line. They found for increasing number of layers, the indirect

transition between the local minimum of the conduction band states at the T point and local maximum of valence band states at the Γ point decreases in energy. The competition between direct

and indirect electronic transitions reduces the photoluminescence quantum efficiency and gives rise

to a new feature at longer wavelength when the WS2 film thickness increases to two or three layers.

The photoluminescence from bulk WS2 is very weak and only shows the indirect transition.

Fig. 2.1 Theoretical values of the Curie temperature for various p-type semiconductors containing 5 % of Mn and 3.5 × 10²⁰ holes per cm³ [7].

Fig. 2.2 X-ray diffraction pattern of Dy-doped ZnO nanowires [56].

Fig. 2.3 High-resolution x-ray diffractometer (HRXRD) patterns of Dy:ZnO thin films deposited under different Dy concentrations (at%) [20].

Fig. 2.4 Variation of resistivity of Dy:ZnO thin films with Dy concentration [20].

Fig. 2.5 Variation of mobility and carrier concentration of Dy:ZnO thin films with Dy concentration [20].

Fig. 2.6 XRD patterns of 0, 1, 2, and 3% Dy-doped ZnO synthesized by a sonochemical method over the 2θ range of 31 ~ 38^{ } [57].

Fig. 2.7 X-ray diffraction pattern of ZnO:Gd nanowires [58].

Fig. 2.8 Rutherford backscattering of Zn_0.95Gd_0.05O [58].

Fig. 2.9 The magnetization versus magnetic field of the as grown and annealed ZnO:Gd nanowires with Gd 5 mol% at 77 K [58].

Fig. 2.10 The magnetization versus magnetic field of the as grown and annealed ZnO:Gd nanowires with Gd 5 mol% at 300 K [58].

Fig. 2.11 (a) The magnetic susceptibility of the as grown and the annealed ZnO:Gd nanowires at room temperature under the applied magnetic field. (b) The effective magnetic moment per Gd atom

peff is a function of the applied magnetic field [58].

Fig. 2.12 XRD data for the ZnO(002) reflection of the 1.3, 4, 7, and 16 % Gd:ZnO films. The FWHM shown in the inset depends strongly on the amount of Gd in the films [59].

Fig. 2.13 XANES and XLD spectra at the Zn K-edge for the ion-implanted and the sputtered Gd:ZnO films [59].

Fig. 2.14 Isotropic (a) XANES, (b) XLD, and (c) XMCD of 1.3, 4, 7, and 16 % Gd:ZnO films at the Gd L₃-edge. The implanted sample is shown for comparison as a dotted line [59].

Fig. 2.15 Element-specific XMCD (H) curves at the Gd L3-edge for the ion-implanted Gd/cm² ZnO and the sputtered Gd:ZnO samples. The XMCD (H) was measured at 6.5 K [59].

Fig. 2.16 XMCD (H) curves taken at 6.4 and 41.5 K and SQUID M H at 5 K. The Brillouin

( )

function for J = =S 7 / 2 is shown as well [59].

Fig. 2.17 SQUID M T behavior recorded after field-cooled and zero-field-cooled conditions. The

( )

diamagnetic back-ground has been subtracted from all data sets [59].

Fig. 2.18 Relaxed structure of 3.7 % Gd at substitutional sites (a) and at interstitial sites (b) for a 3 × 3 ×3 cell of Gd:ZnO. The red atoms are O, gray atoms are Zn, and green atoms are Gd [59].

Fig. 2.19 Room temperature magnetic hysteresis loop of (a) pure and Gd doped ZnO, (b) 0.01, (c) 0.03, and (d) 0.05 mol%. The inset shows the enlarged view of hysteresis loops near the centre [60].

Fig. 2.20 Spectral dependence of transmittance of Dy:ZnO thin films and its inset in the figure shows

(

^{α υ vs. h}

)

²

( )

^hυ plot of Dy:ZnO thin film with 0.45 at% of Dy concentration [20].

Fig. 2.21 Variation of Dy:ZnO thin films band gap with Dy concentration [20].

Fig. 2.22 PL spectra of pure and Dy doped ZnO thin films excited at 325 nm [20].

Fig. 2.23 FTIR spectra of 0, 1, 2, and 3% Dy-doped ZnO synthesized by a sonochemical method [57].

Fig. 2.24 Decolorization efficiencies of MB by 0, 1, 2, and 3% Dy-doped ZnO during irradiation with UV light [57].

Fig. 2.25 Absorption spectra of 5, 10, and 15 mol% Gd doped ZnO nanocrystals (sample A, B, and C) measured at room temperature [21].

Fig. 2.26 PL spectra of the pure and 5 mol% Gd doped ZnO nanocrystals (sample A) at room temperature [21].

