Plasma-assisted chemical vapor synthesis of transparent conducting oxides and their applications as transparent conducting films

Transparent conducting indium tin oxide (ITO), aluminum-doped zinc oxide (AZO), tin-doped zinc oxide (TZO) nanoparticles were synthesized by plasma-assisted chemical vapor synthesis. The injected precursors were vaporized in the plasma flame followed by vapor-phase reaction and subsequent quenching of the vaporized precursors produced nanosized powders. The electrical and optical properties of TCO films synthesized by spin coating a dispersion of nanoparticles vary as a function of dopant. Effects of plasma torch power and plasma-gas flow rate on product phases and grain size were investigated for ITO nanoparticles. The grain size increased with increasing plasma torch power and decreased with an increase in flow rate of plasma gas. The ITO gas sensor was exposed to different concentrations of H2 gas and temperatures to evaluate its gas sensitivity. The optimum operating temperature and gas concentration of H2 showing the highest sensitivity was determined to be 350 °C and 400 ppm, respectively. The linear relation between sensitivity and concentration up to 400 ppm of H2 can benefit the actuator to detect the concentration of H2, and thus making it suitable for high-performance hydrogen gas sensing applications. Synthesized TZO1 (3 atm % Sn) nanoparticles exhibited superior photocatalytic activity compared with ZnO and the improvement was ascribed to increase in specific surface area and enhanced oxygen vacancies as revealed from the XPS O 1s and PL spectra. Hall effect measurements showed that the minimum resistivity of 1.4 х 10-3 Ωcm was obtained for TZO1 film, and all the films exhibited an average transmission of 80 %, indicating their suitability in optoelectronic applications. Optical constants of the films were determined, which varied with doping amount. The photo-current properties of ZnO and TZO films were investigated, and only TZO1 film showed photo response property when irradiated with UV lamp. XRD results of AZO nanoparticles indicate the presence of wurtzite structure without any alumina peaks and SEM micrographs revealed spherical particles. The nanosized AZO would make an excellent material for use as photocatalyst due to high surface to volume ratio. The photocatalytic property of AZO was investigated using the degradation of methylene blue under ultraviolet irradiation. The effects of various parameters, such as catalyst amount, the presence of oxidant, temperature, bubbling of O2 gas, pH, specific surface area, oxygen vacancies, and initial concentration, were studied. The optical study showed that doping leads to a red-shift in band gap. Kinetic analyses indicated that the photodegradation of methylene blue followed pseudo-first order kinetic model using Langmuir-Hinshelwood (L-H) mechanism.

PLASMA-ASSISTED CHEMICAL VAPOR SYNTHESIS OF TRANSPARENT CONDUCTING OXIDES AND THEIR APPLICATIONS AS TRANSPARENT CONDUCTING FILMS by Arun Murali A thesis submitted to the faculty of The University of Utah in partial fulfillment of the requirements for the degree of Master of Science Department of Metallurgical Engineering The University of Utah August 2018 Copyright © Arun Murali 2018 All Rights Reserved The University of Utah Graduate School STATEMENT OF THESIS APPROVAL The thesis of Arun Murali has been approved by the following supervisory committee members: Hong Yong Sohn , Chair 4/24/2018 Date Approved Krista Carlson , Member 4/24/2018 Date Approved Prashant Sarswat , Member 4/24/2018 Date Approved and by Manoranjan Misra the Department/College/School of , Chair/Dean of Metallurgical Engineering and by David B. Kieda, Dean of The Graduate School. ABSTRACT Transparent conducting indium tin oxide (ITO), aluminum-doped zinc oxide (AZO), tin-doped zinc oxide (TZO) nanoparticles were synthesized by plasma-assisted chemical vapor synthesis. The injected precursors were vaporized in the plasma flame followed by vapor-phase reaction and subsequent quenching of the vaporized precursors produced nanosized powders. The electrical and optical properties of TCO films synthesized by spin coating a dispersion of nanoparticles vary as a function of dopant. Effects of plasma torch power and plasma-gas flow rate on product phases and grain size were investigated for ITO nanoparticles. The grain size increased with increasing plasma torch power and decreased with an increase in flow rate of plasma gas. The ITO gas sensor was exposed to different concentrations of H2 gas and temperatures to evaluate its gas sensitivity. The optimum operating temperature and gas concentration of H2 showing the highest sensitivity was determined to be 350 °C and 400 ppm, respectively. The linear relation between sensitivity and concentration up to 400 ppm of H2 can benefit the actuator to detect the concentration of H2, and thus making it suitable for high-performance hydrogen gas sensing applications. Synthesized TZO1 (3 atm % Sn) nanoparticles exhibited superior photocatalytic activity compared with ZnO and the improvement was ascribed to increase in specific surface area and enhanced oxygen vacancies as revealed from the XPS O 1s and PL spectra. Hall effect measurements showed that the minimum resistivity of 1.4 х 10-3 Ωcm was obtained for TZO1 film, and all the films exhibited an average transmission of 80 %, indicating their suitability in optoelectronic applications. Optical constants of the films were determined, which varied with doping amount. The photo-current properties of ZnO and TZO films were investigated, and only TZO1 film showed photo response property when irradiated with UV lamp. XRD results of AZO nanoparticles indicate the presence of wurtzite structure without any alumina peaks and SEM micrographs revealed spherical particles. The nanosized AZO would make an excellent material for use as photocatalyst due to high surface to volume ratio. The photocatalytic property of AZO was investigated using the degradation of methylene blue under ultraviolet irradiation. The effects of various parameters, such as catalyst amount, the presence of oxidant, temperature, bubbling of O2 gas, pH, specific surface area, oxygen vacancies, and initial concentration, were studied. The optical study showed that doping leads to a red-shift in band gap. Kinetic analyses indicated that the photodegradation of methylene blue followed pseudo-first order kinetic model using Langmuir-Hinshelwood (L-H) mechanism. iv This thesis is dedicated to my beloved parents who have always been a source of inspiration, encouragement and stamina to undertake my higher studies and to face the eventualities of life with zeal, enthusiasm, and love of God. TABLE OF CONTENTS ABSTRACT ....................................................................................................................... iii LIST OF TABLES ............................................................................................................. ix NOMENCLATURE ............................................................................................................x ABBREVIATIONS ........................................................................................................... xi ACKNOWLEDGMENTS ................................................................................................ xii Chapters 1. INTRODUCTION ...........................................................................................................1 1.1 Overview of Transparent Conducting Oxides (TCOs) .........................................1 1.2 Opto-electrical Properties of TCOs ......................................................................4 1.3 Industrial Application of TCOs ............................................................................9 1.4 Vapor-phase Synthesis of Nanoparticles ............................................................12 1.5 Recent Examples of Advances in Gas Phase Methods for Preparing Nanoparticles ............................................................................................................12 1.6 Description of Plasma Reactor Process ..............................................................15 1.7 Research Objectives ............................................................................................16 1.8 References ...........................................................................................................18 2. PLASMA ASSISTED CHEMICAL VAPOR SYNTHESIS OF INDIUM TIN OXIDE NANOPOWDER AND HYDROGEN SENSING PROPERTY OF ITO THIN FILM .......................................................................................................................25 2.1 Introduction .........................................................................................................25 2.2 Experimental Set-Up ...........................................................................................28 2.3 Experimental Procedure .....................................................................................29 2.3.1 Analysis Methods for Nanopowders .......................................................29 2.3.2 Preparation of Thin Films .......................................................................30 2.3.3 Analysis Methods for Thin Films ...........................................................30 2.4 Results and Discussions ......................................................................................31 2.4.1 Plasma Process Parameter.......................................................................31 2.4.1.1 Effect of plasma torch power .....................................................31 2.4.1.2 Effect of plasma flow rate at 10 kW ...........................................33 2.4.2 Raman Spectroscopy ...............................................................................33 2.4.3 X-ray Photoelectron Spectroscopy (XPS) ..............................................34 2.4.4 Scanning Electron Microscopy- Energy Dispersive Spectroscopy ........35 2.4.5 Thin Film Characterization Results ........................................................36 2.4.5.1 XRD analysis ..............................................................................36 2.4.5.2 Electrical properties ....................................................................36 2.4.5.3 Optical properties ........................................................................38 2.4.5.4 Gas sensitivity study ...................................................................40 2.5 Conclusions .........................................................................................................44 2.6 References ...........................................................................................................45 3. PHOTOCATALYTIC PROPERTY OF PLASMA-SYNTHESIZED ZINC OXIDE AND TIN-DOPED ZINC OXIDE (TZO) NANOPARTICLES AND THEIR APPLICATION AS TRANSPARENT CONDUCTING FILMS .....................................61 3.1 Introduction .........................................................................................................61 3.2 Experimental Set-Up ...........................................................................................66 3.3 Experimental Procedure .....................................................................................66 3.3.1 Analysis Methods for Nanopowders .......................................................67 3.3.2 Photocatalysis Set-Up .............................................................................68 3.3.3 Measurement of Photocatalytic Activity ................................................69 3.3.4 Preparation of Thin Films .......................................................................70 3.3.5 Analysis Methods for Thin films ............................................................71 3.4 Results and Discussions ......................................................................................71 3.4.1 X-ray Diffraction (XRD) ........................................................................71 3.4.2 Raman Spectroscopy ...............................................................................73 3.4.3 X-ray Photoelectron Spectroscopy (XPS) and Photoluminescence Spectroscopy (PL) ............................................................................................74 3.4.4 Scanning Electron Microscopy- Energy Dispersive Spectroscopy ........76 3.4.5 Photocatalysis Data and Absorption Spectrum Analysis ........................76 3.4.6 Thin Film Characterization Results ........................................................79 3.4.6.1 XRD analysis ..............................................................................79 3.4.6.2 Electrical properties ....................................................................79 3.4.6.3 Optical properties ........................................................................80 3.4.6.4 Photocurrent properties ...............................................................82 3.5 Figure of Merit ....................................................................................................83 3.6 Conclusions .........................................................................................................83 3.7 References ...........................................................................................................84 4. PLASMA- ASSISTED CHEMICAL VAPOR SYNTHESIS OF ALUMINUM- DOPED ZINC OXIDE (AZO) NANOPOWDER AND SYNTHESIS OF AZO FILMS FOR OPTOELECTRONIC APPLICATIONS .........................................................................105 4.1 Introduction .......................................................................................................107 4.2 Experimental Set-Up .........................................................................................107 4.3 Experimental Procedure ...................................................................................108 vii 4.3.1 Analysis Methods for Nanopowders .....................................................109 4.3.2 Preparation of Thin Films .....................................................................109 4.3.3 Analysis Methods for Thin Films .........................................................109 4.4 Results and Discussions ....................................................................................110 4.4.1 X-ray Diffraction (XRD) ......................................................................111 4.4.2 Scanning Electron Microscopy- Energy Dispersive Spectroscopy ......111 4.4.3 X-ray Photoelectron Spectroscopy (XPS) ............................................112 4.4.4 Raman Spectroscopy .............................................................................113 4.4.5 Magnetism Measurements of AZO Samples .......................................114 4.4.6 Thin Film Characterization Results ......................................................116 4.4.6.1 XRD analysis ............................................................................116 4.4.6.2 Electrical properties ..................................................................117 4.4.6.3 Optical properties ......................................................................119 4.5 Figure of Merit ..................................................................................................123 4.6 Conclusions .......................................................................................................123 4.7 References .........................................................................................................124 5. PHOTOCATALYTIC PROPERTY OF PLASMA-SYNTHESIZED ALUMINUM- DOPED ZINC OXIDE NANOPOWDER ..............................................143 5.1 Introduction .......................................................................................................143 5.2 AZO Synthesis and Characterization ................................................................146 5.2.1 Synthesis Procedure ..............................................................................147 5.2.2 Characterization ....................................................................................148 5.2.3 Photocatalysis Set-Up ...........................................................................148 5.2.4 Measurement of Photocatalytic Activity ..............................................149 5.3 Results and Discussions ....................................................................................151 5.3.1 X-ray Diffraction (XRD) ......................................................................151 5.3.2 Scanning Electron Microscopy- Energy Dispersive Spectroscopy ......152 5.3.3 Effect of Catalyst Amount ....................................................................153 5.3.4 Effect of H2O2 Addition ........................................................................154 5.3.5 Effect of Temperature ..........................................................................155 5.3.6 Effect of Bubbling O2 ...........................................................................156 5.3.7 Effect of pH...........................................................................................157 5.3.8 Effect of Oxygen Vacancies .................................................................158 5.3.9 Effect of Specific Surface Area ............................................................159 5.3.10 Effect of Initial Concentrations ...........................................................160 5.3.11 Effect of Doping .................................................................................161 5.4 Conclusions .......................................................................................................163 5.5 References .........................................................................................................164 6. CONCLUSIONS..........................................................................................................184 viii LIST OF TABLES Tables 1.1 TCO compounds and dopants .....................................................................................22 2.1 The electrical properties of ITO films .........................................................................50 3.1 Crystallite size calculated from XRD analysis ............................................................90 3.2 The electrical properties of ZnO and TZO films .........................................................90 3.3 Figure of merit values for ZnO and TZO films ...........................................................90 4.1 Crystallite size calculated from XRD analysis ..........................................................131 4.2 The electrical properties of AZO films ......................................................................131 4.3 Figure of merit values for AZO films ........................................................................132 NOMENCLATURE Symbol Definition σ Electrical conductivity n Carrier density e Electronic charge μ Mobility ρ Electrical resistivity α Absorption coefficient γ Absorption cross-section μi Ionized impurity scattering factor μg Grain boundary scattering factor Eg Band Gap h Planck's Constant ν Frequency of incident photon mc* Conduction band effective mass S Sensitivity k Reaction rate constant A Specific surface area Rs Sheet resistance ABBREVIATIONS AZO Aluminum Doped Zinc Oxide BM Burstein-Moss shift CB Conduction Band CBM Conduction Band Minimum DLE Deep Level Emissions FOM Figure of Merit FTO Fluorine Doped Zinc Oxide GZO Gallium Doped Zinc Oxide ITO Indium Tin Oxide L-H Langmuir-Hinshelwood LO Longitudinal modes MB Methylene Blue NBE Near-Band Edge Emission TCO Transparent Conducting Oxide TO Transverse modes TZO Tin- Doped Zinc Oxide VB Valence Band VBM Valence Band Maximum ACKNOWLEDGMENTS I would like to express my profound gratitude to my supervisor, Professor H.Y. Sohn whose expertise, understanding and patience added considerably to my graduate studies. I appreciate his tireless work ethic, positive life attitude, and more importantly his assistance in writing my first papers in addition to this thesis. I would like to thank my excellent supervisory committee members, Professor Krista Carlson and Professor Prashant Sarswat, for their priceless time and effort to review my work. I would like to thank Dr. Paulo Perez and Dr. Brian Van Devener with the University of Utah Nanofab for assistance with the SEM and XPS analyses. The financial support from NSF/U.S.-Egypt Joint Science and Technology Board under NSF Grant No. IIA-1445577 is gratefully acknowledged. Finally, I would like to also thank my family and friends, especially my parents for their unconditional love and support. I would not be where I am today without their consistent encouragement. CHAPTER 1 INTRODUCTION 1.1 Overview of Transparent Conducting Oxides (TCOs) A TCO is a wide band-gap semiconductor that has high concentration of free electrons in its conduction band. These arise either from defects in the material or from extrinsic dopants, the impurity levels of which lie near the conduction band edge. The high-electron-carrier concentration (the materials will be assumed to be n-type unless otherwise specified) causes absorption of electromagnetic radiation in both the visible and infrared portions of the spectrum for the present purposes, with the former being more important. Because a TCO must necessarily represent a compromise between electrical conductivity and optical transmittance, a careful balance between the properties is required. Reduction of the resistivity involves either an increase in the carrier concentration or in the mobility. Increasing the former also leads to an increase in the visible absorption. Increasing the mobility, however, has no deleterious effect and is probably the best direction to follow. To achieve a high carrier mobility will necessarily improve the optical properties [1]. Generally, TCOs are metal oxides with high optical transmittance and high electrical conductivity [2]. They are also referred to as wide-bandgap oxide semiconductors (band gap > 3.2eV). These materials have high optical transmission at 2 visible wavelengths (400 - 700 nm) and electrical conductivity close to that of metals, which is often induced by doping with other elements. They also reflect the near infrared and infrared wavelengths. Since the band gaps of these materials lie in the ultraviolet wavelength region they hardly absorb visible light; so, they appear to be transparent to the human eye. These unique properties make TCOs widely applicable in modern electronics, which requires optical access behind electrical circuitry. In order to be considered as a TCO substrate, the film needs to possess a low electrical resistivity (~103 - 10-4 Ωcm) as well as a high optical transparency towards visible light (> 80% transmittance) due to their wide band gap (>3.0 eV) [3]. Transparent conductive oxides (TCOs) are used in a wide range of applications including low-e windows, transparent contacts for solar cells, optoelectronic devices, flat panel displays, liquid crystal devices, touch screens, EMI shielding, and automobile window deicing and defogging. These materials have been intensely developed since the late 1970s but have actually been around for a century. Cadmium oxide (CdO) was the first TCO and was used in solar cells in the early 1900s. Tin oxide (SnO2) was first deposited on glass by pyrolysis and CVD in the 1940s for electroluminescent panels. Since then, applications and deposition processes have mushroomed. Transparent thin films and materials are some of the most commonly used materials that we depend on for a wide range of applications. A number of ternary compounds have been developed over the last 10 years including Zn2SnO4, ZnSnO3, MgIn2O4, (GaIn)2O3, Zn2In2O5, and In4Sn3O12. They are critical for energy efficiency applications such as low-e windows, solar cells, and electrochromic windows and are routinely deposited onto plastics and flexible plastics. Transparent conductive thin films are used as the transparent electrical 3 contacts in flat panel displays, sensors, and optical limiters. Newly developing applications are the charge carrier layers in transparent transistors and solar cells, dye sensitive solar cells and organic solar cells. TCOs are now used in such a wide variety of energy-related applications that the cost savings they provide is difficult to quantify. The energy savings in low-e coatings alone is staggering. These coatings can reduce energy loss by as much as 30%. TCOs are used as transparent conductive electrodes in virtually every thin solar cell design and several single crystal cells as well as electrochromic windows [4, 5]. The relative scarcity and high cost associated with indium are significant drawbacks for ITO; whilst fluorine-doped tin oxide (FTO) is far more cost effective but it requires higher deposition temperatures, which limits its applicability in flexible devices. Low cost, high durability and being nontoxic make ZnO an attractive alternative to the commonly used ITO. ZnO has a direct and wide band gap in the near-UV spectral region and a large free-exciton binding energy so that excitonic emission processes can persist at or even above room temperature. One of the key challenges in developing ZnObased TCOs is investigating the best metal dopants and the optimal dopant contents in order to achieve the highest electrical conductivity. Unlike in SnO2- and In2O3-based TCOs, efficient doping of group III elements into the ZnO structure could decrease the resistivity significantly, potentially realizing a future low cost TCO for electronic and optoelectronic applications [6]. In transparent conducting oxides (TCOs), the nonmetal part, B, consists of oxygen. In combination with different metals or metal-combinations, A, they lead to compound semiconductors, AyBz, with different opto-electrical characteristics. These 4 opto-electrical characteristics can be changed by doping, AyBz:D (D = dopant), with metals, metalloids or nonmetals. Hence, metals can be part of the compound semiconductor itself, A, or can be a dopant, D [7]. 