Keywords: chalcogenide glasses, glass ceramic, heterojunction, photovoltaic effect, photocatalysis

Conclusions

Ce premier travail a permis d'établir des relations entre la composition, la microstructure et les propretés photoélectriques des vitrocéramiques dans le système GeSe 2 -Sb 2 Se 3 -CuI. Il a été démontré que les excellentes propriétés de certaines vitrocéramiques sont attribuables à une microstructure inédite contenant des heterojunctions p-n formées par des cristaux de Sb 2 Se 3 dopé par l'iode (semi-conducteur du type n) et des cristaux Cu 2 GeSe 3 (type p). Ces deux cristaux possèdent un gap électronique direct avec une fore absorption de la lumière visible. Ces jonctions favorisent la séparation de charges photogénérées, conduisant à une durée de vie exceptionnellement longue (environ 16 µs) de ces charges. En plus, les cristaux Sb 2 Se 3 orthorhombique sont sous forme d'aiguilles, très efficaces pour former un réseau tridimensionnel interconnecté. Ce réseau est rendu conducteur grâce à la croissance de la phase cubique Cu 2 GeSe 3 , ayant une bonne conductivité, sur les Sb 2 Se 3 nettement moins conducteurs. Cette structure peut donc transporter facilement les charges.

Les matériaux/structures permettant d'absorber efficacement la lumière visible, de séparer et de transporter les charges photogénérées, sont recherchés pour exploiter l'énergie solaire par effet photovoltaïque ou par photocatalyse. Cette étude préliminaire a permis de démontrer l'efficacité de ces nouveaux matériaux/structures pour la photocatalyse. D'autres travaux sont indispensables pour encore mieux comprendre ces matériaux et pour fabriquer des dispositifs afin de démontrer expérimentalement leur efficacité pour convertir de l'énergie solaire.

Introduction

La crise de l'énergie et les problèmes environnementaux sont parmi les grands défis de l'humanité. Une exploitation efficace de l'énergie solaire, grâce à son abondance sur la Terre, pourrait permettre de résoudre les problèmes liés à l'énergie.

L'effet photovoltaïque et la photo-catalyse permettent de transformer directement l'énergie solaire en électricité ou en combustibles chimiques. Ces deux technologies sont parmi les étudiées pour utiliser l'énergie solaire. Elles nécessitent des matériaux semi -conducteurs capables d'absorber efficacement la lumière solaire en générant des paires électron-trou qui doivent avoir une durée de vie suffisamment longue.

Ce travail de thèse, inédit dans le laboratoire des verres et céramiques, est la suite inattendue d'un travail antérieur consacré au développement des vitrocéramiques dans le système GeSe 2 -Sb 2 Se 3 -CuI. L'objectif était de développer des matériaux transparents dans l'infrarouge ayant une résistance mécanique nettement supérieure au verre grâce à une cristallisation contrôlée du verre. Il a été découvert que ces vitrocéramiques possèdent des propriétés optiques inhabituelles et inintéressantes pour l'application envisagée, puisque la transmission diminue quand la longueur d'onde augmente, comme le montre la figure 1. Ces matériaux sont restés dans les « tiroirs » pendant quelques années malgré le fait que ce phénomène déjà connu dans certain semi-conducteurs a été attribué à la présence de porteurs de charge libres dans le matériau. Ce travail de thèse a débuté lorsque ces matériaux ont été testés, par curiosité, par des mesures de photo-courant sous illumination visible.

2 Optimisation de composition dans le système GeSe 2 -Ga 2 Se 3 -CuI Motivés par les premiers résultats intéressants avec un échantillon sélectionné par hasard, nous avons exploré systématiquement ce système ternaire (figure 2) avec la procédure suivante : Plusieurs séries de verres, illustrées dans ce domaine vitreux ont été préparés par la méthode de synthèse couramment utilisée pour les verres de chalcogénure. Elle consiste à mettre les matières de départ dans un tube de silice qui est ensuite scellé sous vide. Ce tube est chauffé dans un four basculant pendant environ 12 heures à une température entre 750-850°C selon la composition. Il est ensuite trempé dans l'eau à la température ambiante. Les températures caractéristiques, notamment la température de transition vitreuse et la ou les températures de cristallisation de tous les verres ont été déterminées par calorimétrie différentielle à balayage (DSC). Il peut être constaté que tous les pics de diffraction de la vitrocéramique obtenue à partir du verre de base sans CuI appartiennent à Sb 2 Se 3 , avec cependant une intensité relative de chaque pic différente par rapport au PDF standard (n ° 01-072-1184). Une étude plus détaillée est nécessaire pour comprendre les raisons. Il est vrai que la cristallisation d'une phase amorphe dans un environnement assez complexe à faible mobilité des atomes peut être très différente de la cristallisation d'un bain fondu avec la même composition.