Fig. 2.27 PL spectra of ZnO:Gd nanocrystals with different mole ratios of Gd in the source materials:

(a) 5, (b) 10, and (c) 15 mol% Gd doped ZnO nanocrystals [21].

Fig. 2.28 Raman scattering spectra of ZnO samples prepared with different Gd amounts, (a) pure, (b) 0.01, (c) 0.03, and (d) 0.05 mol% [60].

Fig. 2.29 Deconvoluted 1LO mode from Raman scattering spectra of (a) pure and Gd doped ZnO nanoparticles (b) 0.01, (c) 0.03, and (d) 0.05 mol% [60].

Fig. 2.30 PL spectra of the pure and Gd doped ZnO, (a) 0.00, (b) 0.01, (c) 0.03, and (d) 0.05 mol% of Gd, measured at room temperature [60].

Fig. 2.31 Synthesis procedure for the atomic layer deposition based WS2 nanosheets [61].

Fig. 2.32 Field-effect transistor structure on the monolayer WS2 [61].

Fig. 2.33 Transfer curve for the field-effect transistor fabricated on a monolayer WS2 nanosheets [61].

Fig. 2.34 A schematic of the device with the principal layers shown [63].

Fig. 2.35 (Left axis) I-V curves for a device on Si/SiO2 taken under illumination at gate voltages from -20 (red) to +20 (blue) in 10-V steps, after doping. The laser illumination energy was 2.54 eV and the power was 10 μW. The curves are linear at low bias but saturate at higher bias due to limited available charge carriers. (Right axis) I-V curves for the same device taken in the dark at gate voltages from -20 (black) to +20 (green) in 20-V steps, after doping [63].

Fig. 2.36 (A) The density of states for monolayer transition metal dichalcogenides: MoS₂, WS₂, and WSe2. Strong peaks are present in all three materials that lead to a strong light-matter interaction. (B) The joint density of states with the same three transition metal dichalcogenides materials [63].

Fig. 2.37 Raman spectra of WS₂ bulk (dotted) and a monolayer (solid red). The inset shows the phonon vibration of E and ¹_{2 g} A modes [64]. _1g

Fig. 2.38 Frequencies of E and ¹_{2 g} A Raman modes (blue) and the difference in peak position _1g ω

∆ (red) as a function of number of WS2 layers [64].

Fig. 2.39 Room-temperature Raman spectra of single-layer WS₂ excited by 514 nm laser line [65].

Fig. 2.40 Raman spectra of single-layer WS₂ with (a) E and (b) _{2 g}¹ A mode measured in a _1g temperature range from 77 to 623 K [65].

Fig. 2.41 Effect of temperature variation on the Raman frequencies of monolayer WS₂ for (a)

( )

2LA M , (b) E , and (c) _{2 g}¹ A modes [65]. _1g

Fig. 2.42 Selected Raman scattering spectra of WSe₂ samples with different number of layers registered with a wide range of excitation wavelengths. The peak 520 cm^-1 comes from the silicon substrate and was used for intensity normalization [66].

Fig. 2.43 Electronic band structure (left) and total density of states (right) for the WS2 (a) bulk and (b) monolayer [64].

Fig. 2.44 Electronic band structure (left) and total density of states (right) for the WSe2 (a) bulk and (b) monolayer [68].

(a) (b)

(a) (b)

Fig. 2.45 (a) The relative photoluminescence intensity of WSe2 multilayer as a function of film thickness. The inset presents photoluminescence spectra from WSe₂ monolayer and bilayer respectively. (b) The normalized photoluminescence spectra (with respect to the peak A) of WSe2

ultrathin films. I labels the luminescence from indirect band gap transition; A and B label the direct band gap transition from the split valence band states edge to the conduction band states edge at K points. Spectra (dashed line) in the zoom windows have been multiplied by a factor as indicated for clarity [69].

Fig. 2.46 Photoluminescence intensities for 1L, 2L, 3L, and bulk using the 488 nm excitation laser line. The positions for the excitons A and B as well as the indirect band gap (I) are labeled [64].

(a) (b) (c)

Chapter 3

Experimental techniques

In this study, we investigate the lattice dynamics and interband transitions of (Dy, Gd) doped

ZnO and monolayer WS2 and WSe2 thin films by Raman scattering, optical transmission, and

spectroscopic ellipsometric spectra. Raman scattering spectra were measured in the range from 70 to

1555 cm^-1. Optical transmission spectra were measured in the range from 190 to 2600 nm (about 0.5

to 6.5 eV). Spectroscopic ellipsometric spectra were measured in the range from 193 to 1700 nm

(about 0.7 to 6.4 eV). In addition, the temperature dependent spectra were measured in the range

from 240 to 700 K by using spectroscopic ellipsometry.

在文檔中 摻雜鑭系元素(鏑,釓)氧化鋅與單層二(硫,硒)化鎢薄膜的光譜性質研究 (頁 49-78)