1.2 Opto-electrical Properties of TCOs The electrical properties of TCOs can be understood by the semiconductor band theory. For electrical conduction to occur within a semiconductor material, ground state electrons must be excited from the valence band to the conduction band minimum (CBM), across the band gap by absorbing photon energy. A wider band gap requires a higher-energy photon in order for an electron to become excited into conduction. Therefore, widening the band gap (i.e, Eg > 3.0 eV) in a material permits transparency to the visible portion of the spectrum by placing a greater separation between the valance band maximum (VBM) and CBM of the material, thus decreasing the probability of exciting an electron into conduction [8]. TCOs have been developed by doping materials in order to facilitate the charge carrier generation within the structure. In the description of the band model, there is an important difference between the fundamental band gap (i.e, the energy separation of the Evb and Ecb; an intrinsic property of the material) and the optical band gap (an extrinsic property), which corresponds to the lowest-energy allowed for an optical transition. The optical band gap determines the transparency of a material which is important in TCO applications. In order to achieve n-type conducting properties, electrons are injected from a nearby defect donor level directly into the conduction band. The point defects in a metal oxide crystal, such as oxygen vacancies, proton or metal interstitials and certain substitutional defects, effectively create an excess 5 of electrons close to the defect site in n-type TCOs. If there is sufficient orbital overlap, it permits delocalization of electrons from the defect sites such that electronic states at the CBM become filled or in other words Fermi level shifts above the CBM. This leads to an effect known as the Moss-Burstein shift, which effectively widens the optical band gap. Eg = ECBM - EVBM (1.1) Eoptg = EMBg + Eg = EF - EVBM (1.2) Since the Moss-Burstein shift is as follows: EMBg = EF - ECBM (1.3) where Eg is the fundamental energy gap separating the VBM and CBM, Eoptg is the optical band gap corresponding to the smallest allowed optical transition from the VB to the CB, EMBg is the Moss-Burstein shift, and EF is the Fermi level shown in Figure 1.1. Thus, lattice defects in TCOs can simultaneously promote both electrical conductivity and optical transparency. Apart from the Moss-Burstein shift, the fundamental band gap is tapered due to the band gap narrowing effect that led by exchange interactions in the free-electron gas and electrostatic interactions between free electrons and ionized impurities [9]. The optical band gap is a key aspect in designing a TCO. However, the CBM depth or electron affinity (EA), in other words the difference between valence energy and CBM that affects the ‘dopability' of the TCO, is also equally important in 6 determining the conducting properties. A higher value of EA indicates greater ease of introducing charge carriers, i.e, a greater dopability [10]. A large separation (Eg > 3.0eV) between the Fermi level in the conduction band and the next electronic energy level (CBM+1) helps to prevent excitation of electrons to higher states within the conduction band, which prevents undesirable optical absorption [11]. The conductivity of a TCO is determined by the number of charge carriers and their mobility within the crystal lattice, which is inversely proportional to their effective mass. The effective mass is a quantity used to express the mass that the electrons appear to have when moving within a crystalline solid in which their mobility is affected by their response to local forces within the crystal, expressed relative to their true mass (me). The local forces in TCO crystal lattice is controlled by the orbital overlap between the metal cation in a host lattice and the oxygen. The electron mobility, electron density and conductivity of inorganic materials is linearly related as described in the Boltzmann equation: σ = neμ (1.4) where σ is the electrical conductivity defined in S cm−1, n is the density of free charge carriers (i.e, electrons in an n-type TCO), e is the electronic charge, and μ is the electron mobility. The electrical resistivity (ρ) is expressed in Ω.cm, as follows: ρ=1/σ (1.5) 7 There are three distinct domains amongst inorganic materials regarding electrical properties: namely the semimetals (high carrier density, low electron mobility), highly conductive metals (both high carrier density and mobility), and semiconductors (low carrier density, high mobility). While the introduction of a donor level close to the conduction band permits a wide optical band gap from the VBM to the CBM, the optical absorption associated with the promotion of electrons from the CBM to higher states places an upper limit on the carrier concentration in the CBM, such that the absorption coefficient α of the TCO is proportional to the density of free electrons n as shown in Eq. 6: α = γn (1.6) where γ is the absorption cross-section, and n is the carrier density. The conductivity  is intrinsically limited for two reasons. First, n and  cannot be independently increased for practical TCOs with relatively high carrier concentrations. At a high conducting electron density, carrier transport is limited primarily by ionized impurity scattering, which is caused by Coulomb interactions between electrons and the dopants. Higher doping concentration reduces carrier mobility to a degree that the conductivity is not increased, and it decreases the optical transmission at the near-infrared edge. With increasing dopant concentration, the resistivity reaches a lower limit and does not decrease beyond it, whereas the optical window becomes narrower. Bellingham et al. [12] were the first to report that the mobility and hence the resistivity of transparent conductive oxides (ITO, SnO2, ZnO) are limited by ionized impurity scattering for carrier concentrations above 8 1020 cm-3. Scattering by the ionized dopant atoms that are homogeneously distributed in the semiconductor is only one of the possible effects that reduce the mobility. In addition to the above-mentioned effects that limit the conductivity, high dopant concentration could lead to clustering of the dopant ions, which increases significantly the scattering rate, and it could also produce nonparabolicity of the conduction band, which has to be taken into account for degenerately doped semiconductors with filled conduction bands. While the development of new TCO materials is mostly dictated by the requirements of specific applications, low resistivity and low optical absorption are always significant prerequisites. There are basically two strategies in managing the task of developing advanced TCOs that could satisfy the requirements. The main strategy dopes known is binary TCOs with other elements, which can increase the density of conducting electrons. As shown in Table 1.1, more than 20 different doped binary TCOs were produced and characterized, of which ITO was preferred, while AZO and gallium doped zinc oxide (GZO) came close to it in their electrical and optical performance. Doping with low metallic ion concentration generates shallow donor levels, forming a carrier population at room temperature. Doping In2O3 with Sn to form ITO substantially increases conductivity. It is believed that substituting Sn4+ for In3+ provides carrier electrons, as Sn4+ is supposed to act as a one-electron donor. Similarly, aluminum is often used for intentional n-type doping of ZnO, but other group III impurities, such as Ga and In, and group IV, such as Sn and Ge, also work. Doping by Al produces AZO with relatively high conductivity. Doping with nonmetallic elements is also common, e.g., ZnO:Ge, SnO2:F and SnO2:Sb. 9 1.3 Industrial Application of TCOs TCOs have diverse industrial applications - some of the more important ones will be described in this section. TCO coatings are applied to transparent materials used for work surfaces and closet doors, particularly in clean rooms used for electronics assembly, in order to prevent harmful static charge buildup. In this application, relatively high surface resistances can be tolerated. Transparent heating elements may be constructed from TCO coatings. These are applied as defrosters in aircraft and vehicular windshields. Their advantages over traditional hot air blowers are that they can have a much shorter effective defrosting time, and work uniformly over large areas. This application requires either the use of very low surface resistance coatings or a high voltage power source. TCO coatings may be used as shielding to decrease electromagnetic radiation interference (EMI) from providing visual access. This may be to keep radiation from escaping an enclosure and interfering with nearby devices. Such coatings will also prevent external radiation from entering an enclosure and interfering with electronic devices within. A potential example is the window of domestic microwave ovens, which today use a perforated metal screen to reduce microwave leakage, which obscures clear visual observation. Radiation leakage must be minimized to prevent harm to the users as well as interference to proliferating wireless devices that use the unlicensed spectral band at 2.45 GHz. While transparent conducting films were proposed 50 years ago, an attempt to introduce microwave windows with TCO coatings into the market was not successful about a decade ago due to the high cost. Low cost designs are currently being developed. 10 The three largest applications of transparent conductive oxide thin films, in terms of the surface area covered and their total value, are flat panel displays, solar cells, and coatings on architectural glass. In general, transparent electrodes are needed for a large variety of electro-optical devices, of which flat panel displays and solar cells are the most important examples. In liquid crystal displays (LCDs), TCO films are needed for both electrodes, in order to allow backlighting to pass through the liquid crystal film while applying voltage to the various pixels. Generally, these electrodes are in the form of a pattern of lines, with the alignment of the lines on the two electrodes perpendicular to each other. This allows addressing individual pixels by applying a voltage to the two lines, which intersect at a given pixel. Thus, patterning the films is required. ITO is the TCO of choice in this application, both because of its electro-optical properties and the relative ease of acid etching. The best LCDs utilize an active matrix comprising one amorphous silicon transistor that occupies a corner of each pixel and, because the silicon is opaque, has reduced light transmission. Recently transparent field effect transistors (FETs) have been developed based on the zinc oxide but using a Cr gate. These zinc oxide FETs have been incorporated into small 220 х 280 10 m pixel active matrix LCDs. Small and medium LCDs are a 25 B$/yr market, which is growing by about 5 %/yr, while large area LCDs have a similar market size and a much higher growth. The explosive growth of demand for ITO coatings for this specific application has generated wide spread concern about indium scarcity in the near future. Most solar cells use TCO films as a transparent electrode. Major considerations in the choice of the TCO for this application, besides the conductivity and transparency, 11 are electronic compatibility with adjacent layers in the cell, processing requirements, and stability under environmental conditions. Often tin oxide-based films are chosen for this application, in as much as patterning is not required, but environmental stability is. TCO films are commonly applied to architectural glass, often as part of multilayer stacks. In window glass applications, usually the conductivity per se is irrelevant, but rather the concurrent high infra-red reflectivity is exploited in order to obtain good light transmission in the visible range, while minimizing heat transmission. This feature is used to minimize air conditioning costs in the summer and heating costs in the winter in buildings equipped with appropriately coated windows. Approximately 25% of flat glass is coated, and energy conserving coatings are now mandated in various regions. Most commonly the coatings are applied by two techniques: (1) Very enduring and inexpensive, but simple, coatings are produced with atmospheric pressure chemical vapor deposition (APCVD), in line with the float glass production process. This ensures a fresh surface and exploits the high temperature of the glass during its production. However, APCVD is not very flexible and there are only limited options available for the coating architecture. (2) A more flexible, but also costlier, process is magnetron sputtering. Commonly multiple (e.g, 20-60) rotary targets are mounted in long modular vacuum systems (e.g, 40-160 m in length), and multilayer stacks are deposited as the glass panels pass beneath the various cathodes, traveling at velocities of ~1 m/s. Typically these systems operate continuously for 2 weeks, after which expended targets are replaced and other maintenance is performed [13 - 15]. 12 1.4 Vapor-phase Synthesis of Nanoparticles In vapor-phase synthesis of nanoparticles, conditions are created where the vapor phase mixture is thermodynamically unstable relative to formation of the solid material to be prepared in nanoparticulate form. This includes usual situation of a supersaturated vapor. It also includes what we might call ‘chemical supersaturation' in which it is thermodynamically favorable for the vapor phase molecules to react chemically to form a condensed phase. If the degree of supersaturation is sufficient, and the reaction/ condensation kinetics permit, particles will nucleate homogeneously. Once nucleation occurs, remaining supersaturation can be relieved by condensation or reaction of the vapor-phase molecules on the resulting particles, and particle growth will occur rather than further nucleation. Therefore, to prepare small particles, one wants to create a high degree of supersaturation, thereby inducing a high nucleation density, and then immediately quench the system, either by removing the source of supersaturation or slowing the kinetics, so that the particles do not grow [16]. 1.5 Recent Examples of Advances in Gas-Phase Methods for Preparing Nanoparticles Inert gas condensation involves vaporizing the solid into a background gas then mixing the vapor with a cold gas to lower the temperature [17]. This method is well suited for production of metal nanoparticles, since many metals vaporize at reasonable rates at attainable temperatures. By including a reactive gas such as oxygen in the cold gas stream, oxides or other compounds of the vaporized materials can also be prepared. Laser ablation uses a laser to vaporize solid materials in an inert flow reactor [18]. 13 This method is good for precursors that have high volatilization temperatures and low vapor pressures such as refractory oxides. The main drawback to this method is that the amount of nanopowder produced is relatively small. Spark discharge generation involves using electrodes made of the metal to be vaporized in the presence of an inert background gas until the breakdown voltage is reached. The arc formed across the electrodes vaporizes a small amount of metal. This produces very small amounts of nanoparticle with relatively reproducible results. Ion sputtering involves the use of high velocity ions of an inert gas to vaporize the solid. The advantage of this method is that the composition of the sputtered material is the same as that of the target. The disadvantage of this method is that low pressures are required, which makes further processing of the nanoparticles in aerosol form difficult. Spray conversion processing involves coordinated steps: preparation and mixing of an appropriate starting solution, spray drying to form a chemically homogeneous precursor powder, and conversion of the precursor into the desired nanophase powder. This process has been applied to the production of nanosized composite powders [19]. It is a scalable technology, which provides the means for producing bulk quantities of nanophase powders at low manufacturing costs. Flame Synthesis involves synthesis of particles within a flame so that the heat needed is produced in situ by the combustion reactions. Lee et al. [20] controlled nanoparticle morphology by using a CO2 laser to reheat flame-synthesized titania and silica nanoparticle agglomerates, thereby sintering them in situ. Wegner et al. [21] controlled the size and morphology of titania nanoparticles by extracting the from the flame through a critical flow nozzle, quenching particle growth and agglomeration. 14 Kammler et al. [22] have shown that they can influence flame conditions and primary particle size by the application of a DC electric field to the particle synthesis flame. Flame spray pyrolysis involves allows the use of precursors that do not have sufficiently high vapor pressure to be delivered as a vapor [23-25]. Madler et al. [26] presented a very detailed study of this method as applied to the synthesis of silica particles from hexamethyl- disiloxane. Low-temperature reactive synthesis makes it is possible to react vapor phase precursors directly without external addition of heat, and without significant production of heat. Sarigiannis et al. [27] produced ZnSe nanoparticles from dimethylzinctrimethylamine and hydrogen selenide by mixing them in a counter-flow jet reactor at room temperature. Laser pyrolysis synthesis involves an alternate means of heating the precursors to induce reaction and homogeneous nucleation is absorption of laser energy. Compared to heating the gases in a furnace, this allows highly localized heating and rapid cooling, since only the gas (or a portion of the gas) is heated, and its heat capacity is small. Heating is generally done using an infrared (CO2) laser, whose energy is either absorbed by one of the precursors or by an inert photosensitizer such as sulfur hexafluoride. A few recent examples are MoS2 nanoparticles prepared by Borsella et al. [28], SiC nanoparticles produced by Kamlag et al. [29], and Si nanoparticles prepared by Ledoux et al. [30]. Ledoux et al. use a pulsed CO2 laser, thereby shortening the reaction time and allowing preparation of even smaller particles. Apparently, the heat of reaction was sufficient to allow crystallization of the particles without increasing the gas temperature. This is an intriguing result, because it is one of few methods reported for vapor-phase preparation 15 of compound semiconductor nanoparticles that are usually produced by colloidal chemistry. Chemical vapor synthesis involves introducing vapor phase precursors into a heated reactor and allowing vapor phase nucleation to occur in the reactor rather than depositing the product as a thin film. The precursors can be solid or liquid in nature but are introduced into the reactor as a vapor by sublimation or evaporation [31]. Metal chlorides are generally preferred as the precursors owing to their low volatilizing temperature. Chemical vapor synthesis has been applied to the synthesis of metallic [32] and intermetallic and alloy powders [33, 34], carbides [35] and their composites [36], and oxide nanopowders [37]. Chemical vapor synthesis assisted by thermal plasma involves the use of a plasma flame as the heat source [38] as shown in Figure 1.2. Plasma flames can facilitate vapor phase reactions by providing sufficient energy for vaporizing precursors and subsequent chemical reactions. The temperature of plasma flame generated is high enough to decompose even reactants of high vaporization temperatures into atoms and radicals, which can then react and condense to form nanosized particles when cooled by mixing with cool gas or expansion through a nozzle. 1.6 Description of the Plasma Reactor Process The plasma reactor system used for the synthesis of aluminum-doped zinc oxide (AZO) nanopowder consisted of a downward plasma torch, a power supply unit, a cylindrical reactor, a powder feeding system, cooling system, powder delivery system, powder collectors, a gas delivery system, an off-gas scrubber, and an off-gas exhaust system, as shown in Figure 1.3. The plasma torch had a copper cathode and a tungsten 16 anode. It was water-cooled and operated at atmospheric pressure. The reactor assembly consisted of a vertical stainless-steel tube, which was water-cooled, and had the dimensions of 15 cm inner diameter and 60 cm length. A graphite cylinder of 7.6 cm inner diameter and 60 cm length was placed inside the stainless-steel tube. A graphite felt filled the gap between the graphite cylinder and the stainless-steel tube. The cooling chamber consisted of two-layer stainless steel to cool the outgoing gases to lower than 150 °C. The precursor was directly fed into the plasma gun using Ar as the carrier gas through a powder feeding system consisting of a test tube filled with the precursor powder, a motor that pushed up the test tube at a constant rate, a carrier gas line that carried the fluidized particles from the top of the particle bed in the test tube, and a vibrator. The product was collected on a Teflon coated polyester filter with a pore size of 1 µm. The off-gas scrubber used a 5 % NaOH solution. 1.7 Research Objectives An innovative plasma processing technique has been developed for the preparation of nano-sized TZO powder by vapor phase reactions. Thermal plasma provides a high processing rate as well as other advantages like good control over size, shape and crystal structure as well as a clean reaction atmosphere that yields high purity products, a high quench rate to form ultra-fine powder, and a wide choice of reactants. Compared with other methods it avoids multiple steps like in mechanical milling, sol gel method, and precipitation method, and it does not require a high liquid volume and surfactants that are involved in a wet chemical process. 17 Thin films of TCOs on glass substrates have been prepared by a variety of physical and chemical deposition techniques for transparent and conductive electrodes, albeit with inherently costly and time-consuming process. Over other techniques, fabricating TCO films from well dispersed nanoparticles is a good alternative because it has the advantages of low processing cost and easy control of composition ratio. In addition to a cheaper processing technique, nanoparticles could be handled efficiently, and loss of raw materials could be minimized. In the present study, the following were performed: ITO nanoparticles with different Sn doping amount were synthesized by plasmaassisted chemical vapor synthesis. Effects of plasma torch power and plasma flow rate on product phases and grain size were investigated. ITO films were made by spin coating a dispersion of nano-sized ITO nanoparticles to investigate their electrical, optical and H2 gas sensing properties and elucidate the feasibility of ITO gas sensor. Zinc oxide and Tin-doped Zinc oxide nanoparticles were synthesized by the plasma process and tested for its photocatalytic property in the degradation of methylene blue. The obtained nanopowders were also used to fabricate ZnO and TZO films to study the effect of Sn doping on their structural, optical, electrical and photo response properties. Aluminum doped Zinc Oxide (AZO) nanopowders were successfully synthesized for the first time by the plasma process using zinc nitrate and aluminum nitrate as precursors. AZO nanoparticles were dispersed in organic solvent to produce transparent conducting layers on glass substrates for optoelectronic applications. The photocatalytic property of AZO was investigated using the degradation of methylene blue under ultraviolet irradiation. The effects of various parameters, such as 18 catalyst amount, the presence of oxidant, temperature, bubbling of O2 gas, pH, specific surface area, oxygen vacancies, and initial concentration were studied. 1.8 References 1. T. J. Coutts, D. L. Young, and X. Li, "Fundamental advances in transparent conducting oxides," in MRS Proc.,2000, vol. 623, pp. 199-209. 2. C. G. Granqvist and A. Hultaker, "Transparent and conducting ITO films: new developments and applications," Thin Solid Films, vol. 411, no. 1, pp. 1-5, 2002. 3. D.S.Y. Jayathilake, T.A. 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Kim, "Processing and properties of nanostructured WC-Co," Nanostructured Materials, vol. 1, no. 2, pp. 119-124, 1992. 20. D. Lee and M. Choi, "Coalescence enhanced synthesis of nanoparticles to control size, morphology and crystalline phase at high concentrations," J. Aerosol Sci., vol. 33, no. 1, pp. 1-16, 2002. 21. K. Wegner, W. J. Stark, and S. E. Pratsinis, "Flame-nozzle synthesis of nanoparticles with closely controlled size, morphology and crystallinity," Mater. Lett., vol. 55, no. 5, pp. 318-321, 2002. 22. H. K. Kammler, S. E. Pratsinis, P. W. Morrison, and B. Hemmerling, "Flame temperature measurements during electrically assisted aerosol synthesis of nanoparticles," Combust. Flame, vol. 128, no. 4, pp. 369-381, 2002. 23. Y. Ji, H. Y. Sohn, H. D. Jang, B. Wan, and T. A. Ring, "Computational fluid dynamic modeling of a flame reaction process for silica nanopowder synthesis from tetraethylorthosilicate," J. Am. Ceram. Soc., vol. 90, no. 12, pp. 3838-3845, 2007. 24. K. Cho, H. Chang, D. S. Kil, J. Park, H. D. Jang, and H. Y. Sohn, "Mechanisms of the formation of silica particles from precursors with different volatilities by flame spray pyrolysis," Aerosol Sci. Technol., vol. 43, no. 9, pp. 911-920, 2009. 20 25. M. Olivas-Martinez, H. Y. Sohn, H. D. Jang, and K.-I. Rhee, "Computational fluid dynamic modeling of the flame spray pyrolysis process for silica nanopowder synthesis," J. Nanopart. Res., vol. 17, no. 7, 2015. 26. L. Madler, H. Kammler, R. Mueller, and S. Pratsinis, "Controlled synthesis of nanostructured particles by flame spray pyrolysis," J. Aerosol Sci., vol. 33, no. 2, pp. 369- 389, 2002. 