Différentes propriétés de cette série de vitrocéramiques ont été mesurées et rassemblées dans le tableau ci-dessous (tableau I). Il est intéressant de noter que les deux vitrocéramiques avec 20% et 30% de CuI donnent un photocourant important. Si on examine ces deux mesures de photocourant plus en détail (figure 5), une grande différence existe. En effet, même si les deux vitrocéramiques ont été mesurées comme des semi-conducteurs du type N, la vitrocéramique avec 20% CuI montre. 3 Optimisation de la cristallisation du verre 40GeSe 2 -40Sb 2 Se 3 -20CuI

Compte tenu de ces résultats, l étude est ensuite concentrée sur la composition 40GeSe 2 -40Sb 2 Se 3 -20CuI. L'étude préliminaire montre également que la cristallisation a une influence déterminante sur les propriétés photoélectriques. Afin de mieux comprendre et de mieux maitriser ce processus, une étude dynamique a été effectuée à l'aide de la calorimétrie et a permis de déterminer des paramètres importants dans le contrôle de ce processus de cristallisation. L'exemple ici est donné par la détermination de la température et du temps optimaux pour la nucléation (figure 6).

Deux stratégies de cristallisation ont été testées. La 1 ère consiste à chauffer le verre à une seule température pendant un temps assez long, de 1 à 9 heures typiquement. La 2 ème méthode, appelée aussi méthode à deux étapes, consiste d'abord à faire nucléer le verre à la température )

Annealing time (h) C'est probablement l'une des raisons pour lesquelles ces matériaux composites présentent une très longue durée de vie des porteurs de charge minoritaire d'environ 16 µs (comme le montre la Figure 14), par rapport à des CIGS, similaires à nos matériaux en constituants et bien étudiés avec une durée de vie généralement plus courts que 0,1 μs. L'inhomogénéité de la durée de vie est associée au caractère composite du matériau avec des canaux conducteurs distribués de façon aléatoire comme précédemment mesurés The objective of this work is to explore a totally new photoelectric material with unique microstructure by controlling the growth of nanocrystals in the GeSe 2 -Ga 2 Se 3 -CuI glass ceramic system. The microstructure of the glass ceramic is mainly determined by the material composition and the crystallizing process of the precursor glass. Consequently emphasis will be put on the establishment of the relationship between the synthesis (composition and process), the microstructure (crystalline phases and composite structure) and the properties (photoelectric and photocatalytic activities) of the glass ceramic. The objective is to understand the mechanism of charge generation, separation and transport in order to design appropriate structures for efficient photon-electron/hole conversion.

The first chapter of this dissertation is on the-state-of-the-art concerning the exploitation of solar energy. The purpose is to situate this work in this context and to explain the motivation.

The GeSe 2 -Sb 2 Se 3 -CuI system has been selected on the base of previous studies devoted to the development of infrared transmitting glass ceramics. It was discovered that these materials have unusual optical properties, suggesting the existence of high concentration of free charge carrier. This kind of materials is of course interesting for photovoltaic or photocatalytic applications.

Chapter 1

Conversion of Solar Energy 1 Introduction

With the development of human civilization, the energy consumption is rising tremendously year by year. In the present scenario, fossil fuels such as coal, oil and gas, are playing lead role to meet the energy demand. However, as the global fossil resources are limited, it is gradually becoming difficult to solve such a contradiction between supply and demand of human being. Furthermore, fossil fuels have been regarded as the largest greenhouse gas emitters in the world, contributing 3/4 of all carbon dioxide, methane and other greenhouse gas emissions. Burning coal, petroleum and other fossil fuels at extremely high temperatures (combustion) is the primary means by which electricity is produced, but it leads also to high concentration of pollutants in our air and water. The combined effects of the depletion of fossil fuels and the gradually emerging concern about environmental degradation have forced humanity to increase the use of renewable energy resources in the 21 st century.