27. D. Sarigiannis, J. D. Peck, G. Kioseoglou, A. Petrou, and T. J. Mountziaris, "Characterization of vapor-phase-grown ZnSe nanoparticles," Appl. Phys. Lett., vol. 80, no. 21, pp. 4024-4026, 2002. 28. E. Borsella, S. Botti, M. Cesile, S. Martelli, and A. Nesterenko, "X-Ray diffraction investigation on MoS2 nanoparticles produced by CO2 laser assisted synthesis," Mater. Sci. Forum, vol. 278-281, pp. 636-641, 1998. 29. Y. Kamlag, A. Goossens, I. Colbeck, and J. Schoonman, "Formation of cubic SiC nanocrystals by laser-assisted CVD," Chem. Vap. Dep., vol. 11, no. PR3, 2001. 30. G. Ledoux, J. Gong, F. Huisken, O. Guillois, and C. Reynaud, "Photoluminescence of size-separated silicon nanocrystals: Confirmation of quantum confinement," Appl. Phys. Lett., vol. 80, no. 25, pp. 4834-4836, 2002. 31. H. Y. Sohn, Chemical vapor synthesis of inorganic nanopowders. New York: Nova Science Publishers, 2012. 32. T. Ryu, H. Y. Sohn, K. S. Hwang, and Z. Z. Fang, "Chemical vapor synthesis (CVS) of tungsten nanopowder in a thermal plasma reactor," Int. J. Refract. Hard Met., vol. 27, no. 1, pp. 149-154, 2009. 33. H. Y. Sohn and S. Paldey, "Synthesis of ultrafine particles of intermetallic compounds by the vapor-phase magnesium reduction of chloride mixtures: Part II. Nickel aluminides," Metall. Mater. Trans. B, vol. 29, no. 2, pp. 465-469, 1998. 34. H. Y. Sohn and S. Paldey, "Synthesis of ultrafine particles and thin films of Ni4Mo by the vapor-phase hydrogen co-reduction of the constituent metal chlorides," Mater. Sci. Eng. A, vol. 247, no. 1-2, pp. 165-172, 1998. 35. T. Ryu, H. Y. Sohn, K. S. Hwang, and Z. Z. Fang, "Tungsten carbide nanopowder by plasma-assisted chemical vapor synthesis from WCl6-CH4-H2 mixtures," J. Mater. Sci., vol. 43, no. 15, pp. 5185-5192, 2008. 36. T. Ryu, H. Y. Sohn, G. Han, Y.-U. Kim, K. S. Hwang, M. Mena, and Z. Z. Fang, "Nanograined WC-Co composite powders by chemical vapor synthesis," Metall. Mater. Trans. B, vol. 39, no. 1, pp. 1-6, 2008. 21 37. T. Ryu, Y. J. Choi, S. Hwang, H. Y. Sohn, and I. Kim, "Synthesis of yttriastabilized zirconia nanopowders by a thermal plasma process," J. Am. Ceram. Soc., vol. 93, no. 10, pp. 3130-3135, 2010. 38. J. H. Seo and B. G. Hong, "Thermal plasma synthesis of nano-sized Powders," Nucl. Eng. Technol., vol. 44, no. 1, pp. 9-20, 2012. 22 Table 1.1 TCO compounds and dopants TCO Dopant SnO2 Sb, F, As, Nb, Ta Al, Ga, B, In, Y, Sn, F, V, Si, Ge,Ti, ZnO Zr, Hf, Mg, As, H In2O3 Sn, Mo,Ta, W, Zr, F, Ge, Nb, Hf, Mg CdO In, Sn Ta2O GaInO3 Sn, Ge CdSb2O3 Y 23 Figure 1.1 Optical widening by the effect of the Moss-Burstein shift. Figure 1.2 Plasma Synthesis Procedure for the Production of Nano-sized Powders. 24 Figure 1.3 Schematic diagram of plasma reactor system: (1) powder feeding s system, (2) plasma gun, (3) reactor chamber, (4) cooling chamber, (5) powder collector, and (6) scrubber. CHAPTER 2 PLASMA-ASSISTED CHEMICAL VAPOR SYNTHESIS OF INDIUM TIN OXIDE (ITO) NANOPOWDER AND HYDROGEN-SENSING PROPERTY OF ITO THIN FILM 2.1 Introduction ITO is an n-type transparent conducting oxide (TCO) formed by substitutional doping of In2O3 in which In+3 sites are substituted with Sn+4 sites. It is used in various optoelectronic applications because it combines high conductivity with optical transparency in the visible region and has attracted a great deal of attention due to its technological application. It is mainly used to make transparent conducting coatings in electronic displays and heat-reflective coatings for architectural, automotive and light bulb glass. ITO is commonly used in TCO's owing to its excellent performance with extensive commercial applications as electrodes in flat-panel displays, electrochromic windows, and solar cells, in addition to being used as heat mirrors for reducing energy consumption due to its high reflectance in the infrared region and as gas sensors [1 - 3]. An innovative plasma processing technique has been developed for the preparation of nano-sized powder by vapor phase reactions. Thermal plasma provides a high processing rate as well as other advantages like good control over size, shape and crystal structure as well as a clean reaction atmosphere that yields high purity products, 26 a high quench rate to form ultra-fine powder, and a wide choice of reactants [4 - 6]. Compared with other methods, it avoids multiple steps like in mechanical milling, solgel method and precipitation method and does not require a high liquid volume and surfactants that are involved in a wet chemical process. Precursors like chloride salts are commonly used in wet chemical methods but the presence of chloride ions form hard agglomerates in oxide particles and also difficult to be rinsed from the colloidal hydroxide precipitate [7]. To avoid contamination, indium and tin alkoxides have been used but the precursors are expensive, water sensitive and difficult to control the synthesis process [8]. Thin films of ITO on glass substrates have been made by a variety of techniques such as chemical vapor deposition [9], spray pyrolysis [10], direct thermal evaporation of indium and tin metal [11], vacuum evaporation [12], magnetron sputtering [13] and electron-beam evaporation [14] for transparent and conductive electrodes, albeit with inherently costly and time-consuming processes. Over other techniques, fabricating ITO films from well dispersed ITO nanoparticles is a good alternative because it has the advantages of low processing cost and easy control of composition ratio. In addition to cheaper processing technique, nanoparticles could be handled efficiently and the loss of raw materials, especially the scarce and costly indium metal, could be minimized [15, 16]. In recent years, research interest in hydrogen as a near future fuel has increased because it is renewable, abundant, efficient, and unlike other alternatives, provides few emissions. Hydrogen can be used either as a fuel for direct combustion or for fuel cells, which in turn generate direct current electricity to power an electric motor. The amount of energy produced by hydrogen, per unit weight of fuel, is about three times the energy contained in an equal weight of gasoline and nearly seven times that of coal [17]. 27 However, the storage and use of hydrogen poses unique challenges due to its ease of leaking, low-energy ignition, a wide range of combustible fuel-air mixture ratio, buoyancy, and its ability to embrittle metals that must be accounted for to ensure safe operation [18]. Much of the previously published reports extensively deal with nanostructured ZnO-based sensor for gas-sensing applications. Doping with noble metals (such as Pt, Pd) or other alternatives (such as Al, Cu) improved the selectivity and sensitivity of ZnO-based sensor and was used to detect several gases. Recent research has shown that transparent and conducting ITO films offer new advantages in the design of metal-oxide based sensor. Patel et al. [19, 20] have reported the detection of methanol, CO2 and CCl4 using ITO as thin film sensor. Sberveglieri et al. [21] reported on RF sputtered ITO films for NO2 detection and reaction sputtered ITO films for NO gas detection [22]. Galkidas et al. [23] investigated the chlorine- gas sensing property of the ITO film. An ITO based quartz microbalance NO gas sensor has been reported by Zhang et al. [24]. Investigations of nano-grained thin film ITO catalyzed with 0.5 wt % Pd as gas sensors for H2 detection in the range of 1000 - 5000 ppm were carried out by Yoo et al. [25]. The reported sensitivity of 0.008 for 1000 ppm was observed at the operating temperature of 300 °C. In this present study, an ITO gas sensor was fabricated without the addition of catalytic metals to detect the H2 gas in the low concentration range of 50 - 600 ppm that ensures high sensitivity, fast response and recovery time. Kim et al. [26] previously showed the preparation of ITO particles by DC thermal plasma but the effects of plasma torch power and plasma flow rate were not investigated. Also, the fabrication of ITO films from plasma-synthesized nanoparticles has not been explored. In this work, ITO nanoparticles were successfully synthesized by a plasma 28 process using indium nitrate and tin nitrate as the precursors, and the effects of plasma torch power and plasma gas flow rate on material properties were investigated. ITO films were made by spin coating a dispersion of nano-sized ITO nanoparticles to investigate their electrical, optical and H2 gas sensing properties and elucidate the feasibility of an ITO gas sensor. 2.2 Experimental Set-Up The plasma reactor system used for the synthesis of indium tin oxide (ITO) nanopowder consisted of a downward plasma torch, a power supply unit, a cylindrical reactor, a powder feeding system, cooling system, powder delivery system, powder collectors, a gas delivery system, an off-gas scrubber, and an off-gas exhaust system. The plasma torch had a copper cathode and a tungsten anode. It was water-cooled and operated at atmospheric pressure. The reactor assembly consisted of a vertical stainlesssteel tube, which was water-cooled, and had the dimensions of 15 cm inner diameter and 60 cm length. A graphite cylinder of 7.6 cm inner diameter and 60 cm length was placed inside the stainless-steel tube. A graphite felt filled the gap between the graphite cylinder and the stainless-steel tube. The cooling chamber consisted of two-layer stainless steel to cool the outgoing gases to lower than 150 °C. The precursor was directly fed into the plasma gun using Ar as the carrier gas through a powder feeding system consisting of a test tube filled with the precursor powder, a motor that pushed up the test tube at a constant rate, a carrier gas line that carried the fluidized particles from the top of the particle bed in the test tube, and a vibrator. The product was collected on a Teflon coated polyester filter with a pore size of 1 µm. The off-gas scrubber used a 5 % NaOH solution. 29 More details about the experimental set-up can be found in a previous publication [4 - 6, 27]. 2.3 Experimental Procedure The precursors used in this work were 1) indium nitrate (In(NO₃)3•5H₂O, Alfa Aesar, Haverhill, MA) and 2) tin nitrate(Sn(NO₃)4, American Elements, Los Angeles, CA). Each precursor was ground by mortar and pestle and sieved until the final size became approximately 50 µm. The milled precursors were kept in a vacuum oven at 50 °C to remove moisture for easy feeding. The anhydrous milled precursors were uniformly mixed using a vibrating mixer. The uniformly mixed precursor was fed into the plasma flame through an internal port of 2 mm diameter in the plasma torch, using the specially designed powder feeder system mentioned above. The experimental conditions used, unless otherwise stated, were as follows: 1) plasma gas flow rate (Ar) of 30 L/min (25 °C and 86.1 kPa total pressure at Salt Lake City), 2) precursor feeding rate of 0.5 g/min, 3) flow rate of carrier gas (Ar) of 3.5 L/min, and 4) plasma power of 10 kW. The amount of tin nitrate was varied to obtain In/Sn atomic ratios of 95:5, 90:10 and 85:15 designated as ITO1, ITO2 and ITO3, respectively. 2.3.1 Analysis Methods for Nanopowders The synthesized powders were then characterized through the X-ray diffraction technique (Rigaku D/Max-2200V) for its structural analysis. The surface morphology of the powder was investigated by High Resolution Field Emission Scanning Electron Microscope (Hitachi S-4800) attached with Energy Dispersive Spectrophotometer (EDS) 30 system. XPS (Kratos Axis Ultra DLD) was utilized to analyze the chemical state of ITO nanopowder. Raman scattering spectra was measured using micro Raman spectroscopy (WITec Alpha SNOM) using a He-Ne laser as the excitation source with holographic grating of 1800 grooves/mm. 2.3.2 Preparation of Thin Films To obtain uniform coating, the glass slides were first washed with deionized water and then sonicated with acetone for 10 min at 50 °C and washed again with deionized water to remove any trace of acetone on the glass slide. The glass slides were dried in a drying oven at 80 °C prior to the coating operation. ITO nanopowder with different Sn amounts synthesized at 10 kW was then mechanically dispersed in ethanol with ammonium polyacrylic acid added as a dispersion agent, as described elsewhere [28], to obtain an ITO sol. The suspension had a blue color and showed no precipitation at a solid content of 10 wt %. The suspension was spread on a 2.5 cm × 2.5 cm borosilicate glass substrate and rotation at 2000 rpm was applied for ITO sol coating on the substrate. The prepared coating was thermally densified at 500 °C in Ar atmosphere for 1 h. 2.3.3 Analysis Methods for Films The ITO films were then analyzed using a Rigaku D/Max-2200V X-ray diffractometer with Cu Kα radiation (λ = 1.5406 Å) from 10.00° to 80.00° at a rate of 0.02°/s. The sheet Resistance of the thin film was measured by the four-probe technique. Hall effect measurements were carried out using the four-terminal method to minimize 31 the Schottky contacts. The optical properties were recorded using a UV-Vis-NIR spectrophotometer (Shimadzu UV-3600). 2.4 Results and Discussions 2.4.1 Plasma Process Parameter 2.4.1.1 Effect of plasma torch power To evaluate the effect of plasma torch powder, experiments were conducted by varying it from 10 to 30 kW under otherwise identical conditions [primary gas flow rate of 30 L/min (25 °C, 86.1 kPa), precursor feeding rate of 1.2 g/min, and carrier gas flow rate of 3.5 L/min (25 °C, 86.1 kPa)]. From the results, the product phase was seen to be affected by plasma torch power and the grain size of ITO increased with the plasma power. Figure 2.1 shows the XRD patterns of ITO nanopowder with various Sn amounts synthesized with different levels of plasma torch power. The grain sizes of synthesized powders were calculated from the position of three high intensity XRD peaks corresponding to (222), (400) and (440) lattice planes by the Scherrer equation. The grain size of ITO1 obtained at a power level of 10 kW was 18 nm, increased to 21 nm at power level of 20 kW, and further increased to 36 nm as the applied power was increased to 30 kW. The grain size also increased with an increase in plasma power for ITO2 and ITO3 samples. As the plasma torch power increases, the temperature of plasma jet also increases. At higher temperatures, the nucleation rate is affected by the degree of supersaturation. In the vapor-phase synthesis of nanoparticles, conditions are created where the vapor phase mixture is thermodynamically unstable relative to the formation of solid material 32 to be prepared in nanoparticulate form. If the degree of supersaturation is sufficient and the reaction/condensation kinetics permit, particles will nucleate homogeneously. Once nucleation occurs, remaining supersaturation can be relieved by the condensation or reaction of vapor-phase molecules on the resulting particles, and particle growth will occur rather than further nucleation [29]. At higher temperatures, the degree of supersaturation is lowered, leading to low nucleation density. Therefore, to prepare small particles, a high degree of supersaturation is required, which induces high nucleation density and then quenching the system immediately so that the particles do not grow. Coagulation rate is also enhanced at high input power promoting particle growth leading to large particle sizes [30, 31]. The plasma torch power affects the quenching rate of the powder since with the increase in the plasma torch power, the temperature gradient between plasma flame and the reactor exit becomes larger because the reactor exit is maintained below 150 °C by forced cooling system that leads to rapid quenching [32]. In ITO3, the diffraction peaks of In3Sn4O12 are observed at 10 kW alongside the main indium oxide peaks, which means extra Sn ions are not completely solvated inside the indium oxide structure. It is of interest to note that in ITO3 formed at 20 kW, SnO2 phase was observed and no In3Sn4O12 phase was formed in the product. This can be attributed to the rapid quenching of formed powder as the applied torch power was increased. Further, at the highest plasma power of 30 kW, SnO phase was observed along with the main In2O3 peaks, because at higher temperature the SnO phase is thermodynamically more stable than the SnO2 phase. 33 2.4.1.2 Effect of plasma gas flow rate at 10 kW Experiments were performed to determine the effect of plasma gas flow rate by varying the flow rate of the primary gas (Ar) from 40 L/min (25 °C, 86.1 kPa) to 80 L/min (25 °C, 86.1 kPa). Other conditions were kept the same [powder type = ITO2; plasma torch power = 10 kW; precursor feeding rate = 1.2 g/min; and the flow rate of carrier gas to carry precursor = 3.5 L/min (25 °C, 86.1 kPa)]. The XRD diffraction patterns of ITO2 obtained at different flow rates are shown in Figure 2.1 (d). The grain size of ITO2 obtained at the primary gas flow rate of 40 L/min was 23 nm and it decreased to 19 nm and 15 nm at flow rate of 60 L/min and 80 L/min. The runs were repeated to confirm the reproducibility. The growth of particles is initiated by the nucleation from the vapor phase followed by particle growth. As the plasma gas flow rate increased, the velocity of thermal plasma was increased, which led to a shorter residence time [33, 34]. Shorter residence time limited the growth of produced ITO nanoparticles. Also, at a higher plasma gas flow rate, the plasma flame length is reduced, which plays an important role on the particle formation and growth by affecting the length of time reactants and particles remain in the reaction zone [35]. 2.4.2 Raman Spectroscopy The structural properties of ITO1 and ITO2 nanopowders synthesized at 10 kW was further examined by Raman spectra, as shown in Figure 2.2. In agreement with the XRD results for ITO1 and ITO2 , the Raman spectra confirmed the existince of cubic bixbyite structure of In2O3. Cubic In2O3 belongs to the Ia3,Th7 space group and this structure has vibrational modes in the symmetry Ag, Eg and Tg, which are Raman active 34 and infared inactive, and Tu vibrational modes are infared active and Raman inactive. For cubic In2O3 the irreducible representation (IR) is given by 4Ag + 4Eg + 14Tg + 5Au + 5Eu + 16Tu [36]. In the present study, the vibrational modes were located at 308, 366, 475, 495, 556 and 631 cm-1. The main contribution to vibrational modes at 308, 366, 495 and 631 cm-1 come from the In2O3 lattice. The modes 475, 556 and 631 cm-1 are related to the Sn-O vibrations. The higher frequency line at 631 cm-1 can be attributed to the superposition of In-O and Sn-O vibration modes. This finding is consistent with the previously reported results [37, 38]. 2.4.3 X-ray Photoelectron Spectroscopy (XPS) As shown in Figure 2.3 (a), the peaks located at 446.4 eV and 452.1 eV correspond to the In 3d5/2 and In 3d3/2 states, respectively, which represent the In+3 bonding states from In2O3. Figure 2.3 (b) shows the peaks of Sn 3d5/2 and Sn 3d3/2 at 486.8 eV and 495.2 eV, respectively, which indicate the Sn+4 bonding states. The Sn 3d peaks are sharp without peak splitting, which confirms the absence of Sn+2 bonding state [39]. The intensity of Sn 3d increases with an increase in doping concentration, as is observed in Sn 3d plotted spectra of ITO1 and ITO2 synthesized at 10 kW. The XPS results show an asymmetric shape of O1s spectra, as shown in Figure 2.3 (c), which was deconvoluted into three peaks using Gaussian fitting. The main peak centered at 529.92 eV is associated with the bulk oxygen that refers to O-2 species in the cubic bixbyite structure of In2O3.The shoulder at higher binding energy contains the contributions from two components that are oxygen vacancies or defects and chemisorbed oxygen species [40, 41]. The intermediate binding energy peak (531.4 eV) is associated with oxygen 35 vacancies, and the higher binding energy peak (532.4 eV) is related to chemisorbed oxygen species or hydroxyl groups on the surface of nanopowder. 2.4.4 Scanning Electron Microscopy - Energy Dispersive Spectroscopy (SEM-EDS) Figure 2.4 shows the representative SEM micrographs of ITO1 synthesized at 10 kW and 20 kW. The particles have hexagonal and nearly spherical morphology. It appears that some of the particles are well connected to one other, indicating that the ITO particles are agglomerated to some extent. This is due to the surface forces such as Van der Waals force competing against other forces to minimize the surface free energy. Different morphologies of ITO nanopowders like nanorods [42], cubic like structures [43], nanowires [44] have been reported earlier. The particle size from the SEM micrograph of ITO1 synthesized at 10 and 20 kW is in the range of 15 - 25 nm, which is in agreement with the result obtained from the XRD analysis. Energy-dispersive spectrometer analysis was employed to investigate the composition of ITO powders synthesized at different plasma power and plasma flow rates. Figure 2.4 (c) shows the EDS spectrum of ITO1 synthesized at 10 kW. It indicated the presence of indium and tin elements where the In Lα1 and Sn Lα1 peaks appeared at 3.3 keV and 3.4 keV, respectively. The obtained EDS results for all samples indicated that the elemental distribution of In and Sn is uniform and the Sn/In (atomic %) in the product is close to the initially designed composition of ITO nanopowder. 36 2.4.5 Thin Film Characterization Results 2.4.5.1 XRD analysis X-ray diffraction was used to the study the crystal structure of prepared ITO thin films. ITO1 and ITO2 show diffraction peaks along (211), (222), (400), (440) and (602) directions with preferred orientation along the (222) direction as shown in Figure 2.5. The peak is verified to be the (222) peak of cubic bixbyite structure of In2O3 [45]. It is observed that in an ITO3 thin film, no diffraction peaks are observed except for a broad hump. This indicated the amorphous structure in the as-deposited ITO3 film. The formation of amorphous ITO3 film is due to lattice distortion caused by excess SnO2 content segregating to the noncrystalline regions in the grain boundary, which acts as scattering centers to reduce the preferred (222) orientation. 2.4.5.2 Electrical properties The electrical properties of the thin films were measured using the conventional four probe technique at room temperature. The I-V characteristics of ITO thin films shown in Figure 2.6 show ohmic behavior. The decrease in resistivity with Sn doping is due to the replacement of In+3 ions by Sn+4 ions, leading to an increase in the density of charge carriers as more number of free electrons are available for conduction. As shown in Table 2.1, the carrier density increased from 1.8 х 1020 cm-3 in ITO1 to 5.5 х 1020 cm3 in ITO2 and a consequent increase in conductivity was observed. Kim et al. [46] reported a similar trend with an initial decrease in resistivity upon increasing SnO 2 content and then an increase in resistivity beyond certain SnO2 content. The minimum resistivity obtained in the present study was 6.65 х 10-4 Ωcm in ITO2 and the resistivity 37 values obtained was close to the standard value of 10-4 Ωcm commonly achieved by sputtering. However, at a higher Sn content as in ITO3, an increase in resistivity was observed due to the increased lattice disorder by grain boundary scattering and ionized impurity scattering [47, 48], which provided strong scattering centers for charge carriers and resulted in a decrease in the mobility. This result can be related to the amorphous structure in the thin film as observed from the XRD results. In the amorphous structure of films, some electrons are bound in short range by the net nonuniform structure and when the structure of the film becomes crystalline, these trapped electrons are released from the bound [49]. At higher dopant levels, the excess Sn atoms forms defects like SnO2 or SnO phases that segregate to grain boundaries and acts as carrier traps to modify the grain boundary potential barrier [46, 50]. These phases also no longer contribute to the formation of free electrons. The decrease of hall mobility in ITO3 is due to the scattering from the grain boundaries and ionized impurity scattering. The hall mobility is expressed as [51] 𝟏 𝝁𝑯 𝟏 𝟏 =𝝁 +𝝁 𝒊 (2.1) 𝒈 where 𝜇𝑖 and 𝜇𝑔 are mobilities due to ionized impurity scattering and grain boundary scattering, respectively. 38 2.4.5.3 Optical properties Optical transmission spectra of ITO films are recorded in the wavelength region of 200 - 800 nm. Achieving an ITO thin film for transparent electrode with low resistivity and high transparency has always been the goal of the process. Figure 2.7 (a) shows the transmission spectra of ITO films. ITO1 and ITO2 films exhibited maximum transmittance of 85 %, whereas ITO3 showed a maximum transmission of 75 %. The decrease in the transmission is associated with loss of light due an increase in the scattering centers and increasing defects in grain boundaries at the higher Sn concentration. The optical absorption coefficient α of a direct band gap semiconductor near the band edge, for a photon energy hν greater than the band-gap energy Eg of the semiconductor, is given by the relation [52, 53] (αhν) = A (hν - Eg)1/N (2.2) where h is Planck's constant, and ν is the frequency of the incident photon. The constant N depends on the nature of electronic transition. In the case of ITO films, N is equal to 2 for direct allowed transition. The Tauc plot of (αhν)2 versus energy hν for all the ITO films are shown in Figure 2.7 (b). The band gap energy was obtained by extrapolating the linear plot of the Tauc plot curves to the intercept with the energy axis (at αhν= 0). It was observed that the band gap increased from 3.6 eV to 3.7 eV with increasing Sn concentration, as in ITO2 film. The blue shift exhibited on increasing the Sn concentration is associated with the Moss 39 Burstein (BM) effect. According to this well-known effect, the conduction band of the degenerate semiconductor is filled with high carrier concentration and the lowest valence energy states are blocked, leading to the lifting of the Fermi level into the conduction band and widening of the optical band gap [54]. The results are in good in agreement with the results obtained from the Hall effect measurements where the carrier concentration of ITO2 was higher than that of ITO1 film. The increase in band gap in ITO films due to an increase in the carrier concentration values have previously been reported by Kim et al. [46]. The expression for the shift in band gap Eopt is given by the following equation [55, 56] Δ Eopt= (h2/8mc*) (3N/𝝅)2/3 (2.3) where N is the carrier concentration; h is the Plank constant; and mc* is the conduction band effective mass. This suggests a direct relationship between the carrier concentration and the BM shift and an inverse relationship between the BM shift and carrier effective masses. In ITO3, there was a decrease in the transmittance observed in the visible range, which is due to the high Sn concentration that induces a negative influence on the transmittance. At a high doping concentration, that is in ITO3, the decrease in transmittance is due to excess Sn atoms segregating to grain boundaries, which can be associated with an amorphous structure of the ITO3 film. The decrease in band gap in ITO3 to 3.65 eV can be again explained by the decrease in carrier concentration, which causes a red-shift, indicating that the variation in band gap is dependent on the change in the electron concentration. 40 2.4.5.4 Gas sensitivity study After depositing ITO1 film on a glass substrate, the device was fabricated as a resistive type gas sensor. The substrate was 2.5 cm х 2.5 cm in size. An ITO1 film of thickness 300 nm deposited by the spin-coating process was used to fabricate the device. Highly conductive silver paste was used on both ends of the film for ohmic contacts to permit electrical measurements. The sensor was then placed in an alumina tube inserted coaxially inside a tubular furnace so as to adjust and optimize the temperature of the sensor. For response measurements, a known volume of gas was flowed inside the alumina tube and subsequently a decrease in the electrical resistance of the film was observed. The gas concentration was controlled by mass flow controller, and the electrical resistance before and after exposure to H2 gas was measured using a Keithley 2000 DMM. The film was tested for the operating temperature in the range of 100 °C to 500 °C in the steps of 50 °C, and the gas concentration was varied from 50 ppm to 600 ppm. The resistance of the sensor in H2 gas was measured as a function of time up to its saturation in response and then air was flowed into the tube. The resistance in air was measured until the recovery of sensor resistance to its original value in air. Different kinds of oxygen adsorbates (O2¯, O¯, and O2¯) are present, which cover the surface of the sensing material. O¯ is the most reactive species with reducing gases at operating temperatures from 300 °C to 450 °C. The response of the ITO sensor to H2 gas is explained as follows. Initially, oxygen is chemisorbed on the ITO surface. Depending upon the operating temperature, the chemisorption of oxygen forms different oxygen species, as shown below [57, 58]. 41 O2 (gas) → O2 (ads) (2.4) O2 (gas) + e⁻ → O⁻2 (ads) (2.5) O2 (gas) + e⁻ → 2O⁻ (ads) (2.6) 2O⁻ (ads) + e⁻ → O2⎯ (ads) (2.7) where O⎯2(ads), 2O⁻(ads) and O2-(ads) stand for the different physisorbed and chemisorbed surface oxygen species. Above the temperature of 450 °C, O2-(ads) ions are the major oxygen species, which are not adsorbed because these species are not stable and are usually trapped by oxygen vacancies. In the temperature range of 150 °C to 450 °C, O⁻(ads) plays the major role. The chemisorption of oxygen species leads to the increase in the sensor resistance as electrons are acquired from the conduction band and creates depletion region on the surface of ITO film to increase the potential barrier [59, 60]. When hydrogen is exposed to an ITO sensor, hydrogen molecules react with O¯(ads) to reinject the electrons into the conduction band. This increases the electron concentrations, which in turn leads to a decrease in sensor resistance. H2(ads) + O¯(ads) →H2O(gas) + e− (2.8) The hydrogen gas sensing properties of ITO gas sensor were measured in two ways: 1) by varying the sensor operating temperature and keeping the gas concentration fixed, and 2) by varying the gas concentration and keeping the sensor operating temperature fixed. 42 The resistance of the ITO sensor was measured in air as Ra before exposing to H2 gas. That of the ITO film exposed to H2 gas was measured as Rg. The sensitivity of gas sensor was then calculated from [61] 𝑺= 𝑹𝒂 −𝑹𝒈 𝑹𝒈 х 100. (2.9) The maximum response to a gas of interest depends on the operating temperature of the gas sensor. Figure 2.8 (a) shows the relationship between sensitivity and operating temperature for the ITO1 thin film at 400 ppm H2. It can be seen that the sensitivity goes on increasing with operating temperature and attains a maximum at 350 °C. It then decreases with an increase in the operating temperature. At low operating temperatures, the gas molecules do not have enough thermal energy to overcome the potential barrier developed during the adsorption of oxygen on the film surface, and thus the speed of the chemical reaction is restricted leading to low gas response. At high temperatures, the thermal energy obtained is high to overcome the potential barrier. This increases the electron concentrations significantly and leads to an increase in the response of gas sensor [62]. The temperature at which the gas sensitivity reaches a maximum is the thermal energy needed for the reaction to proceed spontaneously. It was observed from the variation of sensitivity with temperature that the sensitivity increases with an increase in the operating temperature of the gas sensor, reaches a maximum value, and then falls when the operating temperature is further increased. This is because at a higher operating temperature beyond the optimum temperature, the oxygen molecules are desorbed from the surface of the sensor [63]. The desorption of surface adsorbed oxygen tends to 43 decrease the active sites for the chemical reaction with H2 gas molecules, which in turn reduces the film resistance and sensitivity of the gas sensor. Sahay et al. [58] investigated the gas sensing properties of AZO films to methanol vapors. They observed that there existed an optimum operating temperature of a sensor to achieve the maximum response to a gas of interest, dependent on the kind of gases, that is, the mechanism of dissociation and further chemisorption of a gas on the particular sensor. It was observed from Figure 2.8 (b) that the sensitivity increases rapidly with an increase in H2 concentration at 350 °C, and finally saturation of the sensor response was observed beyond the tested H2 concentration of 400 ppm. This is attributed to the fact at lower H2 gas concentrations, the surface reaction proceeds slowly due to the smaller coverage of H2 molecules. With a further increase in gas concentration, the surface reaction increases due to the larger coverage of H2 molecules and thereby resulting in an increase in the gas sensitivity. However, at high concentrations beyond the optimum range, a multilayer of gas molecules is formed, resulting in the saturation in response [58]. The excess gas molecules remain idle and do not reach the surface-active sites necessary for surface reaction to take place. Thereby, no increase in the sensitivity of gas sensor was observed beyond a certain concentration. Shukla et al. [17] investigated the hydrogen gas sensitivity of an ITO film (6.5 mol % SnO2) with and without Pt- sputtering and observed a parabolic relationship between the sensitivity and gas concentration contrary to linear relationship obtained in our study. Along with the operating temperature and gas concentration, the effectiveness of the gas sensor for practical application also depends on the response and the recovery time as well. These are defined as the time required for resistance to undergo 90 % 44 variation of its equilibrium value, respectively, in the presence and absence of a test gas [64]. Figure 2.8 (c) shows the transient response and recovery characteristic of ITO1 to 400 ppm H2 gas at the operating temperature of 350 °C, and it was found that the injection of 400 ppm H2 induced a significant decrease in electrical resistance of the layer. It is observed from the figure that the ITO1 film showed fast response as well as recovery time of 28 s and 40 s, respectively, for 400 ppm at 350 °C. 2.5. Conclusions The thermal plasma process performed in this work has shown its potential as an efficient technique for synthesizing ITO nanopowder. This process is also suitable for large scale production of nano-sized powders due to the availability of high temperatures for volatizing reactants rapidly, followed by vapor phase reactions and rapid quenching to yield nano-sized powder. The experimental results showed that the grain size of ITO nanopowder increased with an increase in plasma torch power, and also product phases were affected by the plasma torch power. Grain size decreased with an increase in plasma gas flow rate within the range tested. ITO films were prepared on glass substrates by spin coating using ITO sol solution. Optical transmittance in the visible region approached 85 % in ITO1 and ITO2 films. The deposited films showed enhanced electric properties with resistivity in the order of 10-3 -10-4 Ωcm, and in particular ITO2 showed the best performance: a high carrier concentration of 5.5 х 1020 cm-3 and a low electrical resistivity of 6.65 х 10-4 Ωcm. 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Nath, "Al-doped ZnO thin films as methanol sensors," Sens. Actuators B, vol. 134, no. 2, pp. 654-659, 2008. 59. N. Srinatha, Y. S. No, V. B. Kamble, S. Chakravarty, N. Suriyamurthy, B. Angadi, A. M. Umarji, and W. K. Choi, "Effect of RF power on the structural, optical and gas sensing properties of RF-sputtered Al doped ZnO thin films," RSC Adv., vol. 6, no. 12, pp. 9779-9788, 2016. 60. N. G. Patel, P. D. Patel, and V. S. Vaishnav, "Indium tin oxide (ITO) thin film gas sensor for detection of methanol at room temperature," Sens. Actuators B, vol. 96, no. 1-2, pp. 180-189, 2003. 61. M. A. Basyooni, M. Shaban, and A. M. E. Sayed, "Enhanced gas sensing properties of spin-coated Na-doped ZnO nanostructured films," Sci. Rep., vol. 7, p. 41716, Jan. 2017. 62. P. P. Sahay and R. K. Nath, "Al-doped zinc oxide thin films for liquid petroleum gas (LPG) sensors," Sens. Actuators B, vol. 133, no. 1, pp. 222-227, 2008. 50 63. H. Windischmann, "A model for the operation of a thin-film SnOx conductancemodulation carbon monoxide sensor," J. Electrochem. Soc., vol. 126, no. 4, p. 627, 1979. 64. S. Isik, O. Coban, C. Shafai, S. Tuzemen, and E. Gur, "Growth conditions effects on the H2 and CO2 gas sensing properties of indium tin oxide," J. Phys.:Conf. Ser., vol. 707, p. 012021, 2016. Table 2.1 The electrical properties of ITO films Magnetic Field (T) Type of Film Carrier Density (cm-3) Mobility (cm2/V-s) Resistivity (Ωcm) Conductivity (S/cm) 0.35 0.35 0.35 ITO1 ITO2 ITO3 1.8 х 1020 5.5 х 1020 2.2 х 1020 14.0 17.1 3.64 2.47 х 10-3 6.65 х 10-4 7.79 х 10-3 404.8 1503.8 128.4 51 Figure 2.1 X-ray patterns of ITO1, ITO2, and ITO3 obtained at different plasma torch power and gas flow rates. 2.1 (a) X-ray diffraction patterns of ITO1, ITO2, and ITO3 obtained at 10 kW plasma torch power. The crystal planes are indexed corresponding to the cubic bixbyite structure of In2O3. Diffraction peaks of In3Sn4O12 were observed around the main indium oxide peaks in ITO3 sample. 2.1 (b) X-ray diffraction patterns of ITO1, ITO2, and ITO3 obtained at 20 kW plasma torch power. The crystal planes are indexed corresponding to the cubic bixbyite structure of In2O3. SnO2 peaks were observed at higher plasma power of 20 kW in ITO3 sample. 2.1 (c) X-ray diffraction patterns of ITO1, ITO2, and ITO3 obtained at 30 kW plasma torch power. The crystal planes are indexed corresponding to the cubic bixbyite structure of In2O3. SnO peaks were observed at higher plasma power of 30 kW in ITO3 sample. 2.1 (d) X-ray diffraction patterns of ITO2 obtained at different plasma flow rates: 40 L/min, 60 L/min, and 80 L/min at applied plasma torch power of 10 kW. The crystal planes are indexed corresponding to the cubic bixbyite structure of In2O3. SnO2 peaks were observed at higher plasma flow rate of 80 L/min. 52 53 720 308 366 475 495 Intensity (a.u.) 660 600 (b) 556 631 540 830 308 366 747 495 475 664 (a) 631 581 300 400 500 600 -1 Wavenumber (cm ) Figure 2.2 Room temperature Raman spectra of a) ITO1 and b) ITO2. The main contribution to vibrational modes at 308, 366, 495, and 631 cm-1 comes from the In2O3 lattice and the modes 475, 556, and 631 cm-1 are related to the Sn-O vibrations. 54 Figure 2.3 XPS spectra of samples. 2.3 (a) XPS In 3d narrow spectra of samples: ITO1 and ITO2. The peaks located at 446.4 eV and 452.1 eV correspond to the In 3d5/2 and In 3d3/2 states, respectively, which represent the In+3 bonding states from In2O3. 2.3 (b) XPS In3d narrow spectra of samples: ITO1 and ITO2. The peaks located at 446.4 eV and 452.1 eV correspond to the In 3d5/2 and In 3d3/2 states, respectively, which represent the In+3 bonding states from In2O3. XPS O 1s narrow spectra of samples: ITO1 and ITO2. 2.3 (c) The oxygen peak was deconvoluted into three components. The component with the lowest binding energy, centered at 529.92 eV, was attributed to O-2 ions on wurtzite structure of hexagonal Zn+2 of the metal oxide. The component with medium binding energy, centered at 53.4 eV, was associated with O-2 ions in the oxygen deficient regions within the matrix of ZnO, and the component at highest binding energy centered at 534.4 eV is attributed to the chemisorbed oxygen ions. 55 56 Figure 2.4 SEM micrographs and EDS spectra. 2.4 (a) SEM micrographs of ITO1 nanopowder synthesized at 10 kW. 2.4 (b) SEM micrographs of ITO1 nanopowder at 20 kW. Particles with hexagonal and nearly spherical morphology were observed for all compositions of ITO samples. 2.4 (c) Energy-dispersive X-ray Spectrum of ITO1 nanopowder synthesized at 10 kW. 57 Figure 2.5 X-ray diffraction pattern of films: a) ITO1, b) ITO2, and c) ITO3. Amorphous structure is observed at higher doping amount in ITO3 film. 58 Figure 2.6 I-V characteristics of films: a) ITO1, b) ITO2, and c) ITO3.The lowest resistivity of 6.65 х 10-4 Ωcm is obtained with 10 % Sn in ITO2 film, and the highest resistivity of 7.79 х 10-3 Ωcm is obtained with higher doping amount of 15 % Sn in ITO3 film. Linear fit overlaps with the I-V curves. 59 Figure 2.7 Transmission curve and Band gap measurements. 2.7 (a) Transmission curves of ITO films. ITO1 and ITO2 films exhibited a maximum transmittance of 85 % whereas optical transmittance decreased to 75 % in ITO3 film. 2.7 (b) Tauc plots to determine the band gap. Band gap increased from 3.6 eV to 3.7 eV in ITO2 film followed by a decrease in band gap to 3.65 eV in ITO3 film. 60 Figure 2.8 H2 gas sensitivity analysis. 2.8 (a) The variation of sensitivity with temperature. The sensitivity goes on increasing with operating temperature and attains a maximum at 350 °C and then decreases with further increase in temperature. 2.8 (b) The variation of sensitivity with gas concentration. The sensitivity increases rapidly with an increase in H2 concentration at 350 °C, and finally saturation of the sensor response was observed beyond the tested H2 concentration of 400 ppm. 2.8 (c) Transient response and recovery characteristic of ITO1 to 400 ppm H2 gas at the operating temperature of 350 °C. CHAPTER 3 PHOTOCATALYTIC PROPERTY OF PLASMA-SYNTHESIZED ZINC OXIDE AND TIN-DOPED ZINC OXIDE (TZO) NANOPOWDERS AND THEIR APPLICATIONS AS TRANSPARENT CONDUCTING FILMS 3.1 Introduction Transparent conducting oxides (TCO) are of great scientific and commercial importance because they combine high conductivity with optical transparency in the visible region. TCOs are commonly used in organic light emitting diodes, transparent transistors, electro-optical devices and gas sensitive devices. Zinc oxide (ZnO) is a II-VI group compound semiconductor, which crystallize in a wurtzite structure belonging to the space group P63mc. ZnO has characteristics of high transparency, good UV trapping properties, nontoxicity, natural abundance etc., which are important properties of optoelectronic and piezo electronic materials owing to its large band gap of 3.37 eV and large exciton binding energy of 60 meV [1, 2]. However, ZnO films have poor conductivity and doping with various dopants are usually necessary to improve the conductivity for as use as TCO film. Doping of ZnO by replacing Zn+2 ions with higher valence ions such as Al+3 and Sn+4 ions can in general induce dramatic changes in its electrical and optical properties [3, 4]. At present, tin-doped indium oxide (indium tin oxide, or ITO) is the most commonly used TCOs, but because of concerns of 62 the supply of world indium reserves and the cost of indium, there has been an increasing interest in alternatives [5]. Tin-doped zinc oxide (TZO) is one of the most important alternatives to ITO and is widely used as transparent electrode in various kind of devices. When Sn was added into ZnO for doping, Sn+4 substitutes Zn+2 sites in the ZnO crystal structure resulting in two more free electrons to contribute to the electric conduction. The ionic radius of Sn+4 (0.069 nm) is smaller than Zn+2 (0.074 nm); therefore, Sn+4 ions can replace Zn+2 ions in substitution sites. The electrical conductivity, transparency, thermal stability, and durability make this material interesting and attractive [6, 7]. Also, oxides of TZO can be used as photocatalyst because of their high activity and chemical stability. Dyes are an important source of environmental pollution. It is used in many industrial processes, namely cosmetic, textile, and printing. About 15 % of the world production of dyes is lost during the dying process and is released in textile effluents [8]. Methylene Blue is a cationic dye that is usually used in the textile industry for dying linen, wool and silk. The discharge of large amounts of effluent containing different dyes is harmful to microbes, the aquatic system, and human health. The consequence of colored water is detrimental to environment because color obstructs the sunlight access to aquatic organism and plants, and it diminishes photosynthesis and affects the ecosystem [9, 10]. In recent years, an advanced oxidation process utilizing photocatalyst has attracted a great deal of attention in wastewater treatment because of its high activity, mild reaction conditions, and low energy consumption compared to conventional methods [11, 12]. 63 The mechanism of photocatalytic degradation is explained as follows [13, 14]. When the photocatalyst is irradiated with photons of energy equal to or greater than the band gap energy of the photocatalyst, the electrons get excited from the valence band to the conduction band with the reaction of holes (h+) in the valence band. Photocatalyst + hv → eCB- + hVB+ (3.1) The electrons are trapped by the dissolved O2 or by the adsorbed O2 to give rise to superoxide radicals: eCB- + O2 → O2• - (3.2) The superoxide radicals can react with H2O2 to form hydroperoxyl radicals (HO2•) and hydroxyl radicals(OH•), which are strong oxidizing agents to decompose the organic molecules. O2• - + H2O→HO2•+ OH • (3.3) The photoinduced holes can be trapped by hydroxyl groups (or H2O) on the photocatalyst surface to yield hydroxyl radicals: hVB+ + OH- → •OHads (3.4) 64 hVB+ + H2O→•OH + H+ (3.5) Thereby organic molecules will get oxidized to yield CO2 and H2O as follows: • OH + organic molecules + O2 →products (CO2 + H2O) (3.6) Meanwhile, the recombination of positive hole and electron will take place, which reduces the photocatalytic activity of prepared TZO photocatalyst. eCB- + hVB+→ Photocatalyst (3.7) An innovative plasma processing technique has been developed for the preparation of nano-sized TZO powder by vapor phase reactions. Thermal plasma provides a high processing rate as well as other advantages like good control over size, shape and crystal structure as well as a clean reaction atmosphere that yields high purity products, a high quench rate to form ultra-fine powder, and a wide choice of reactants [15]. Compared with other methods it avoids multiple steps like in mechanical milling, sol -gel method, and precipitation method and does not require a high liquid volume and surfactants, which are involved in a wet chemical process. Thin films of TCOs on glass substrates have been prepared by a variety of physical and chemical deposition techniques for transparent and conductive electrodes, albeit with inherently costly and time-consuming process. Over other techniques, fabricating TZO films from well-dispersed TZO nanoparticles is a good alternative 65 because it has the advantages of low processing cost and easy control of composition ratio. In addition to cheaper processing technique, nanoparticles could be handled efficiently, and loss of raw materials could be minimized [16, 17]. Different synthesis techniques of TZO nanoparticles have been reported. Wu et al. [18] fabricated Sn-doped ZnO nanorods by hydrothermal treatment and used methyl orange as the probe molecule to evaluate its photocatalytic activity. Wang et al. [19] prepared tin-doped zinc oxide nanoparticles in organic solution, with metal acetylacetonate as the precursor and oleyl amine as the solvent. Javid et al. [20] synthesized Sn-doped ZnO nanoparticles by chemical solution method using zinc nitrate and NaOH as precursors. Junlabhut et al. [21] synthesized Sn-doped ZnO nanopowders by coprecipitation method with various Sn additives from 0- 50 wt %. Verma et al. [22] reported the structure property correlation in pure and Sn-doped ZnO nanocrystalline materials prepared by coprecipitation. Li et al. [23] used the p-type Si (100) substrate for the synthesis of TZO nanowires using the vapor-liquid-solid growth process. In this study, for the first-time zinc oxide and tin-doped zinc oxide nanopowders were synthesized by the plasma process and tested for its photocatalytic property in the degradation of methylene blue. The obtained nanopowders were also used to fabricate ZnO and TZO films to study the effect of Sn doping on their structural, optical, electrical, and photocurrent properties. The high figure of merit and obtained resistivity values were much lower than that of reported results on TZO films [24 - 26], indicating their suitability for use as transparent electrical contacts in optoelectronic devices. 66 3.2 Experimental Apparatus The plasma reactor system used for the synthesis of zinc oxide and tin-doped zinc oxide nanopowder consisted of a downward plasma torch, a power supply unit, a cylindrical reactor, a powder feeding system, a cooling system, powder delivery system, powder collectors, a gas delivery system, an off-gas scrubber and an off-gas exhaust system. The plasma torch had a copper cathode and a tungsten anode. It was water-cooled and operated at atmospheric pressure. The reactor assembly consisted of a vertical stainless-steel tube, which was water-cooled, and had the dimensions of 15 cm inner diameter and 60 cm length. A graphite cylinder of 7.6 cm inner diameter and 60 cm length was placed inside the stainless-steel tube. A graphite felt filled the gap between the graphite cylinder and the stainless-steel tube. The cooling chamber consisted of two-layer stainless steel to cool the outgoing gases to lower than 150 °C. The precursor was directly fed into the plasma gun using Ar as the carrier gas through a powder feeding system consisting of a test tube filled with the precursor powder, a motor that pushed up the test tube at a constant rate, a carrier gas line that carried the fluidized particles from the top of the particle bed in the test tube, and a vibrator. The product was collected on a Teflon coated polyester filter with a pore size of 1 µm. The off-gas scrubber used a 5 % NaOH solution. More details about the experimental set-up can be found in a previous publication [27 - 30]. 3.3 Experimental Procedure The precursors used for synthesizing zinc oxide and tin-doped zinc oxide in this work were as follows: 1) zinc nitrate [Zn(NO₃)₂•6H₂O , Alfa Aesar, Haverhill, MA] powder for zinc oxide, and 2) a mixture of zinc nitrate [Zn(NO₃)₂•6H₂O, Alfa Aesar, 67 Haverhill, MA) and tin nitrate [Sn(NO₃)4, American Elements, Los Ange, CA] powders for tin-doped zinc oxide. Each precursor was ground and crushed by mortar and pestle and sieved until the final particle size was around 50 µm. The milled precursors were kept in a vacuum oven at 50 °C to remove the moisture for ease of powder feeding. The anhydrous milled precursors were uniformly mixed using a vibrating mixer. The uniformly mixed precursor was fed into the plasma flame through the internal port in the plasma torch of 2 mm in diameter, using a specially designed powder feeding system mentioned above. The experimental conditions were as follows: 1) plasma gas flow rate (Ar): 40 L/min (25 °C and 86.1 kPa total pressure at Salt Lake City); 2) precursor feeding rate: 0.5 g/min; 3) flow rate of carrier gas (Ar): 3.5 L/min and 4) plasma power of 15 kW. The amount of tin nitrate was varied to obtain 3 and 5 atomic percent Sn designated as TZO1 and TZO2, respectively. 3.3.1 Analysis Methods for Nanopowders The synthesized powders were characterized through an X-ray diffraction technique (Rigaku D/Max-2200V) for structural analysis. The surface morphology of the powder was investigated by a High-Resolution Field Emission Scanning Electron Microscope (Hitachi S-4800) equipped with an Energy Dispersive Spectrophotometer (EDS) system. XPS (Kratos Axis Ultra DLD) was utilized to analyze the chemical state of ZnO and TZO nanopowders. Raman scattering spectra were measured with a micro Raman spectrometer (WITec Alpha SNOM) using a He-Ne laser as the excitation source with holographic grating of 1800 grooves/mm. The absorption spectra of ZnO and TZO 68 with various dopant concentrations were investigated by diffuse reflectance spectroscopy (DRS) in which the scanning was done in the wavelength range of 300-800 nm. Photoluminescence (PL) spectra were examined by Perkin Elmer spectrophotometer with a Xe lamp using an excitation wavelength of 350 nm. 3.3.2 Photocatalysis Set-Up Photocatalytic degradation of methylene blue by zinc oxide and tin-doped zinc oxide with various doping concentrations were carried out in 1-liter volume reaction vessel, as shown by a schematic sketch in Figure 3.1. The reaction vessel is fabricated of borosilicate glass to accommodate the immersion well. The double-walled, quartz immersion well consists of a small diameter inner tube, which extends down the annular space to insure flow of coolant from the bottom of the well upward to an outlet. Cooling water was circulated to control the solution temperature. A 450 W medium pressure quartz mercury vapor lamp was inserted vertically in the immersion well. Approximately 60 % of the radiated energy is in the ultraviolent portion of the spectrum, 35 % in the visible region, and the balance in the infrared range. A 6-foot power cord allowed for lowering the lamp into the well. The bottom of the reactor was flat to allow the use of a magnetic stirrer. The working volume in the reactive area of the lamp was approximately 40-50% of the total volume. The reaction vessel had a 14/20 size taper joint for withdrawing the solution after certain time intervals using sparger tubes and one side arm provision for thermometer insertion. The transformer operating at 115 V and 60 Hz was used to supply the extra voltage and current required to initiate the arc. 69 3.3.3 Measurement of Photocatalytic Activity The photocatalytic activity of the prepared slurry was evaluated by the degradation of a methylene blue solution. Prior to irradiation, the slurry was magnetically stirred for 30 min to obtain the equilibrium of adsorption and desorption. The slurry was prepared with 0.1 g of the photocatalyst dispersed in 500 ml of 85 µM methylene blue (MB) solution. At time intervals of 20 min, 10 ml of the solution were collected from the slurry suspension and centrifuged at 5000 rpm for 10 min. The centrifuged MB solution was filtered by a Millipore filter to remove particles, and the filtrate was analyzed using a UV-VIS spectrophotometer (Shimadzu UV-3600) by measuring the light absorption intensity at 664 nm where the maximum absorption intensity was attained. A calibration plot based on Beer-Lambert's law was established by relating absorbance to the concentration. The reaction kinetics of the photocatalysis in general follows the following Langmuir-Hinshelwood (L-H) equation: 𝑹= 𝒅𝑪 𝒅𝒕 =− 𝒌𝒔 𝑲𝑪 (𝟏+𝑲𝑪) (3.8) where ks is the surface reaction rate constant, K is the adsorption coefficient of the reactant, and C is the concentration of methylene blue at any time t. At low concentrations of methylene blue and weak adsorption (KC << 1), Eq. 3.8 is simplified to R= where k = Kks. 𝒅𝑪 𝒅𝒕 = -kt (3.9) 70 The integration of Eq. (3.9) gives: Ln (C) = -kt + Ln (C0) (3.10) where C0 is the initial concentration of methylene blue, k is the pseudo-first-order reaction rate constant (min-1), and t is the reaction time (min). The reaction rate constant (k) was calculated from the slope of ln(C) vs. time plot. The percentage degradation was determined using the following formula: % degradation = 𝑪𝒊 −𝑪𝒇 𝑪𝒊 х 100 (3.11) where Ci and Cf are the initial and final concentrations of the dye. 3.3.4 Preparation of Thin Films ZnO and TZO nanopowders with different dopant concentrations were mechanically dispersed in polyethylene glycol (PEG600) and carbonic acid was added as a dispersion agent. The wetted powder was then ground in mortar to obtain a homogenous paste. The paste was then dissolved in 1-propanol. ZnO and TZO coatings were prepared by the spin coating process on a 2.5 cm × 2.5 cm borosilicate glass substrate using the dispersion at 2000 rev/min for 60 s. The prepared coating was thermally densified at a temperature of 500 °C in an inert atmosphere. 71 3.3.5 Analysis Method for Thin Films Deposited ZnO and TZO films were then analyzed using a Rigaku D/Max-2200V X-ray diffractometer with Cu Kα radiation (λ = 1.5406 Å) from 10.00° to 80.00°at a rate of 0.02°/s. The sheet resistance of the thin film was measured by the four-probe technique. Hall effect measurements were carried out using the four-terminal method to minimize Schottky contacts. The optical properties were recorded using a UV-Vis-NIR spectrophotometer (Shimadzu UV-3600). 3.4 Results and Discussions 3.4.1 X-Ray Diffraction (XRD) Figure 3.2 shows the XRD patterns of ZnO, TZO1 and TZO2. The peaks were indexed using the X'Pert High Score Plus, and the peaks correspond to the hexagonal wurtzite structure of ZnO. Narrow and sharp peaks confirm the good crystallinity of synthesized nanopowder. No SnO2 or other Sn phases were detected in TZO1 sample but in TZO2 sample, SnO2 phase was observed indicating that the Sn content has exceed its maximum solubility in ZnO. The presence of SnO2 phase as observed in TZO2 lowers the crystallinity of TZO nanopowder, which is evident from the decrease in the intensity of peaks. Small shifts in the diffraction peaks are observed in TZO1 and TZO2 to higher diffraction angles with the doping of tin into the zinc oxide lattice. Crystallite size is decreased with Sn concentration at 3 atm % (TZO1), which is due to lattice shrinkage from the difference in the ionic radii of Zn+2 and Sn+4 ions. However, an increase in the crystallite size is observed in TZO2 (5 atm %), which is due to excess Sn+4 ions segregating to the grain boundaries. The reflections (100), (002), and (101) are of high 72 intensity and other reflections (102), (110), (103), and (112) are of lower intensity. Since the (101) reflection is the highest in all the samples, we can conclude the particles are orientated mostly in the [101] direction. In TZO2, excess Sn+4 serves as an inhibitor for the growth of particles in the [101] direction. The crystallite size was determined by the following Debye-Scherrer equation [31]: D = Kλ/ (β cos θ) (3.12) where β =√(β2(FWHM) - β02) the peak broadening after removing the instrumental broadening, β (FWHM) is the full width at half maximum, and β0 is the correction factor (0.005 rad). The crystallite size is determined from the three peaks (100), (002), and (101) and the average size is calculated. The lattice constants are calculated by the following equation [32, 33] 𝟏 = 𝟒 𝒉𝟐 +𝒉𝒌+𝒌𝟐 𝒅(𝒉𝒌𝒍)𝟐 𝟑 ( 𝒂𝟐 ) + 𝒍𝟐 𝒄𝟐 (3.13) where a and c are the lattice constants, and d (hkl) is the crystalline plane distance for indices (hkl) calculated from the XRD pattern. The position of the most intense peaks are slightly shifted towards a higher angle as compared to the ZnO XRD pattern. This indicates the replacement of Zn+2 by Sn+4 ions in the ZnO lattice. The crystallite size is 73 decreased on increasing the doping amount to 3 atm % as in TZO1, and a simultaneous decrease in FWHM also observed, as shown in Table 3.1. 3.4.2 Raman Spectroscopy ZnO has a wurtzite structure and the symmetry of the ZnO crystal structure lies in the spatial group of C46v with two formula units in the primate cell. Therefore, its optical properties at the center of Brillouin zone (I-) is represented by the following irreducible relation: I- = 1A1 + 2B1 + 1E1 + 2E2 (3.14) The polar characteristics of A1 and E1 modes are split into LO (longitudinal) and TO (Transverse) modes, with different frequencies. Modes A1, E1, and E2 are Raman and infrared active whereas the two B1 modes are infrared and Raman inactive and are silent modes [34, 35]. A single broad LO peak is observed at 588 cm-1, which is due to the overlapping of A1(LO) and E1(LO) modes. In this study, as shown in Figure 3.3, a strong and sharp peak at 437 cm-1 is attributed to the high frequency of E2(high) mode and is the characteristic of hexagonal wurtzite structure of ZnO [36]. ZnO & TZO1 exhibited high intense Raman peaks suggesting improved crystal quality, whereas for TZO2, the intensity of E2(high) mode decreased, indicating that the crystallinity of TZO2 nanopowder gets deteriorated at a higher doping amount, and this observation corroborates with the XRD data. Junlabhut et al. [21] also reported that the Raman signals obtained in TZO nanopowders are highly sensitive to crystalline nature. A small peak 74 shift was observed in the E2(high) mode for TZO2 sample with the usual position of ZnO at 437 cm-1, indicating the decreased order of wurtzite phase when Sn+4 ions are introduced into the ZnO lattice. The peak at 670 cm-1 corresponding to A1 mode appears only in TZO1 and TZO2 [21, 37]. The intensity of this peak is significantly stronger in TZO2 than in other samples. 3.4.3 X-ray Photoelectron Spectroscopy (XPS) and Photoluminescence Spectroscopy (PL) XPS analysis was carried out to verify the concentration of Sn and to ascertain the valence of Sn in TZO1 and TZO2. The binding energy difference of 23.10 eV between Zn 2p3/2 and Zn 2p1/2 and the binding energy position indicate that Zn in ZnO, TZO1 and TZO2 nanopowder exist in +2 oxidized state. To ascertain the valence state of Sn, Sn 3d spectra was acquired from the samples under study. The Sn 3d5/2 and Sn 3d3/2 peaks are located at 486.6 eV and 495.2 eV, respectively. The Sn 3d3/2 signal was unusually intense due to the Auger ZnL3M45M45 transition, which made it difficult to directly compare the spectral parameters of Sn 3d3/2 with those of the reference standard. However, the binding energy of 486.6 eV indicated that Sn is incorporated in the form of Sn+4 bonding state from SnO2 [38 - 40]. This further gives support to the Sn+4 ions substituting for Zn+2 sites. The observed Sn 3d5/2 peak also ruled out the possibility of segregation of metallic Sn or the presence of other phases like SnO. The Zn 2p spectra are located at 1021.42 eV and 1044.34 eV, as shown in Figure 3.4 (a), which corresponds to Zn 2p3/2 and Zn 2p1/2, respectively. The intensity of Sn 3d5/2 increased with an increase in doping concentration, as was observed in Sn 3d plotted spectra, shown in Figure 3.4 (b). The XPS results showed an asymmetric shape of O 1s spectra as shown in Figure 3.5, which was 75 deconvoluted into three peaks using Gaussian fitting. The main peak centered at 529.92 eV is associated with the bulk oxygen that refers to O-2 species in the hexagonal structure of ZnO. The shoulder at higher binding energy contains the contributions from two components that are oxygen vacancies or defects and chemisorbed oxygen species. The intermediate binding energy peak (531.4 eV) is associated with the oxygen vacancies, and the higher binding energy peak (532.4 eV) is connected with chemisorbed oxygen species or hydroxyl groups on the surface of the nanopowder [41 - 44]. It was observed that with Sn addition, a relative increase in intensity of component Ov to OL-2component was observed, indicating that the concentration of oxygen vacancies varies with doping. The changes in the intensity of Ov component is in connection with the variation in the concentration of oxygen vacancies [45, 46]. The PL spectra of undoped and Sn-doped ZnO nanopowders are shown in Figure 3.6. All of the spectra showed a strong UV emission peak and a green emission peak (deep-level emission) in the visible region. The UV emission also known as near-band edge emission arises from the band edge recombination of free excitations and green band emission peak is related to the singly ionized oxygen vacancy in ZnO, and results specifically from the recombination of a photo generated hole with singly ionized charge state of this defect [47, 48]. It is evident that the ratio of DLE/UV is larger for TZO1 followed by TZO2 and ZnO nanopowders. Thus, the weak green band edge emission in the PL spectrum for undoped ZnO indicates low concentration of oxygen vacancies present. This further confirms that in TZO, Sn atoms exist substitutionally, sharing the oxygen with Zn atoms and hence increasing the defect of oxygen vacancy which enhances the intensity of green band emission. 76 3.4.4 Scanning Electron Microscopy- Energy Dispersive Spectroscopy (SEM-EDS) Figure 3.7 (a) shows the typical SEM micrographs of TZO1 particles synthesized at 10 kW. The formation of nearly spherical shaped morphology was observed for all compositions of TZO particles. Different morphologies of TZO nanopowders have been reported earlier like nanorods [18], cubic like structures [21], and nanowires [23]. Figure 3.7 (b) shows the EDS spectrum for TZO1 synthesized at 10 kW. It indicated the presence of indium and tin elements where the Zn Lα1 and Sn Lα1 peaks appeared at 1.1 keV and 2.7 keV, respectively. The obtained EDS results for all samples indicated that the elemental distribution of Zn and Sn is uniform and the Sn/Zn (atomic %) in the product is close to the initially designed composition of TZO nanopowder. 3.4.5 Photocatalysis Data and Absorption Spectrum Analysis The photocatalytic activity of ZnO, TZO1, and TZO2 nanopowders were investigated using MB as an indicator under UV light. The results demonstrated that TZO1 is a superior photocatalyst to ZnO. From Figure 3.8, it is clear that the degradation of MB is enhanced significantly when catalyzed by Sn-doped ZnO nanopowder. The percentage degradation increased from 42.0 % to 94.6 %. The degradation rate constant for TZO1 increased to 0.0339 min-1 from 0.0106 min-1 as shown in Figure 3.9 (a). The enhanced photocatalytic efficiency of TZO1 is attributed to the increase in the specific surface area and band gap of the nanopowder. Upon doping Sn in ZnO, the BET specific surface area increased from 18.5 m2/g to 48.6 m2/g. The enhanced S/V ratio facilitated the increase in the number of the active states on the photocatalyst surface and thereby 77 increased the concentration of photo generated carriers. Thus, the relative number of free radicals attacking the dye molecules increases. Another factor influencing the rate of degradation is the band gap. Figure 3.9 (b) shows TZO1 nanoparticles have maximum absorbance at 395 nm corresponding to a band gap of 3.16 eV, which was redshifted compared to undoped ZnO nanoparticles that showed a maximum absorbance at 383 nm corresponding to a band gap of 3.22 eV. It clearly indicates that TZO nanoparticles have higher absorption than ZnO nanoparticles, that is, TZO nanoparticles have higher photocatalytic activity than ZnO nanoparticles. The red shift observed with lower band energy was also observed for other doped ZnO samples. Ravishankar et al. [49] reported that the band gap of Ag-doped ZnO nanoparticles was red-shifted compared to that of undoped ZnO sample. Rajbongshi et al. [50] reported a red shift in N-doped ZnO compared to that of pristine ZnO. When doped ZnO gets excited by photons with energy greater than band gap energy, a larger number of electrons are produced from VB to CB of ZnO and SnO2, leading to generation of electron-hole pairs. The electrons transfer from the CB of ZnO to CB of SnO2 and conversely, the holes transfer from the CB of SnO2 to CB of ZnO decreases the recombination rate of electron-hole pairs [51]. Thus, by separating the arrival time of electrons and holes at the surface of the photocatalyst, the probability of recombination of pairs is restrained, facilitating the enhancement of photocatalytic activity especially in TZO1. Higher photocatalytic activity in TZO1 is also attributed to higher oxygen vacancy, revealed from the PL spectra. All of the PL spectra showed UV emission peak and a green emission peak (deep-level emission) in the visible region. Oxygen vacancy is one of the important factors for narrowing the band gap, as observed above in conjunction with Figure 3.6. Ansari et al. [52] suggested that the band 78 tail was formed below the CB of ZnO due to O-vacancy level energies, and Wang et al. [53] reported that an impurity level was created by O-vacancy level energies nearer to the valence band than the undoped ZnO nanopowder. Oxygen vacancies facilitate the reduction in the recombination rate of electrons and holes. These vacancies act as electron acceptors during the photocatalytic reaction, thereby ensuring that holes in the valence band react with the hydroxyl ions to produce hydroxyl radicals, which are the main oxidant species [52 - 54]. Wu et al. [18] reported that the photocatalytic activity of Sndoped ZnO catalyst increased with an increase in Sn content. In this range of Sn contents, no peaks corresponding to SnO2 crystal or other Sn phases were detected. They concluded that an increase in the catalytic activity with Sn content within their range was because of an increase in the singly ionized oxygen vacancy. To study the optimum Sn doping amount, the photocatalytic activity of TZO2 nanopowder was investigated in the present work. It was apparent that the photocatalytic activity of TZO2 was lower than that of TZO1. This is due to the fact that at higher doping concentrations, the excess Sn cannot enter the lattice of ZnO crystal or further alter the band gap and segregates at grain boundaries to form defect clusters. These defect clusters serve as recombination centers and against the separation of excited electron and hole [55]. The ratio of the intensities of DLE to UV emission peaks obtained was lower for TZO2 (5 at. % Sn) sample than for TZO1 (3 at. % Sn), indicating a decrease in the concentration of oxygen vacancies and corresponding decrease in the photocatalytic activity of TZO2 sample. 79 3.4.6 Thin Films Analysis 3.4.6.1 X-ray diffraction (XRD) The crystalline quality and orientation of the ZnO and TZO thin films prepared in this work were investigated by means of XRD. All the films exhibited a sharp peak near 34.54° corresponding to the (002) plane of the hexagonal wurtzite structure of ZnO and exhibited a c-axis orientation perpendicular to the substrate surface, which is evident from the sharp and strong (002) diffraction peak as shown in Figure 3.10. It is clear that with the addition of Sn content, the crystalline quality of the TZO1 thin film remained the same as the ZnO film. In TZO1, the (002) plane becomes sharper and narrower as compared to the TZO2 film. The decrease in intensity at the (002) plane for the TZO2 film and increase in the intensity of (101) and (004) indicated that excess Sn deteriorates the crystalline quality of thin film due to lattice distortion. Bedia et al. [25] also reported the variations in the intensity of (002) plane with Sn content and indicated that the quality of film was improved with 1 % Sn concentration. The lattice distortion is due to excess Sn segregating to the noncrystalline regions in the grain boundary, which acts as scattering centers to reduce the preferred c-axis orientation [56, 57]. 3.4.6.2 Electrical properties The electrical properties of the thin films were measured using the conventional four probe technique at room temperature. The current vs. voltage curve showed a linear relationship, which demonstrated the ohmic nature of undoped and doped films, as shown in Figure 3.11. The resistivity of TZO1 decreased from 8.2 х 10-3 Ω cm to 1.4 х 10-3 Ω cm as compared to that of the undoped ZnO thin film. The decrease in the resistivity of 80 TZO1 is due to the increase in the crystallinity of the thin film, which is in agreement with the XRD results as discussed above and due to the replacement of divalent Zn+2 with the tetravalent Sn+4 ions, more free electrons are generated for conduction. No further decrease in the resistivity of the thin film was observed with a higher Sn content in TZO2. The increase in the resistivity of TZO2 film is attributed to the segregation of excess Sn dopant at grain boundaries, which act as carrier traps [58]. Shelke et al. [24] observed a similar decrease in resistivity with the incorporation of Sn dopants compared to undoped ZnO film and an increase in the resistivity at higher doping concentration. The minimum resistivity reported by them was 3.11 Ω cm noted for 4 atm % Sn. Nasir et al. [26] reported a minimum resistivity of 3.08 х 103 Ω cm at 4 atm % Sn and Bedia et al. [25] observed a minimum resistivity of 2.1 х 10-2 Ω cm at 1 atm % Sn. The reported minimum resistivity values were much higher than our current study. Hall measurement derived electrical properties of films is shown in Table 3.2, and it is evident that the mobility reached maximum value in TZO1 for which the resistivity was calculated to have a minimum value. 3.4.6.3 Optical properties The optical properties of the thin films were determined in the wavelength range of 300 to 800 nm. Figure 3.12 shows the transmission spectra of the films. The transmission spectra of all the films exhibited an average transmission of 80 %. The absorption edge slightly shifted to a higher wavelength in the TZO1 film compared with the ZnO film. All films exhibited a ripple pattern, which revealed a homogeneous surface and good adhesion on glass substrates. Similar results were obtained for films deposited 81 on glass substrates by spray pyrolysis [25]. Shelke et al. [24] also reported a similar oscillating nature in TZO films with optical transmittance above 85 % in the visible region and observed a decrease in the transmittance at higher doping concentration. Refractive index and extinction coefficient are important parameters for various optoelectronic applications and in integrated optical devices (switches, modulators, filters, etc.). The extraction coefficient defined as the fraction of energy lost due to scattering and absorption per unit thickness in a particular medium was also calculated from [59]: 𝑘=𝛼𝜆/4𝜋 (3.15) The spectral dependence of k in the visible region of spectrum is shown in Figure 3.13 (a). The small values of k obtained in the visible region of the spectrum indicated the high transparency of the prepared thin films. In the case of the TZO2 film, a small increase in the k value at a higher wavelength indicated that the excess Sn segregated to the grain boundaries acted as scattering centers, which corroborated with XRD data as extinction coefficient is directly related to the creation of defects and absorption centers. The refractive index was calculated from the following equation [59]: 𝑛 = (1+𝑅/1−𝑅) + [4𝑅 /(R−1)2 -𝑘2]1/2 where R is the reflectance and k is the extinction coefficient. (3.16) 82 The variation of the refractive index of ZnO and TZO1 with wavelength is shown in Figure 3.13 (b). It is clear that the refractive index increases sharply near the optical absorption edge and then decreases as the wavelength increases. This is due to the increase in transmissivity and decrease in absorption coefficient in the visible region of spectrum. The decrease in refractive index observed in TZO1 as compared to ZnO is due to an increase in the carrier concentration of the film since variation in refractive index is mainly because of the interactions between photons and electrons in the films [60]. 3.4.6.4 Photocurrent properties The photocurrent characteristics of the prepared film is shown in Figure 3.14. Only the prepared TZO1 thin film showed an enhanced increase in current response when illuminated with UV light. The increase in photocurrent is due to the photo excitation of electrons from VB into CB. When the films are illuminated with UV light, photon energy is absorbed. Because the photon energy of UV light is higher than the optical band gap energy of the films, the electron-hole pairs are created, and electrons gain enough energy for excitation leaving behind holes in the valence band [61]. The magnitude of photocurrent obtained was higher at all bias voltages for TZO1 as compared to the I-V spectra for the TZO1 film, which was obtained without UV illumination. The I-V spectra obtained for the ZnO and TZO2 films with and without UV illumination exhibited no significant changes in the photocurrent values. Mammat et al. [61] investigated the photocurrent properties of AZO films at different doping content of Al with and without UV light. It was observed that maximum photocurrent was obtained with 1 % Al, and minimum photocurrent was observed with undoped ZnO film. In our study also, the ZnO 83 thin film exhibited insignificant rise in photocurrent due to its low carrier density, and in the TZO2 film, the low mobility value and more defects arising out of excess Sn affected the photocurrent properties. 3.5 Figure of Merit For applications as transparent contacts, a film must have a low resistivity and high transmissivity in the visible region. The figure of merit defined by Haacke [62] is calculated as the criterion to determine the performance of transparent conducting oxides, as follows: FOM =T10/Rs (3.17) where T is the average transmittance in the visible region, and Rs is the sheet resistance. As shown in Table 3.3, the highest figure of merit 3.2 х 10-3 is obtained for TZO1. High FOM values imply that the prepared films are suitable in various optoelectronic applications. 3.6. Conclusions The thermal plasma process performed in this work has shown its potential as an efficient technique for synthesizing TZO nanopowder. This process is also suitable for large scale production of nano-sized powders due to the availability of high temperature for volatizing reactants rapidly followed by vapor phase reactions and rapid quenching to yield nano-sized powder. XRD results revealed that the doped and undoped ZnO 84 nanopowder are purely crystalline, which belong to the hexagonal wurtzite structure of ZnO. 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Phys., vol. 47, no. 9, pp. 4086-4089, 1976. 90 Table 3.1 Crystallite size calculated from XRD analysis Type SIZE SIZE SIZE (100) nm (002) nm (101) nm Average Crystallite Size (nm) Average Lattice constant ‘a' (Å) ZnO 56.3 59.4 56.2 57.3 3.28 TZO1 21.4 24.7 20.2 22.1 3.27 TZO2 38.5 38.4 38.3 38.4 3.27 Table 3.2 The electrical properties of ZnO and TZO films Magnetc Type of Field Film 0.35 ZnO 0.35 0.35 Carrier Resistivity Mobility (Ωcm) (cm2/V-s) 6.6 х 1019 8.2 х 10-3 11.6 TZO1 7.6 х 1019 1.4 х 10-3 58.7 TZO2 1.8 х 1020 2.3 х 10-3 15.1 Density 3 (cm ) Table 3.3 Figure of merit values for ZnO and TZO films TYPE T10/ Rs (Ω-1) ZnO 5.2 х 10-4 TZO1 3.2 х 10-3 TZO2 1.9 х 10-3 91 Figure 3.1 Schematic diagram of the experimental set-up used for photocatalysis tests. 92 Figure 3.2 XRD diffraction patterns of a) ZnO, b) TZO1, and c) TZO2. All the peaks correspond to the hexagonal wurtzite structure of ZnO. SnO2 peaks are observed in TZO2 sample. 93 Figure 3.3 Raman spectra of a) ZnO, b) TZO1, and c) TZO2. Decrease in the intensity of E2(high) mode in TZO2 sample indicates the decrease in the crystal quality at higher doping amount of 5 atm % Sn. 94 Figure 3.4 XPS spectra. 3.4 (a) XPS Zn 2p core level spectra of a) TZO1, and b) TZO2. The binding energy difference of 23.10 eV between Zn 2p3/2 and Zn 2p1/2 and binding energy positions indicates that Zn in TZO nanopowder exist in +2 oxidization state. 3.4 (b). XPS Sn 3d spectra of a) TZO1, and b) TZO2. The binding energy of 486.6 eV indicates that Sn is incorporated in the form of Sn+4 bonding state from SnO2 and the intensity of Sn 3d spectra increases with an increase in doping amount. 95 Figure 3.5 XPS O1s spectra for a) TZO1, and b) TZO2. The oxygen peak was deconvoluted into three components with different binding energies. A relative increase in intensity of component Ov to OL-2 component indicates that the concentration of oxygen vacancies varies with Sn addition. 96 Figure 3.6 PL spectra for a) ZnO, b) TZO1, and c) TZO2. Oxygen vacancies enhance the intensity of green band emission and the ratio of DLE to UV is larger for TZO1 sample followed by TZO2 and ZnO samples. 97 Figure 3.7 SEM micrographs and EDS spectra. 3.7 (a) SEM micrographs of TZO1 particles synthesized at 10 kW. 3.7 (b) EDS spectrum for TZO1 sample synthesized at 10 kW. Spherical shaped morphology with no cluster formation observed for all the compositions of TZO and for ZnO samples. 98 Figure 3.8 Absorption spectra of methylene blue at different intervals of time using photocatalyst a) ZnO, b) TZO1, and c) TZO2. The degradation of methylene blue is enhanced significantly when TZO1 sample was used as catalyst. 99 Figure 3.9 Photocatalysis and absorbance data. 3.9 (a) Plot showing the linear regression curve fit for the natural logarithm of absorbance of the methylene blue concentration against irradiation time for a) ZnO, b) TZO1, and c) TZO2. The TZO1 sample exhibits superior photocatalytic activity compared to ZnO and TZO2 samples. 3.9 (b) Absorbance spectrum of a) ZnO, b) TZO1, and c) TZO2. ZnO nanoparticles show maximum absorbance at 392 nm corresponding to a band gap of 3.22 eV, which is redshifted to a band gap of 3.16 eV in the case of TZO1. 100 Figure 3.10 XRD diffraction patterns of thin films a) ZnO, b) TZO1, and c) TZO2. Sharp and strong (002) peak corresponds to the hexagonal wurtzite of ZnO. The decrease in intensity at the (002) plane for the TZO2 film indicates that the crystalline quality of thin film is deteriorated at higher doping amount. 101 Figure 3.11 Current-Voltage characteristics of films a) ZnO, b) TZO1, and c) TZO2. The lowest resistivity of 1.4 х 10-3 Ω cm is obtained with 5 % Sn in TZO1 film, and the highest resistivity of 8.2 х 10-3 Ω cm is obtained with undoped ZnO film. Linear fit overlaps with the I-V curves. 102 Figure 3.12 Plot of transmission curves versus wavelength for a) ZnO, b) TZO1, and c) TZO2. All films exhibit an average transmittance of 80 %. 103 Figure 3.13 Optical constants of films. 3.13 (a) Plot of extinction coefficient versus wavelength for a) ZnO, b) TZO1, and c) TZO2. The small values of k obtained in the visible region of the spectrum indicate the high transparency of the prepared thin films. 3.13 (b) The variation of refractive index with wavelength a) ZnO, and b) TZO1. The decrease in refractive index is observed in TZO1 as compared to ZnO is due to an increase in the carrier concentration of the film. 104 Figure 3.14 Current-Voltage characteristic curve of TZO1 film under UV light illumination. An increase in photocurrent values at all bias voltages for TZO1 was obtained as compared to the I-V spectra for the TZO1 film obtained without UV illumination. Linear fit overlaps with the I-V curves. CHAPTER 4 PLASMA-ASSISTED CHEMICAL VAPOR SYNTHESIS OF ALUMINUMDOPED ZINC OXIDE NANOPWDER AND SYNTHESIS OF AZO FILMS FOR OPTOELECTRONIC APPLICATIONS 4.1 Introduction Transparent conducting oxides (TCO) are of great scientific and commercial importance because they combine high conductivity with optical transparency in the visible region. TCOs are commonly used in organic light emitting diodes, transparent transistors, electro-optical devices, and gas sensitive devices. Zinc oxide (ZnO) is a II-VI group compound semiconductor, which crystallize in a wurtzite structure belonging to the space group P63mc. ZnO has the characteristics of high transparency, good UV trapping properties, nontoxicity, and natural abundance, which are important properties of optoelectronic and piezo electronic materials owing to its large band gap of 3.37 eV and large exciton binding energy of 60 meV [1-3]. However, ZnO films have poor conductivity, and doping with various dopants is usually necessary to improve the conductivity for use as TCO films. Doping of ZnO by replacing Zn+2 ions with higher valence ions such as Al+3 and Sn+4 can in general induce dramatic changes in electrical conductivity by increasing the charge carrier density [4]. At present, tin-doped indium oxide (indium tin oxide or ITO) is the most commonly used TCOs, but because of 106 concerns of the supply of world indium reserves and the cost of indium, there has been an increasing interest in alternatives [5]. Aluminum-doped zinc oxide (AZO) is one of the most important alternatives to ITO and is widely used as a transparent electrode in various kinds of devices. When Al is doped into ZnO, Al+3 substitutes Zn+2 sites in the ZnO crystal structure resulting in one free electron to contribute to the electric conduction. The ionic radius of Al+3 (0.053 nm) is smaller than that of Zn+2 (0.074 nm); therefore, Al+3 ions can replace Zn+2 ions in substitution sites. The electrical conductivity, transparency, thermal stability, and durability make this material attractive [6, 7]. In addition, AZO also can be used as a photocatalyst because of their high activity and chemical stability. Various methods have been reported for preparing AZO nanopowder, including sol-gel [8], spray pyrolysis [9], precipitation [10], and hydrothermal processes [11]. In this study, an innovative plasma processing technique has been developed for the preparation of nano-sized AZO powder by vapor phase reactions. Thermal plasma provides a high processing rate as well as other advantages such as good control over size, shape, and crystal structure. Apart from these advantages, a clean reaction atmosphere yields high purity products, a high quench rate to form ultrafine powder, and a wide choice of reactants [12]. The films of TCOs on glass substrates have been prepared by a variety of physical and chemical deposition techniques for transparent and conductive electrodes, albeit with inherently costly and time-consuming processes. Over other techniques, fabricating AZO films from well-dispersed AZO nanoparticles is a good alternative because it has the advantages of low processing cost and easy control of composition. In addition to low cost 107 processing technique, nanoparticles can be handled efficiently, and loss of raw materials could be minimized [13, 14]. In this study, AZO nanopowders were synthesized for the first time by the plasma process using zinc nitrate and aluminum nitrate as the precursors. Further, AZO nanoparticles were dispersed in an organic solvent to produce transparent conducting layers on glass substrates. These layers had a high figure of merit and their resistivity values in the order of 10-3-10-4 Ωcm were much lower than that of other reported values for AZO films [15, 16], indicating their suitability for use as transparent electrodes. It was found that the electrical and optical properties of deposited AZO films are sensitive to the doping amount. 4.2. Experimental Set -Up The plasma reactor system used for the synthesis of aluminum-doped zinc oxide (AZO) nanopowder consisted of a downward plasma torch, a power supply unit, a cylindrical reactor, a powder feeding system, a cooling chamber, cooling system, powder delivery system, powder collectors, a gas delivery system, and an off-gas scrubber, and an off-gas exhaust system. The plasma torch had a copper cathode and a tungsten anode. It was water-cooled and operated at atmospheric pressure. The reactor assembly consisted of a vertical stainless-steel tube, which was water-cooled, and had the dimensions of 15 cm inner diameter and 60 cm length. A graphite cylinder of 7.6 cm inner diameter and 60 cm length was placed inside the stainless-steel tube. A graphite felt filled the gap between the graphite cylinder and the stainless-steel tube. The cooling chamber consisted of two-layer stainless steel to cool the outgoing gases to lower than 150 °C. The precursor was directly 108 fed into the plasma gun using Ar as the carrier gas through a powder feeding system consisting of a test tube filled with the precursor powder, a motor that pushed up the test tube at a constant rate, a carrier gas line that carried the fluidized particles from the top of the particle bed in the test tube, and a vibrator. The product was collected on a Teflon coated polyester filter with a pore size of 1 µm. The off-gas scrubber used a 5 % NaOH solution. More details about the experimental set-up can be found in a previous publication [17- 20]. 4.3 Experimental Procedure The precursors used in this work were 1) Zinc Nitrate (Zn(NO₃)₂•6H₂O;Alfa Aesar, Haverhill, MA), and 2) Aluminum Nitrate (Al(NO₃)₃·9H₂O; Alfa Aesar, Haverhill, MA). Each precursor was ground by mortar and pestle and sieved until the final size approximately 50 µm. The milled precursors were kept in a vacuum oven at 50 °C to remove the moisture for ease of powder feeding. The anhydrous milled precursors were uniformly mixed using a vibrating mixer. The uniformly mixed precursor was fed into the plasma flame through an internal port of 2 mm diameter in the plasma torch, using a specially designed powder feeder system mentioned above. The experimental conditions were as follows: 1) plasma gas flow rate (Ar) of 40 L/min (25 °C and 86.1 kPa total pressure at Salt Lake City), 2) precursor feeding rate of 0.5 g/min, 3) flow rate of carrier gas (Ar) to carry the precursor of 3.5 L/min, and 4) plasma power of 15 kW. The amount of aluminum nitrate was varied to obtain the doping levels of 2, 4, and 8 atm % Al designated as AZO1, AZO2 and AZO3, respectively. 109 4.3.1 Analysis Methods for Nanopowders The synthesized powders collected were then characterized through the X-ray diffraction technique (Rigaku D/Max-2200V) for its structural analysis. The surface morphology of the powder was investigated by a High-Resolution Field Emission Scanning Electron Microscope (Hitachi S-4800) equipped with an Energy Dispersive Spectrophotometer(EDS) system. XPS (Kratos Axis Ultra DLD) was utilized to analyze the chemical state of the AZO nanopowder. Raman scattering spectra were measured by micro Raman spectroscopy (WITec Alpha SNOM) using a He-Ne laser as the excitation source with holographic grating of 1800 grooves/mm. Magnetic properties were analyzed at room temperature using a vibrating sample magnetometer (Microsense FCM-10) with magnetic field ranging from -10000 T to 10000 T. 4.3.2 Preparation of Thin Films AZO nanopowder was mechanically dispersed in ethylene glycol using an ultrasonicator, and carbonic acid was added as dispersion agent. The solution was then added to ethanol. AZO coatings were prepared by a spin coating process on a 2.5 cm ×2.5 cm borosilicate glass substrate using 0.8 ml of the prepared solution, and the speed was set to 2000 rev/min for 60 s. The prepared coatings were annealed at 500 °C in an inert atmosphere for thermal densification. 4.3.3 Analysis Methods for Thin Films AZO film was then analyzed using a Rigaku D/Max-2200V X-ray diffractometer with Cu Kα radiation (λ = 1.5406 Å) from 10.00°to 80.00°at a rate of 0.02°/s. The surface 110 morphology was measured by AFM (Bruker dimension Icon using the ScanAsyst mode). Sheet Resistance of thin film was measured by the four-probe technique and after annealing in hydrogen (H2) atmosphere. Hall effect measurements were carried out using the fourterminal method to minimize Schottky contacts. The optical properties were recorded using a UV-Vis-NIR spectrophotometer (Shimadzu UV-3600) and by RT-PL using He-Cd laser. 4.4 Results and Discussions 4.4.1 X-Ray Diffraction (XRD) Figure 4.1 shows the XRD pattern for AZO nanopowders with different Al concentrations. All the diffraction peaks were indexed using the PDXL software and the peaks confirmed the hexagonal wurtzite structure of ZnO (JCPDS 36-1451) with space group as P63mc. For AZO1 with Al concentration of 2 atm % and AZO2 with Al concentration of 4 atm %, no additional peaks apart from ZnO peaks were observed. In sample AZO3 with 8 atm % Al, there were additional peaks observed and were indexed as gahnite phase ZnAl2O4 (JCPDS 5-0669). In samples AZO1 and AZO2, it was concluded that all aluminum ions were incorporated in the ZnO crystal structure because no Al peaks were observed. In sample AZO3, the excessive Al must have reacted with ZnO to form the gahnite phase implying that the Al content must have exceeded the solubility of Al in ZnO. The reflections (100), (002), and (101) are of high intensity and other reflections (102), (110), (103), and (112) are of lower intensity. Since the (101) reflection is the highest in all the samples, we can conclude the particles are orientated mostly in the [101] direction. The crystallite size was determined by the following Debye-Scherrer equation [21] 111 D = Kλ/(β cos θ) (4.1) where β =√(β2(FWHM) - β20) the peak broadening after removing the instrumental broadening, β (FWHM) is the full width at half maximum and β0 is the correction factor (0.005 rad, determined previously for this instrument). Crystallite size is determined from the three peaks (100), (002), and (101), and the average size is calculated. The lattice constants are calculated by the following equation [22, 23] 𝟏 𝒅(𝒉𝒌𝒍)𝟐 𝟒 𝒉𝟐 +𝒉𝒌+𝒌𝟐 = 𝟑( 𝒂𝟐 )+ 𝒍𝟐 𝒄𝟐 (4.2) where a and c are the lattice constants, and d (hkl) is the crystalline plane distance for indices (hkl) calculated from the XRD pattern. The position of the most intense peak (101) was slightly shifted towards a higher angle as compared to the standard reference of ZnO pattern. This indicated the replacement of larger Zn+2 (72 pm) by smaller Al+3 (53 pm) ions in the ZnO lattice [24, 25]. The crystallite size was increased on increasing the doping amount, and a simultaneous decrease in FWHM was also observed as shown in Table 4.1. 4.4.2 Scanning Electron Microscopy- Energy Dispersive Spectroscopy (SEM-EDS) Figure 4.2 (a) shows typical SEM micrographs of AZO1 particles synthesized at 10kW. Nearly spherical shaped morphology was observed for all compositions of AZO particles. Different morphologies of AZO nanopowders have been reported earlier, such as 112 polyhedral [26], nanorods [27], hexagonal rod like shape [11], and nanowires [28]. Figure 4.2 (b) shows the EDS spectrum for AZO1 synthesized at 10 kW. It indicates the presence of zinc and aluminum elements where the Zn Lα1 and Al Kα1 peaks appear at 1.1 keV and 1.5 keV, respectively. The obtained EDS results for all samples indicate that the elemental distribution of Zn and Al is uniform, and the Al/Zn (atomic %) in the product is close to the initially designed composition of AZO powders. 4.4.3 X-ray Photoelectron Spectroscopy (XPS) XPS analysis was carried out to verify the concentration of Al and to ascertain the valence of Al in ZnO in AZO1 and AZO2. Al 2p core levels were observed at ~ 73.8 eV, which is the characteristic peak attributed to the Al 2p core level doublet, corresponding to the oxidation state close to Al+3 in Al2O3 [29]. This further gives support to the Al atoms substituting Zn atoms, which act as an Al donor. The Zn 2p spectra were located at 1021.42 eV and 1044.34 eV, as shown in Figure 4.3 (a), which correspond to Zn 2p3/2 and Zn 2p1/2, respectively. The binding energy difference of 23.10 eV between Zn 2p3/2 &Zn 2p1/2 and the binding energy position indicate that Zn in AZO nanopowder exist in +2 oxidized state, which is also in agreement with the standard reference value of ZnO. The intensity of Al 2p increased with an increase in doping concentration, as observed in Al 2p plotted spectra, shown in Figure 4.3 (b). The observed Al 2p peak also ruled out the possibility of segregation of metallic Al or the presence of other phases. The XPS results showed an asymmetric shape of O 1s spectra as shown in Figure 4.4, which was deconvoluted into three peaks using Gaussian fitting. The main peak centered at 529.92 eV is associated with the bulk oxygen that refers to O-2 species in the hexagonal structure of ZnO [30]. The 113 shoulder at higher binding energy contains the contributions from two components that are oxygen vacancies or defects and chemisorbed oxygen species. The intermediate binding energy peak (531.4 eV) is associated with the oxygen vacancies, and the higher binding energy peak (532.4 eV) is attributed to chemisorbed oxygen species or hydroxyl groups on the surface of nanopowder [31- 34]. The changes in the intensity of the Ov component are in connection with the variation in the concentration of oxygen vacancies [35]. To further study the effect of hydrogen annealing on the concentration of oxygen vacancies, AZO1 nanopowder was annealed in hydrogen atmosphere at 400 °C for 2 h. The results from XPS O 1s spectra clearly indicate a significant increase in the relative intensity of the Ov component. Thus, hydrogen annealing influences the concentration of oxygen vacancies by producing an oxygen deficient state of the surface after the removal of oxygen atoms. Park et al. [35] reported an increase in the intensity of Ov component upon annealing a ZnO film and attributed it to the increase in the amount of oxygen vacancies. 4.4.4 Raman Spectroscopy ZnO has a wurtzite structure and the symmetry of the ZnO lies in the spatial group of C46v with two formula units in the primate cell. Therefore, its optical properties at the center of Brillouin zone (I-) is represented by the following irreducible relation. I-=1A1+ 2B1+ IE1+ 2E2 (4.3) The polar characteristics of A1 and E1 modes are split into LO (longitudinal) and TO (Transverse) modes, with different frequencies. Modes A1, E1, and E2 are Raman and 114 infrared active, whereas the two B1 modes are infrared and Raman inactive, and are silent modes [36, 37]. A single broad LO peak was observed at 588 cm-1,which is due to the overlapping of A1(LO) and E1(LO) modes. In this study as shown in Figure 4.5, AZO1 exhibited highly intense Raman peaks suggesting improved crystal quality, whereas for AZO3, the reduced intensity of Raman signals indicates that the partial crystallinity of AZO nanopowder was deteriorated at higher doping amount, which is corroborated by the XRD data. The relative intensity ratio of A1 (LO) to E2 (high) increased with Al concentration, suggesting the modification of ZnO crystalline lattice. A small peak shift was observed in the E2 (high) mode with the usual position of ZnO at 435 cm-1, indicating the decreased order of wurtzite phase when Al+3 ions were introduced into the ZnO lattice. Lupan et al. [38] also reported that the Raman signals obtained in AZO nanopowder are highly sensitive to crystalline nature. Yun et al. [39] reported that the aluminum doping in ZnO depresses the E1 modes of ZnO, indicating that the incorporation of Al sites in ZnO are Zn sites. The bands above 800 cm-1 are due to the second order Raman modes. The most prominent second-order features occur in the high-frequency region such as broad, intense peak observed at 1158 cm−1 and correspond to LO overtones and combinations involving LO modes [40, 41]. 4.4.5 Magnetism Measurements of AZO Samples Magnetic properties of AZO1 and AZO2 were recorded at room temperature and associated M(H) data is shown in Figure 4.6. The M(H) curves of AZO2 sample clearly indicate paramagnetic behavior in the high field region and in the low field region and a weak hysteresis is observed. The coexistence of paramagnetism and ferromagnetism has 115 been reported by other authors [42- 44]. AZO1 sample showed a more significant ferromagnetic contribution, evident by its S-shaped curve. The saturation magnetization was attainable in the high field region, which was absent and almost like a linear shape in AZO2. M(H) behavior in the high field region for sample AZO2 was due to the increased occurrence of antiferromagnetic coupling between the Al+3 pairs. At a higher doping concentration, the average distance between Al+3 ions decrease, resulting in an enhancement of antiferromagnetic contribution. Shatwani et al. [45] investigated the influence of Mn doping on the magnetic properties of ZnO nanocrystalline particles and reported the ferromagnetic contribution from the possible intrinsic defect that play a major factor in obtaining room temperature ferromagnetism. The room temperature ferromagnetism mainly arises from two causes: One is intrinsic magnetism, and the other is extrinsic magnetism [46]. Extrinsic magnetism arises from the formation of secondary phases or the formation of clusters of transition elements. Extrinsic magnetism can be ruled out as there was no formation of secondary phases, as observed from the XRD results of AZO1 and AZO2.The intrinsic magnetism in doped ZnO arises from exchange interaction between the local spin polarized electrons of Al+3 ions and the conductive electrons. Consequently, the polarized conductive electrons undergo an exchange interaction with local spin-polarized electrons of other Al+3 ions and thus after successive long exchange interaction, almost all Al+3 ions exhibit same spin direction, resulting in ferromagnetism [47]. The ferromagnetic coupling between the Al dopants and intrinsic defects, such as oxygen vacancies or zinc interstitials, is also responsible for room temperature ferromagnetism [48- 50]. It has been reported that Al doping decreases the concentration of oxygen vacancies [51], which is shown by the reaction 116 Al2O3+Vo•• →2Al•Zn + 3Oox. (4.4) Thus, the weak ferromagnetism observed in AZO2 is due to a decrease in oxygen vacancies because of Al incorporation in ZnO. 4.4.6 Thin Films Analysis 4.4.6.1 X-ray diffraction (XRD) All the films exhibited a sharp peak at near 35.21o, corresponding to the (002) plane of hexagonal wurtzite structure of ZnO. XRD analysis suggest that all the AZO thin films exhibited a c-axis orientation perpendicular to the substrate surface, which is evident from the sharp and strong (002) diffraction peak. It is clear that with an increase in Al content, the crystalline quality of AZO thin film is improved, as shown in Figure 4.7. In AZO3, the decrease in intensity at (002) plane and increase in the intensity of (101) and (004) indicated that excess Al deteriorates the crystalline quality of thin film due to lattice distortion. The lattice distortion is due to excess Al segregating to the noncrystalline regions in the grain boundary, which acts as scattering centers to reduce the preferred c-axis orientation [52]. Deng et al. [53] investigated the properties of heavily Al-doped ZnO thin film by RF magnetron sputtering and observed that the intensity of (002) peak was decreased beyond the optimum doping amount, indicating that an increase in doping concentration deteriorates the crystallinity of the films. 117 4.4.6.2 Electrical properties The I-V characteristics of the AZO thin films at room temperature were measured using the conventional four-probe technique. The I-V curve shows the ohmic behavior [Figure 4.8 (a)]. The Al doping mechanism is as follows [54]: 2Al2O3 ZnO⇒ 4Al•Zn+ 4OOx+ 2O2(g) + 4e− (4.5) This mechanism suggests that a decrease in resistivity with Al doping is due to the replacement of Zn+2 ions by Al+3 ions. Aluminum atoms are incorporated into the ZnO and contribute to the conduction electrons. As is observed, with an increase in the dopant concentration the resistivity of the thin film was decreased and the minimum resistivity of 9.9 х 10-4 Ωcm was obtained with 4 atm % Al. With further increase in dopant concentration (Al 8 atm %), the reduction in resistivity of thin film was not observed. This indicates that the doped Al+3 ions act as electrical dopant at the initial doping concentration, but at a higher doping concentration, the decrease in resistivity is not observed due to the formation of gahnite phase, which no longer contributes to the formation of free electrons. Mousavi et al. [15] observed a similar decrease in resistivity with an increase in Al doping concentration and an increase in the resistivity at higher doping concentration. The minimum resistivity reported by them was 1.7 х 10-2 Ωcm noted for 3 atm % Al. Zhai et al. [16] reported a minimum resistivity of 1.28 х 10-2 Ω cm at 4.9 atm % Al. The reported minimum resistivity values were much higher than from our current study. Hall coefficient, carrier concentration, and carrier mobility were calculated from the current versus voltage response of the AZO thin films and shown in Table 4.2. From 118 the hall measurements it was observed that the Hall coefficient was negative, indicating the presence of negative charge carriers. It is clear that with an increase in doping, the carrier concentration increases due to the substitution of more Al+3 ions at Zn+2 sites of the ZnO structure. For AZO3, the resistivity of the thin film is increased because of decrease in the carrier concentration and mobility values. The lowering in the carrier concentration and mobility is attributed to the ZnO structure deterioration, caused by excessive Al atoms producing defects. The distortion of crystal structure is due to the segregation of Al dopant at grain boundaries, forming the gahnite phase (ZnAl2O4) that acts as scattering centers giving rise to different scattering mechanisms and as carrier traps for charge carriers. This in turn results in a decrease in mobility and modifies the potential barrier for charge transport across the grains [15, 55, 56]. Annealing of AZO1 thin film was carried out in hydrogen gas at 400 °C for 2 h. Hydrogen annealing significantly improved the electrical properties of the thin film. As shown in Figure 4.8 (b), the resistivity of AZO1 decreased to 8.68 х 10-4 Ωcm after annealing in hydrogen, indicating the generation of free charge carriers. The annealing process can be described by the Kroeger-Vink notation to explain the generation of free charge carriers [54]. 2ZnZnx + 2OOx+ H2(g) ⇔ 2ZnZnx+ 2OHO•+ 2e− (4.6) Another reason for the decrease in resistivity of thin film is that hydrogen annealing removes the adsorbed oxygen on the surface. Adsorbed oxygen acts as an electron trap and 119 forms a depletion region. Because the charge transfer is dominated by the tunneling effect, the large barrier affects the mobility of electrons [57 - 59]. O- + H2 = H2O + e− (4.7) Oh et al. [57] reported a decrease in resistivity in the hydrogen-annealed ZnO:Al film from 4.80 х 10-3 Ωcm to 8.30 х 10-4 Ωcm, which is accompanied by an increase in the carrier concentration. Oxygen vacancies formed as shown in the PL spectra of the thin film also acted as an electron donor and permitted the electron movements in the conduction band, thereby improving the conductivity of film. These oxygen vacancies act as electron donors and enhance the transfer of electrons into conduction band, leading to enhanced carrier concentration and low resistivity [60]. 4.4.6.3 Optical properties The optical transmission spectra of AZO thin films are shown in Figure 4.9 (a). The AZO1 and AZO2 thin films exhibited a transmittance of nearly 80 % in the visible region, whereas the AZO3 film showed a transmittance of 70 % in the visible region. The optical absorption coefficient α of a direct band gap semiconductor near the band edge, for a photon energy hν greater than the band-gap energy Eg of the semiconductor, is given by the relation [61, 62] 120 (αhν) = A (hν - Eg)1/N (4.8) where h is Planck's constant and ν is the frequency of the incident photon. The constant N depends on the nature of electronic transition. In the case of AZO films, N is equal to 2 for direct allowed transition. The Tauc plot of (αhν)2 versus energy hν for all the AZO films are shown in Figure 4.9 (b). The band gap energy was obtained by extrapolating the linear plot of the Tauc plot curves to the intercept with the energy axis (at αhν = 0). It was observed that the band gap increased from 3.2 eV to 3.28 eV with increasing Al concentration, as in AZO2 film. The blue shift exhibited on increasing the Al concentration is associated with the Moss Burstein effect (BM) [63 - 65]. According to this well-known effect, the conduction band of the degenerate semiconductor is filled with high carrier concentration, and the lowest valence energy states are blocked, leading to the lifting of the Fermi level into the conduction band and widening of the optical band gap [66]. The results are in good agreement with the results obtained from the Hall effect measurements where the carrier concentration of AZO2 was higher than that of AZO1 film. Kim et al. [64] investigated the structural, optical, and electrical properties of aluminumdoped zinc oxide prepared by radio frequency magnetron sputtering and observed an increase in band gap upon Al doping followed by a decrease in band gap, which was in accord with a decrease in carrier concentration. The expression for the shift in band gap Eopt is given by the following equation [67] 121 Δ Eopt= (h2/8mc*) (3N/𝝅)2/3 (4.9) where N is the carrier concentration; h is the Plank constant; and mc* is the conduction band effective mass. This suggests a direct relationship between the carrier concentration and the BM shift and an inverse relationship between the BM shift and carrier effective masses. In AZO3, there was a decrease in the transmittance observed in the visible range, and the decrease is due to the high Al concentration, which induces a negative influence on the transmittance. At a high doping concentration for the case of AZO3, the decrease in transmittance is due to excess Al atoms segregating to grain boundaries, which can be associated with a decrease in the crystallinity of AZO3. This is in agreement with the findings from the XRD data, which showed a decrease in intensity of the (002) plane orientation, which is a significant parameter for understanding the quality of thin films. Also, the interference fringes that were observed in AZO1 and AZO2 films were suppressed in AZO3, which revealed that the layers in AZO3 lack homogeneity, thereby generating more scattering centers because of rough surface. The decrease in band gap in AZO3 can be again explained from the decrease in carrier concentration, which causes a red-shift. At higher Al doping, the band gap shrinkage is also due to the carrier-carrier and carrier-impurity interaction. This phenomenon is valid only when carrier concentration is increased beyond the Mott critical density [68, 69]. Because the carrier concentration obtained in this study is of the order of 1020 cm-3, less than the critical density, the possibility of increased carrier masses influencing the interaction can be ruled out in AZO3, and the variation in band gap is dependent on the change in the carrier concentration values. 122 All the films display a sharp peak of UV near-band edge (NBE) emission as shown in Figure 4.10. The crystal quality affects the origin of green emission, hence the improvement of crystal quality (reduction in structural defects such as oxygen vacancies) enhances the near band edge emission with reduction or vanishing of the green emission. The ultraviolet emission is the characteristic emission of ZnO and is attributed to the band edge transitions or the exciton recombination [70, 71]. UV near-band edge (NBE) emission is subjected to a blue shift for AZO2 film followed by a red shift for AZO3 film. Blue shift in the UV emission for AZO2 thin film can be attributed to the Burstein-Moss effect. When group III elements such as ‘Al' are doped into ZnO, there will be excess of one donor. The Pauli principle prevents the state from being double occupied. Thus, the Fermi level shifts to higher energy leading to the blue shift in the UV emission [72]. Lo et al. [71] and Hou et al. [73] also reported a similar blue shift in UV emission from the near-band edge emission of ZnO upon increasing the Al doping concentration. The increase in FWHM of the UV peak emission can be accounted for by the broadening of the band edges due to potential fluctuations induced by the high concentration of Al. A green emission peak is observed at 540nm, which arises primarily from intrinsic defects. These intrinsic defects are associated with deep level emissions such as oxygen vacancies and the peaks result from the recombination of a photo generated hole with a single ionized oxygen vacancy [74]. Low intensity indicates low concentration of oxygen vacancies in the as-synthesized thin film and thereby improved crystallinity. 123 4.5 Figure of Merit The high transparency and lowest sheet resistance result in high values of the figure of merit calculated by the formula derived by Haacke [75], as expressed below: FOM =T10/Rs (4.10) where T is the average transmittance in the visible region, and Rs is the sheet resistance. As shown in Table 4.3, the highest figure of merit of 3.4 х 10-3 is obtained for AZO1. High FOM values imply that the prepared films are suitable in various optoelectronic applications. 4.6 Conclusions The thermal plasma process used in this work has shown its potential as an efficient technique for synthesizing AZO nanopowder. This process is also suitable for large scale production of nano-sized powders due to the availability of high temperature for volatizing reactants rapidly followed by vapor phase reactions and rapid quenching to yield nanosized products. XRD results revealed that the synthesized AZO nanopowder are purely crystalline and of the hexagonal wurtzite structure, which was also confirmed by Raman measurements. Binding energies of Zn 2p3/2 and Zn 2p1/2 were observed at 1021.42 eV and 1044.34 eV, and binding energy of Al 2p core level was observed at 73.8 eV, corresponding to Al+3 in Al2O3. Room temperature ferromagnetism was observed in the AZO1 nanopowders, and the contribution of the paramagnetic phase becomes dominant over ferromagnetic phase in the AZO2 nanopowders. AZO films were prepared on the glass 124 substrates by spin coating using dispersion of nanoparticles. The effects of Al doping on the structural, optical, and electrical properties of AZO films were investigated. XRD results indicated that the crystalline properties were deteriorated at higher doping concentrations. The deposited films showed enhanced electric properties with high carrier concentrations and low resistivity. 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The crystal planes are indexed corresponding to the hexagonal wurtzite structure of ZnO (JCPDS 361451). The diffraction peaks of ZnAl2O4 were observed only in AZO3 sample. 134 Figure 4.2 SEM micrographs and EDS spectrum. 4.2 (a) SEM micrographs of AZO1 nanoparticles. Spherically shaped morphology with no cluster formation is observed for all the compositions of AZO samples. 4.2 (b) EDS spectrum for AZO1 synthesized at 10 kW. 135 Figure 4.3 XPS spectra. 4.3 (a) XPS Zn 2p narrow spectra of samples AZO1, and AZO2. The binding energy difference of 23.10 eV between Zn 2p3/2 and Zn 2p1/2 and binding energy position indicate that Zn in AZO nanofpowder exists in +2 oxidized state. 4.3 (b) XPS Al 2p narrow spectra of samples AZO1 and AZO2. The binding energy of Al2p at 73.8 eV indicates that Al is incorporated in the form of Al+3 bonding state from Al2O3, and the intensity of Al 2p spectra increases with an increase in doping amount. 136 Figure 4.4 XPS O1s spectra of samples: a) AZO1, b) AZO2, and c) H2 annealed AZO1. The oxygen peak was deconvoluted into three components with different binding energies. A relative increase in the intensity of component Ov to OL-2 component in H2 annealed AZO1 sample indicates that the concentration of oxygen vacancies increases upon hydrogen annealing. 137 Figure 4.5 Raman spectra of a) AZO1, b) AZO2, and c) AZO3. The intensity of the Raman peaks associated to wurtzite ZnO decrease as the concentration of Al doping increases. 138 Figure 4.6 M(H) data for a) AZO1, and b) AZO2. AZO1 nanoparticles exhibit ferromagnetic behavior, whereas the small kink at around an origin indicates that the AZO2 sample has a small ferromagnetic component along with a significant paramagnetic component. 139 Figure 4.7 XRD diffraction patterns of thin films a) AZO1, b) AZO2, and c) AZO3. Sharp and strong (002) peak correspond to the hexagonal wurtzite of ZnO. 140 Figure 4.8 Electrical Properties of AZO films. 4.8 (a) I-V characteristics of thin films: AZO1, AZO2, and c) AZO3. The lowest resistivity of 9.9 х 10-4 Ωcm was obtained with 4 % Al in AZO2 film, and the highest resistivity of 1.7 х 10-3 Ωcm was obtained at higher doping amount of 8 % Al in AZO3 film. 4.8 (b) Hall Voltage vs. Current for films: AZO1, AZO2, AZO3, and H2 annealed AZO1. The lowest resistivity of 8.68 х 10-4 Ωcm was obtained with H2 annealed AZO1 film which was accompanied by increases in carrier concentration and mobility values. Linear fits obtained overlaps with the current-voltage curves for all films. 141 Figure 4.9 Optical properties of AZO films. 4.9 (a) Transmission curves for AZO1, AZO2, and AZO3. AZO1, and AZO2 films exhibit a maximum transmittance of 85 %, whereas the transmittance decreases to 75 % in ITO3 film. 4.9 (b) Tauc plots to determine the band gap for AZO1, AZO2, and AZO3. Band gap increased from 2.93 eV to 3.28 eV in AZO2 film followed by a decrease to 3.16 eV in AZO3 film. 142 Figure 4.10 PL spectra for a) AZO1, b) AZO2, and c) AZO3. The blue shift in UV emission peak in AZO2 film is followed by red shift in AZO3 film. CHAPTER 5 PHOTOCATALYTIC PROPERTY OF PLASMA-SYNTHESIZED ALUMINUM- DOPED ZINC OXIDE NANOPOWDER 5.1 Introduction Transparent conducting oxides like AZO are currently of great commercial and scientific importance in applications such as organic light emitting diodes, transparent transistors, electro-optical devices and gas sensitive devices. Additionally, oxides like AZO can be used as a photocatalyst because of their high activity and chemical stability. An innovative plasma processing technique has been developed for the preparation of nano-sized powder by vapor phase reactions. Thermal plasma provides a high processing rate as well as other advantages like good control over size, shape, and crystal structure, as well as a clean reaction atmosphere that yields high purity products, a high quench rate to form ultra-fine powder, and a wide choice of reactants [1]. Compared with other methods, it avoids multiple steps like in mechanical milling, sol -gel method and precipitation method and does not require a high liquid volume and surfactants which are involved in a wet chemical process. Dyes are an important source of environmental pollution. It is used in many industrial processes, namely cosmetic, textile, and printing. About 15 % of the world production of dyes is lost during the dying process and is released in textile effluents [2]. 144 Methylene Blue is a cationic dye that is usually used in the textile industry for dying linen, wool and silk. The discharge of large amounts of effluent containing different dyes is harmful to microbes, the aquatic system, and human health. The consequence of colored water is detrimental to environment because color obstructs the sunlight access to aquatic organism and plants, and it diminishes photosynthesis and affects the ecosystem [3, 4]. In recent years, an advanced oxidation process utilizing photocatalyst has attracted a great deal of attention in wastewater treatment because of its high activity, mild reaction conditions, and low energy consumption compared to conventional methods [5, 6]. In this work AZO nanoparticles were successfully synthesized by the plasma process. The obtained nanopowders was tested for its photocatalytic property for the degradation of methylene blue under ultraviolet irradiation and various factors that influence the reaction was studied. The mechanism of photocatalytic degradation is explained as follows [7-9]. When the photocatalyst is irradiated with photons of energy equal to or greater than the band gap energy of the photocatalyst, the electrons get excited from the valence band to the conduction band with the reaction of holes (h+) in the valence band. AZO + hv → eCB- + hVB+ (5.1) The electrons are trapped by the dissolved O2 or by the adsorbed O2 to give rise to superoxide radicals: eCB- + O2 → O2• - (5.2) 145 The superoxide radicals can react with H2O2 to form hydroperoxyl radicals (HO2•) and hydroxyl radicals(OH•) which are strong oxidizing agents to decompose the organic molecules. O2 + H2O → HO2•+ OH• (5.3) The photoinduced holes can be trapped by hydroxyl groups (or H2O) on the photocatalyst surface to yield hydroxyl radicals: hVB+ + OH- → •OHads (5.4) hVB+ + H2O → •OH+H+ (5.5) Thereby organic molecules will get oxidized to yield CO2 and H2O as follows: • OH + organic molecules + O2 → products (CO2 + H2O) (5.6) Meanwhile, the recombination of positive hole and electron will take place, which reduces the photocatalytic activity of prepared AZO photocatalyst. eCB- + hVB+→AZO (5.7) Zhang et al. [9] synthesized AZO conductive nanopowders by coprecipitation method and investigated the relationship between the electrical conductivity and 146 photocatalytic activity. Sun et al. [10] prepared porous AZO nanosheets and reported an enhanced photocatalytic activity compared with commercial P25 TiO2 nanoparticles. Pal et al. [11] synthesized Al doped nanostructured films and tested the films photocatalytic activity towards the Rh- 6G dye decomposition. Wang et al. [12] reported different catalytic efficiencies in AZO nanopowders prepared by hydrothermal-facile template method. Pradhan et al. [13] produced ZnO and Al doped ZnO films by spray pyrolysis and evaluated its photocatalytic activity in the degradation of methyl orange dye under different light sources. Zhang et al. [14] prepared ZnO and AZO nanoparticles via sol-gel combustion method and reported an increase in photocatalytic activity for AZO nanoparticles in the degradation of methyl orange under the irradiation of UV light. However, to the best of our knowledge, the report on the influence of all important factors influencing the photocatalytic performance of AZO nanopowders and the kinetic study has not been reported so far in the literature. Herein, AZO nanoparticles were synthesized for the first time by plasma-assisted chemical vapor synthesis route and photocatalytic performance was investigated for the removal of methylene blue dye. 5.2 AZO Synthesis and Characterization The plasma reactor system used for the synthesis of aluminum-doped zinc oxide (AZO) nanopowder consisted of a downward plasma torch, a power supply unit, a cylindrical reactor, a powder feeding system, cooling system, powder delivery system, powder collectors, a gas delivery system, an off-gas scrubber, and an off-gas exhaust system. The plasma torch had a copper cathode and a tungsten anode. It was water-cooled and operated at atmospheric pressure. The reactor assembly consisted of a vertical 147 stainless-steel tube, which was water-cooled, and had the dimensions of 15 cm inner diameter and 60 cm length. A graphite cylinder of 7.6 cm inner diameter and 60 cm length was placed inside the stainless-steel tube. A graphite felt filled the gap between the graphite cylinder and the stainless-steel tube. The cooling chamber consisted of two-layer stainless steel to cool the outgoing gases to lower than 150 °C. The precursor was directly fed into the plasma gun using Ar as the carrier gas through a powder feeding system consisting of a test tube filled with the precursor powder, a motor that pushed up the test tube at a constant rate, a carrier gas line that carried the fluidized particles from the top of the particle bed in the test tube, and a vibrator. The product was collected on a Teflon coated polyester filter with a pore size of 1 µm. The off-gas scrubber used a 5 % NaOH solution. More details about the experimental set-up can be found in a previous publication [15- 18]. 5.2.1 Synthesis Procedure The precursors used in this work were 1) zinc nitrate [Zn(NO₃)₂•6H₂O; Alfa Aesar, Haverhill, MA] and 2) aluminum nitrate [Al(NO₃)₃·9H₂O; Alfa Aesar, Haverhill, MA]. Each precursor was ground by mortar and pestle and sieved until the final size decreased to approximately 50 µm. The milled precursors were kept in a vacuum oven at 50 °C to remove moisture for ease of powder feeding. The anhydrous milled precursors were uniformly mixed using a vibrating mixer. The uniformly mixed precursor was fed into the plasma flame through an internal port of 2 mm diameter in the plasma torch, using a specially designed powder feeder system as mentioned above. The experimental conditions were as follows: 1) plasma gas flow rate (Ar) of 40 L/min (25 °C and 86.1 kPa total pressure at Salt Lake City), 2) precursor feeding rate of 0.5 g/min and 3) flow rate of carrier gas 148 (Ar) to carry the precursor of 3.5 L/min. The amount of aluminum nitrate was fixed to obtain the doping level of 2 atm % Al in AZO. As an operating variable, the plasma power was set at 10, 20, and 30 kW. To compare the photocatalytic activity of AZO with that of ZnO, zinc oxide was synthesized under the same experimental conditions at plasma power of 10 kW. 5.2.2 Characterization Powder X- Ray diffraction (XRD) was recorded on Rigaku D/Max-2200V with a voltage of 40 kV using Cu-Kα radiation (λ = 1.5406 Å) in 2θ range from 10° to 80° with a step size of 0.050°. The surface morphology of the powder was examined with a HighResolution Field Emission Scanning Electron Microscope (Hitachi S-4800) equipped with an Energy Dispersive Spectrophotometer(EDS) system. The micrographs were taken at voltage of 10 kV. A BET surface area analyzer (Micrometric, ASAP 2010, Norcross, GA), which was degassed at 200 °C for 6 h before the measurement, was used to measure the specific surface area of the produced AZO nanopowders. The absorption spectra of ZnO and AZO with various dopant concentrations were investigated by diffuse reflectance spectroscopy (DRS) in which the scanning was done in the wavelength range of 300 - 800 nm. Photoluminescence (PL) spectra were examined by Perkin Elmer spectrophotometer with a Xe lamp using an excitation wavelength of 350 nm. 5.2.3. Photocatalysis Set-Up Photocatalytic degradation of methylene blue by AZO was carried out in a 1liter volume reaction vessel, as shown by a schematic sketch in Figure 5.1. The reaction vessel 149 is fabricated of borosilicate glass to accommodate the immersion well. The double-walled, quartz immersion well consists of a small diameter inner tube, which extends down the annular space to insure flow of coolant from the bottom of the well upward to an outlet. Cooling water was circulated to control the solution temperature. A 450 W medium pressure quartz mercury vapor lamp was inserted vertically in the immersion well. Approximately 60 % of the radiated energy is in the ultraviolent portion of the spectrum, 35 % in the visible region, and the balance in the infrared range. A 6-foot power cord allowed for lowering the lamp into the well. The bottom of the reactor was flat to allow the use of a magnetic stirrer. The working volume in the reactive area of the lamp was approximately 40 - 50 % of the total volume. The reaction vessel had a 14/20 size taper joint for withdrawing the solution after certain time intervals using sparger tubes and one side arm provision for thermometer insertion. The transformer operating at 115 V and 60 Hz was used to supply the extra voltage and current required to initiate the arc. 5.2.4 Measurement of Photocatalytic Activity The photocatalytic activity of the prepared AZO nanopowder was evaluated by the degradation of a methylene blue solution. Prior to irradiation, the slurry solution was magnetically stirred for 30 min to obtain the equilibrium of adsorption and desorption. In this work, 0.1 g of the AZO nanopowder synthesized at 10 kW plasma power was dispersed in 500 mL of 85 µM methylene blue (MB) solution by stirring with a magnetic stirrer. At certain time intervals, 10 mL of the solution were collected from the slurry suspension and centrifuged at 5000 rpm for 10 min. The centrifuged methylene blue solution was filtered by a Millipore filter to remove particles, and the filtrate was analyzed using a UV-VIS 150 spectrophotometer (Shimadzu UV-3600), scanning from 200 nm to 800 nm, and the maximum absorption intensity at 664 nm was measured. A calibration plot based on BeerLambert's law was established by relating the absorbance to the concentration. The reaction kinetics of the photocatalysis in general follows the following LangmuirHinshelwood (L-H) equation: 𝑹= 𝒅𝑪 𝒅𝒕 =− 𝒌𝒔 𝑲𝑪 (𝟏+𝑲𝑪) (5.8) where ks is the surface reaction rate constant, K is the adsorption coefficient of the reactant, and C is the concentration of methylene blue at any time t. At low concentrations of methylene blue and weak adsorption (KC << 1), Eq. 5.8 is simplified to 𝒅𝑪 = -kt (5.9) Ln (C) = -kt + Ln (C0) (5.10) R= 𝒅𝒕 where k = Kks. The integration of Eq. (5.9) gives: where C0 is the initial concentration of methylene blue, k is the pseudo-first-order reaction rate constant (min-1), and t is the reaction time (min). The reaction rate constant (k) was calculated from the slope of ln(C) vs. time plot. The percentage degradation was determined using the following formula: 151 % degradation = 𝑪𝒊 −𝑪𝒇 𝑪𝒊 х 100 (5.11) where Ci and Cf are the initial and final concentration of the dye. The photocatalytic degradation kinetics of methylene Blue (MB) were investigated by analyzing the following parameters: 1) effect of catalyst amount, 2) effect of H2O2 addition, 3) effect of solution temperature, 4) effect of bubbling of O2, 5) effect of pH of the solution, 6) effect of oxygen vacancies, 7) effect of specific surface area, 8) effect of initial concentration, and 9) effect of doping. 5.3. Results and Discussions 5.3.1 XRD Results Figure 5.2 shows the XRD pattern for AZO nanopowder synthesized at 10 kW. All the diffraction peaks were indexed using the PDXL software, and the peaks confirmed the hexagonal wurtzite structure of ZnO (JCPDS 36-1451) with space group as P63mc. No alumina peaks were observed, indicating that all the aluminum oxide was incorporated in the ZnO crystal structure. AZO samples showed preferred (101) orientation with intense and sharp peaks indicating good crystallinity. 152 5.3.2 Scanning Electron Microscopy- Energy Dispersive Spectroscopy (SEM-EDS) Figure 5.3 shows the typical SEM micrographs of AZO particles synthesized at 10 and 20 kW. The formation of nearly spherical shaped morphology with no cluster formation was observed at all plasma powers of 10, 20, and 30 kW. EDS spectra confirmed the existence of elements Al, O, and Zn in all AZO samples. The particle size of AZO powder was measured using the SEM micrographs based on the manual measurement of the limited number of particles and using the ImageJ software for thresholding the images. At plasma power of 10 kW, the average particle size obtained was 30 nm, and the particle size increased with increasing plasma power. Average particle sizes of 50 nm and 130 nm were obtained at plasma power of 20 and 30 kW, respectively. The surface area of the AZO powder produced with varying plasma power was measured using BET to be 48.5 m2/g (at 10 kW), 18.5 m2/g (at 20 kW), and 8.8 m2/g (at 30 kW). The average size of particles can be obtained from the specific surface area based on the following equation under the assumption of spherical particles: 𝑫 = 𝟔/(𝑨) (5.12) where D is the average diameter of the AZO particles produced, A is the specific surface of the powder (m2/g), and  is the density of AZO powder. The average particle size of AZO thus determined was close to average particle size values that were observed from the SEM micrographs. This means that the particles were well dispersed and not agglomerated. 153 5.3.3 Effect of Catalyst Amount The effect of catalyst loading on the photo-degradation of methylene blue was determined by varying the amount of AZO catalyst at 0.050 g, 0.1 g, and 0.5 g in 500 mL of 85 µM dye solution under UV light at pH 7.5 for 120 min. The obtained results are shown in Figure 5.4. With increasing amount of AZO from 0.050 g to 0.1 g, the reaction rate constant increased but with further increase in the catalyst amount, it decreased. The initial increase in rate constant from 0.0244 min-1 to 0.0303 min-1 and the corresponding increase in the degradation of methylene blue from 94.2 % to 97.3 % were because of the increase in the number of active sites on AZO surface. When the amount of AZO was increased further to 0.5 g, the reaction rate constant decreased to 0.202 min-1 and the percentage degradation was 91.3 %. The decrease in reaction rate constant and percentage degradation of methylene blue at a higher dosage of 0.5 g were due to the fact that the opacity of the aqueous medium is decreased and some AZO particles did not receive enough energy for the methylene blue oxidation [19- 20]. Pouretedal et al. [21] investigated the photodegradation of methylene blue and safranin at different dosages of zinc sulfide nanocatalyst doped with manganese, nickel and copper and also observed a decrease in the degradation rate as the loading was increased beyond the optimum amount. Also, at these higher dosages, more light is reflected by the excess AZO nanoparticles and the chances of agglomeration and sedimentation of AZO nanoparticles become greater, leading to a net decrease in the interfacial area involved in the reaction and in the number of active sites [22- 25]. Therefore, the most effective photo-degradation rate of methylene blue is observed at 0.1 g in 500 mL of solution. 154 5.3.4 Effect of H2O2 Addition The effect of H2O2 addition was investigated by using two different amounts of H2O2, 10 mmol, and 30 mmol, with the catalyst amount fixed at 0.050 g in 500 mL of 85 µM dye solution under UV irradiation at pH 7.5 for 80 min. The results obtained are shown in Figure 5.5. The addition of H2O2, a strong electron acceptor, had a positive effect on the degradation of methylene blue. The reaction rate increased from 0.0232 min-1 without the addition of H2O2 to 0.0356 min-1 with the addition of 10 mmol H2O2. The percentage degradation increased from 89.5 % to 98.0 %, indicating that the presence of H2O2 increased the concentration of hydroxyl radicals, thereby increasing the rate of photocatalytic degradation of methylene blue. The addition of H2O2 also inhibits the electron-hole recombination, which is shown by the following equation [26- 28]: H2O2 + e− → ·OH + OH− (5.13) H2O2 also reacts with superoxide anions to form hydroxyl radicals, as follows: H2O2 + O2·− → ·OH + OH− + O2 (5.14) However, with a higher concentration of H2O2 at 30 mmol, it was observed that there was a slight decrease in the photocatalytic degradation efficiency of AZO. The degradation of methylene blue decreased to about 92.7 % and a decline in the rate constant from 0.0356 min-1 at 10 mmol to 0.0325 min-1 at 30 mmol, still a significantly higher k than one without 155 the addition of H2O2. The excess H2O2 consumes hydroxyl radicals by the following two reactions [29] H2O2 + ·OH → H2O + OH2· (5.15) HO2· + ·OH → H2O + O2 (5.16) OH2· has a lower oxidizing ability than hydroxyl radicals, and exceeding the optimum dosage of H2O2 can scavenge the •OH radicals [30, 31]. •OH and hVB+ are strong oxidants for organic pollutants, and the photocatalytic degradation of methylene blue is inhibited under excess H2O2. 5.3.5 Effect of Temperature The effect of increasing temperature on the degradation of methylene blue shown in Figure 5.6 was determined by varying the temperature from 25 °C to 70 °C while keeping the catalyst amount fixed at 0.050 g in 500 mL of 85 µM methylene blue solutions for 150 min. A significant increase in the rate constant from 0.0242 min-1 at 25 °C to 0.0307 min-1 was observed at 50 °C, as shown in Figure 5.7. The percentage degradation of methylene blue also increased to 98.7 % at 50 °C from 96.3 % at room temperature (25 °C). An increase in temperature helps overcome the electron-hole recombination. The reaction rate exhibited no further increase at high temperature of 70 °C, since there is not only an increase in the oxidation rate of organic compounds but at the same time a negative effect on the solubility of dissolved oxygen [33]. Zhou et al. [32] reported that increasing the reaction temperature increased the oxidation rate of organic compounds at the interface, 156 but it also reduced the adsorptive capacities associated with the organics and dissolved oxygen. Dissolved oxygen is important to maintain the concentration of ·OH radicals. 5.3.6 Effect of Bubbling O2 In this work, two experiments were carried out with and without O2 gas bubbling, as shown in Figure 5.8. The dissolved oxygen is important for the degradation of methylene blue, as molecular oxygen adsorbed on the surface of photo catalyst reacts with the electrons forming O2·− radicals, as follows: O2+ e− → O2·− (5.17) These radicals are an important oxidant species involved in the degradation of dye molecules and help slow down the recombination rate of electrons and holes by serving as electron acceptors [34-36]. Zhang et al. [7] reported a 12.7 % improvement in the degradation rate of methylene orange (MO) upon bubbling air, demonstrating that bubbling air is able to considerably promote the efficiency of MO photocatalytic degradation. From the results obtained in our study, the degradation rate constant and the efficiency of degradation were not altered much from that of the experiment without O2 gas bubbling. A small increase in the degradation rate constant from 0.0239 min-1 to 0.0289 min-1 and in percentage degradation from 96.2% to 97.1 % indicated the existence of enough dissolved oxygen in the solution even without the O2 gas bubbling. Excess O2 is redundant since the e− concentration is restrained as the amount of catalyst is fixed and so were other parameters like the concentration of dye and temperature. 157 5.3.7 Effect of pH The effect of pH shown in Figure 5.9 was investigated at different media by keeping the catalyst amount and the concentration of methylene blue fixed, respectively, at 0.050 g per 500 mL and 85 µM and at room temperature. The pH was varied by using small amounts of HCl and NaOH. The degradation rate of methylene blue increased with pH of the solution around pH 10. At pH 13, the degradation rate of methylene blue is decreased compared with at pH 10. In an alkaline medium, there is more generation of hydroxyl radicals due to the presence of more hydroxyl ions, by the following reaction: h+VB + OH− → •OH (5.18) The photocatalytic oxidation of hydroxyl ions leads to more hydroxyl radicals, which are the main oxidizing species involved in the degradation of methylene blue. The percentage degradation of methylene blue increased from 96.8 % at pH 3 to 98.2 % at pH 10. A significant increase in the rate constant from 0.0297 min-1 to 0.0419 min-1 was observed with an increase in pH of the solution from 3 to 10. At pH 3, there is more electrostatic repulsion between the cations of the dye molecule and the positively charged AZO surface, leading to less effective adsorption. A subsequent decrease in the rate constant and percentage degradation was observed at pH 13, indicating that degradation of methylene blue is not fast and favorable in a highly alkaline medium. The variation in pH value influences the charge distribution on the AZO surface. The increase in pH makes the AZO surface negatively charged due to the presence of more hydroxyl ions, and as a result there is strong adsorption of cationic methylene blue molecule on the photo catalyst surface. 158 Therefore, the maximum degradation rate of methylene blue is observed at pH 10. In a highly alkaline medium (at pH 13), reduction in degradation rate is due to excessive repulsion between hydroxyl ions and negatively charged AZO surface. This finding is consistent with the previously reported results [37- 39]. At a higher pH, also high adsorption of methylene blue dye molecules on the surface of the catalyst leads to the inhibition of light penetration [40], leading to a decrease in photocatalytic oxidation rate of hydroxyl ions. 5.3.8 Effect of Oxygen Vacancies The photocatalytic degradation of methylene blue was investigated with assynthesized and hydrogen annealed AZO nanopowder shown in Figure 5.10. The annealing of AZO nanopowder was carried out under hydrogen gas for 2 h at 400 °C. The annealed nanopowder showed a slightly higher degradation rate constant, and the percentage degradation of methylene blue also marginally increased. Hydrogen annealing induces oxygen vacancies, according to [41, 42]. 2ZnZnx + 2OOx+ H2(g) ⇔ 2ZnZnx+ 2OHO•+ 2e− 2OHO•⇔ Oxo + V••o + H2O (5.19) (5.20) Oxygen vacancies reduce the recombination rate of electrons and holes. These vacancies act as electron acceptors during the photocatalytic reaction, thereby ensuring that holes in the valence band react with the hydroxyl ions to produce hydroxyl radicals, which are the main oxidant species [43, 44]. Upon UV light irradiation, electrons get excited from VB 159 into oxygen vacancies, and electrons trapped in these vacancies subsequently react with adsorbed oxygen on the AZO surface to generate the oxygen radicals, which in turn produces peroxide ions that oxidize the methylene dye molecules as discussed under reaction mechanism. An increase in the degradation rate constant from 0.0239 min-1 to 0.0279 min-1 and in percentage degradation from 96.2 % to 97.6 % was observed when H2 annealed catalyst was used. This indicates that oxygen vacancies in AZO somewhat influenced the photocatalytic degradation of methylene blue. Pal et al. [11] also reported the highest rate of photodecomposition in AZO sample that had the highest presence of defects (oxygen vacancies). 5.3.9 Effect of Specific Surface Area The effect of specific surface area of the particles on the photocatalytic degradation of methylene blue were studied by keeping the initial methylene blue concentration at 85 µM in 500 mL of solution, the catalyst amount at 0.050 g, and temperature kept constant at 25 °C. There is a significant increase in the rate constant and percentage degradation of methylene blue when particles (synthesized at 20 kW) having specific surface area of 18.5 m2/g were used over particles (synthesized at 30 Kw) with 8.8 m2/g as specific surface area. The percentage degradation of methylene blue increased from 80.0 % to 89.5 %, and the increase in the rate constant was significant from 0.0159 min-1 to 0.0205 min-1. A further increase in rate constant to 0.0238 min-1 was observed when particles (synthesized at 10 kW) having specific surface area of 48.5m2/g were used. The increase in rate constant with decreasing particle size is because of the increase in active sites on the surface [45, 46]. The results are shown in Figure 5.11. 160 5.3.10 Effect of Initial Concentration The initial concentration of dye was varied while keeping the amount of catalyst constant at 0.050 g per 500 mL and temperature fixed at 25 °C. The effect of initial concentration was investigated at three different levels (85 µM, 130 µM and 170 µM). There was an appreciable decrease in the degradation rate constant of methylene blue from 0.0244 min-1 to 0.0173 min-1 when the concentration of the solution was increased to 130 µM. The maximum degradation rate was observed at lower concentration; in this work it was at 85 µM. Further decrease in the percentage degradation and rate constant were recorded at much higher concentration of dye at 170 µM. It is known that degradation rate constant does not vary with the initial concentration according to the first- order reaction law. However, the rate constant obtained vary with initial concentrations as shown in the Figure 5.12. The intensity of UV light, which is an important factor affecting the photocatalytic degradation rates, does not appear in the mentioned first- order kinetics and thereby leading to the variation in the rate constant. The penetration of UV light is important for the electron excitation form VB to CB and for the generation of hydroxyl radicals from the presence of holes in the VB. As the initial concentration of dye molecules increases, more and more dye molecules are adsorbed on the surface of the catalyst and significant amount of UV light is absorbed by the dye molecules rather than AZO nanoparticles [47]. Also, at higher concentration, the requirement for more catalyst surface increases and because the amount of catalyst is fixed, the amount of •OH radicals generated is less in comparison to the dye molecules. Therefore, relative number of radicals oxidizing the dye molecules decreases when the initial dye concentration of solution is increased. 161 Zhang et al. [9] also noted that the degradation of the methyl orange (MO) using AZO catalyst was enhanced with the decrease in the initial concentration of MO. 5.3.11 Effect of Doping The photocatalytic activity of ZnO nanopowder was investigated using MB as an indicator under UV light. The results demonstrated that AZO is superior to ZnO under identical conditions. From Figure 5.13, it is clear that the degradation of MB is enhanced significantly when catalyzed by AZO. The percentage degradation increased from 42.1 % to 96.3 %, and the degradation rate constant increased from 0.0106 min-1 to 0.0244 min-1 upon using AZO nanopowder as photocatalyst. The enhanced photocatalytic efficiency of AZO is attributed to the increase in the specific surface area and band gap of the nanopowder. Upon doping Al in ZnO, the BET specific surface area increased from 27.1 m2/g to 48.6 m2/g. The enhanced S/V ratio facilitated the increase in the number of the active sites on the photocatalyst surface and thereby increased the concentration of photo generated carriers. Thus, the relative number of free radicals attacking the dye molecules increases. The PL spectra of ZnO and AZO nanopowder are shown in Figure 5.14. All of the spectra showed UV emission peak and a green emission peak (deep-level emission) in the visible region. The UV emission also known as near-band edge emission arises from the band edge recombination of free excitations, and green band emission peak is related to the singly ionized oxygen vacancy in ZnO, and results specifically from the recombination of a photo generated hole with singly ionized charge state of this defect [48, 49]. It is evident that the ratio of DLE to UV is higher for AZO than for ZnO nanopowder. 162 Also, from the absorption spectra, the optical band gap was determined from Tauc's relationship [50]. (αhν) = A (hν - Eg)1/N (5.21) The Tauc plot of (αhν)2 versus energy hν is shown in Figure 5.15 assuming direct transition between the edges of the valence and the conduction band. The band gap energy was obtained by extrapolating the linear plot of the Tauc plot curves to the intercept with the energy axis (at αhν= 0). The calculated band gap from Tauc's plot is 3.06 eV for ZnO and 2.91 eV for AZO. The red shift observed with lower band energy in doped sample was also observed for other doped ZnO samples [51, 52]. This indicates that AZO nanoparticles have stronger absorption than ZnO nanoparticles, that is AZO nanoparticles possesses higher photocatalytic activity than undoped ZnO nanoparticles. Oxygen vacancy is also one of the important factors for narrowing the band gap, as observed above in conjunction with Figure 5.15. Ansari et al. [53] suggested that the band tail was formed below the CB of ZnO due to O-vacancy level energies, and Wang et al. [54] reported that an impurity level was created by O-vacancy level energies nearer to the valence band than the undoped ZnO nanopowder. In AZO samples, the replacement of Zn+2 ions by Al+3 ions generate more free electrons leading to an increase in the carrier concentration, and during the UV light irradiation, these free electrons are easily removed from the surface of AZO resulting in the appearance of positively-charged holes [9]. The photocatalytic activity of AZO is enhanced by the efficient transfer of the free electrons to dissolved oxygen molecules to 163 form superoxide anion radicals and consequently suppressing the recombination of photogenerated carriers [55]. 5.4 Conclusions The thermal plasma process performed in this work has shown its potential as an efficient technique for synthesizing AZO nanopowder. This process is also suitable for large scale production of nano-sized powders due to the availability of high temperature for volatizing reactants rapidly followed by vapor phase reactions and rapid quenching to yield nano-sized powders. In this work, methylene blue was degraded by the synthesized AZO photocatalyst. The kinetics of photodegradation of methylene blue dye showed first order kinetics. The effects of various parameters on the photocatalytic activity of AZO photocatalyst were studied. The results showed that the optimum catalyst loading was 0.1 g in 500 mL of 85 µM dye solution. With respect to the initial dye concentration, it is observed that at higher initial dye concentration, degradation rate is decreased due to weakening of light intensity, and the initial concentration had a strong influence on the degradation of methylene blue. The degradation rate increased with an increase in catalyst surface area and optimum concentration of H2O2 for enhanced degradation was 10 mmol. Photodegradation efficiency is favored at pH 10 and at temperature of 70 °C. The photocatalysis study also revealed that AZO nanoparticles had higher photocatalytic activity than ZnO. 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Cho, "Highly visible light active Ag@ZnO nanocomposites synthesized by gel-combustion route," J. Ind. Eng. Chem., vol. 20, no. 4, pp. 1602-1607, 2014. 54. C. Wang, D. Wu, P. Wang, Y. Ao, J. Hou, and J. Qian, "Effect of oxygen vacancy on enhanced photocatalytic activity of reduced ZnO nanorod arrays," Appl. Surf. Sci., vol. 325, pp. 112-116, 2015. 55. R. Mahdavi and S. S. A. Talesh, "Sol-gel synthesis, structural and enhanced photocatalytic performance of Al doped ZnO nanoparticles," Adv. Powder Technol., vol. 28, no. 5, pp. 1418-1425, 2017. 169 Figure 5.1 Schematic diagram of the experimental set- up used for photocatalysis tests. 170 Figure 5.2 XRD pattern of AZO synthesized at 10 kW. The crystal planes are indexed corresponding to the hexagonal wurtzite structure of ZnO (JCPDS 36-1451). 171 Figure 5.3 SEM images of AZO nanopowder synthesized under a) 10 kW and b) 20 kW. Spherical shaped morphology with no cluster formation were observed at all plasma powers of 10, 20, and 30 kW. 172 Figure 5.4 Effect of catalyst loading on the degradation of methylene blue, MB concentration - 85 µM, pH -7.5, temperature- 25 oC and specific surface area (SSA)- 48.5 m2/g. 173 Figure 5.5 Effect of addition of H2O2 concentration on the photodegradation of methylene blue, MB concentration - 85 µM, catalyst amount- 0.050 g/500 mL, pH - 7.5, temperature - 25oC, and SSA- 48.5 m2/g. 174 Figure 5.6 Effect of temperature on the photodegradation of methylene blue, MB concentration - 85 µM, catalyst amount- 0.050 g/500 mL, pH - 7.5, and SSA- 48.5 m2/g. 175 Figure 5.7 Plot of ln (k) vs 1/T. Degradation rate increased with the increase in temperature from 25°C to 50°C, and no further increase was observed at temperature of 70°C. 176 Figure 5.8 Effect of oxygen bubbling on the photodegradation of methylene blue, MB concentration- 85 µM, catalyst amount- 0.050 g/500 mL, pH - 7.5, temperature - 25oC, and SSA- 48.5 m2/g. 177 Figure 5.9 Effect of solution pH on the photodegradation of methylene blue, MB concentration-85 µM, catalyst amount- 0.050 g/500 mL, temperature - 25oC, and SSA48.5 m2/g. 178 Figure 5.10 Effect of using hydrogen annealed AZO photocatalyst on the photodegradation of methylene blue, MB concentration- 85 µM, catalyst amount- 0.050 g/500 mL, pH - 7.5, temperature - 25oC, and SSA- 48.5 m2/g. 179 Figure 5.11 Effect of specific surface area of AZO photocatalyst on the photodegradation of methylene blue, MB concentration - 85 µM, catalyst amount- 0.050 g/500 mL, pH - 7.5, and temperature - 25oC. 180 Figure 5.12 Effect of various initial concentrations on the photodegradation of methylene blue, catalyst amount- 0.050 g/500 mL, pH - 7.5, temperature - 25oC, and SSA- 48.5 m2/g. 181 Figure 5.13 Absorption spectra of methylene blue at different intervals of time using photocatalyst a) ZnO and b) AZO. The degradation of methylene blue is enhanced significantly when AZO nanoparticles were used as catalyst. 182 Figure 5.14 PL spectrum for a) ZnO and b) AZO. Oxygen vacancies enhance the intensity of green band emission and the ratio of DLE to UV is larger in AZO sample than in ZnO sample. 183 Figure 5.15 Tauc's plot for a) ZnO and b) AZO. Band gap decreased from 3.06 eV in ZnO to 2.91 eV in AZO. CHAPTER 6 CONCLUSIONS The thermal plasma process performed in this work has shown its potential as an efficient technique for synthesizing ITO nanopowder. This process is also suitable for large scale production of nano-sized powders due to the availability of high temperatures for volatizing reactants rapidly, followed by vapor phase reactions and rapid quenching to yield nano-sized powder. The experimental results showed that the grain size of ITO nanopowder increased with an increase in plasma torch power, and also product phases were affected by the plasma torch power. Grain size decreased with an increase in plasma gas flow rate within the range tested. ITO films were prepared on glass substrates by spin coating using ITO sol solution. Optical transmittance in the visible region approached 85 % in ITO1 and ITO2 films. The deposited films showed enhanced electric properties with resistivity in the order of 10-3 -10-4 Ωcm, and in particular ITO2 showed the best performance: a high carrier concentration of 5.5 х 1020 cm-3 and a low electrical resistivity of 6.65 х 10-4 Ωcm. The ITO1 film exhibited good sensitivity to H2 gas, and the sensitivity increased with increases in gas concentration and temperature and reached maximum with 400 ppm of H2 gas at an operating temperature of 350 °C. For the synthesis of ZnO and TZO nanopowders, XRD results revealed that the doped and undoped ZnO nanopowders are purely crystalline, which belong to the 185 hexagonal wurtzite structure of ZnO. The presence of tin oxide peaks in TZO2 indicates that the Sn content has exceeded the maximum solubility in ZnO. TZO1 (3 atm % Sn) exhibited higher photocatalytic activity than ZnO in the degradation of methylene blue due to increased specific surface area and higher oxygen vacancy. ZnO and TZO films were prepared on glass substrates by spin coating using dispersion of nanoparticles. The deposited films showed enhanced electric properties with resistivity in the order of 10 -3 Ωcm, and in particular TZO1 showed the best characteristics: a high carrier concentration of 7.6 х 1019 cm-3 and a low electrical resistivity of 1.4 х 10-3 Ωcm, and the optical transmission approached 80 % in the visible range for all the films, thus making it suitable for optoelectronic applications The thermal plasma process performed in this work has also shown its potential as an efficient technique for synthesizing AZO nanopowder. XRD results revealed that the AZO nanopowders are purely crystalline, which belong to the hexagonal wurtzite structure of ZnO and is also confirmed by Raman measurements. Binding energies of Zn 2p3/2 and Zn 2p1/2 were observed at 1021.42 eV and 1044.34 eV, and binding energy of Al 2p core level was observed at 73.8 eV, corresponding to Al+3 in Al2O3. Room temperature ferromagnetism is observed in AZO1 nanopowder, and contribution of paramagnetic phase becomes dominant over ferromagnetic phase in AZO2 nanopowder. AZO films were prepared on glass substrates by spin coating using dispersion of nanoparticles. The effects of Al doping on the structural, optical, and electrical properties of AZO films were investigated. XRD results indicated that the crystalline properties were deteriorated at higher doping concentration. The deposited films showed enhanced electric properties with high carrier concentration and low resistivity, and in particular AZO2 showed the best 186 performance: a high carrier concentration of 7.20 х 1020 cm-3 and a low electrical resistivity of 9.90 х 10-4 Ωcm. Hydrogen annealing of AZO1 film enhanced the electrical properties with high carrier concentration of 2.24 х 1020 cm-3 and mobility value of 32 cm2 /V-s. Optical transmittance in the visible region approached 80 % in AZO1 and AZO2 films. The blue shift in the band gap from 3.2 eV in AZO1 film to 3.28 eV in AZO2 film is associated with Burstein Moss effect, and band shrinkage observed in AZO3 film is due to the decrease in carrier concentration. In this work methylene blue was degraded by the synthesized AZO photocatalyst. The kinetics of photodegradation of methylene blue dye showed first order kinetics. The effects of various parameters on the photocatalytic activity of AZO photocatalyst were studied. The results showed that the optimum catalyst loading was 0.1 g in 500 mL of 85 µM dye solution. With respect to the initial dye concentration, it is observed that at higher initial dye concentration, degradation rate is decreased due to weakening of light intensity, and the initial concentration had a strong influence on the degradation of methylene blue. The degradation rate increased with an increase in catalyst surface area, and optimum concentration of H2O2 for enhanced degradation was 10 mmol. Photodegradation efficiency is favored at pH 10 and at temperature of 70 °C. The photocatalysis study also revealed that AZO nanoparticles had higher photocatalytic activity than ZnO. The enhanced photocatalytic activity in AZO is ascribed to the increase in specific surface area, enhanced oxygen vacancies as revealed from PL spectra and stronger absorption.