Nowadays, looking for new, non-polluting and renewable energy resources have become a worldwide issue.

Renewable energy sources include solar energy, biomass energy, wind energy, biogas, etc.

These energy sources are directly or indirectly derived from solar radiation. Biomass energy is produced by plants as a result of photosynthesis generated by the sun; wind energy is also indirectly derived from the heat generated by the sun. Among these renewable energy sources, solar energy has received much attention in solving the environmental and energy crisis due to its clean and inexhaustible feature, and thus occupies a central position as renewable energy.

It is predicted that by 2100 solar energy will account for 70% in the energy consumption structure of the world (as shown in Figure 1-1) [1]. It is evident that, despite the attenuation, the total amount of solar energy available on the Earth is still of an enormous amount, but owing to its low-density and intermittency, it needs to be transformed and stored efficiently. The commonly used techniques for transforming solar energy are photothermal conversion, photochemical conversion and photovoltaic conversion. For the photothermal conversion, the solar energy is transformed into heat which may be subsequently used directly, stored or transformed into mechanical, electrical or chemical work. While for the photochemical conversion and photovoltaic conversion, they have a similar mechanism. The absorbers based on quantum processes transform part of the solar energy directly into electrical energy (as it happens in a photovoltaic cell) or chemical energy (as it happens in case of photocatalytic water splitting into oxygen and hydrogen).

2 Solar energy conversion

Photothermal conversion

Photothermal conversion of the solar energy can be divided into different applications according to the working temperature. Low temperature thermal application (less than 80°C) mainly includes heating, water heating and air conditioning, etc. The temperature zone of medium temperature application is typically from 80 °C to 250 °C, and the main applications are solar cooling, solar boiler and solar desalination. High temperature thermal application mainly refers to the use of high temperature (over 250 °C) for electric power generation as an example. Figure 1-3 summarizes the current major solar energy utilization with photothermal conversion [3]. The solar energy has low flux density (about 1353W/m 2 at atmospheric peripheral, and less than 1000W/m 2 at the ground), low energy grade and intermittent changes in circadian.

Therefore, the use of concentrating technology can improve the conversion temperature and get high quality heat, allowing mid-high temperature thermal applications of solar energy. In the mid-high temperature solar energy utilization, the current research directions mainly include the medium temperature solar industrial heating, solar refrigeration and air conditioning and high temperature solar thermal power technology. There are typically four types of concentrating systems: trough, tower, dish and Fresnel reflection (Figure 12345).

Among them, the trough and tower thermal power technology has achieved commercial operation. The dish thermal power system is still in the demonstration phase due to the high cost of the Stirling engine and its high technical complexity. The maximum consumption of the medium temperature solar energy is mainly in industrial sectors to meet the thermal requirements for food processing, plastics processing, glass processing, chemical industry, paper industry, wood processing, textile industry and other industries. Solar thermal power generation is an important direction of solar energy utilization.

Compared with silicon, solar thermal power generation can significantly reduce the cost of solar power and achieve controllable solar thermal power generation through the energy storage materials. The energy conversion process is as following: the solar radiation is first converted into heat, and then the heat energy is converted into mechanical energy and finally the mechanical energy converted into electrical energy. Currently, the large-scale solar thermal power generation is based on central power tower. Tower solar collector system (as shown in Figure 123456) is an array of heliostats on the ground with each heliostat tracking the sun on two axes to continuously reflect the solar radiation onto the tower-top receiver. There are various dual-axis tracking modes, such as azimuth-elevation, radial-pitch-roll, azimuthpitch-roll, oriented polar and receiver [6]. Tower collector concept was first raised by Soviet scientists in the 1950s. In 1965, the Italian scientist Giovanni designed and manufactured the first tower solar collector, and the mirror farm consisted of 120 circular mirrors, which can reflect the solar radiation to the steamer in the tower and generate the superheated steam with a temperature of 500 °C and a pressure of 10MPa [7]. In 1982, the United States built a 10MWe massive solar tower power plant called Solar One (as shown in Figure 1234567). The receiver consists of 24 panels (receiver diameter: 7 m), six of which are for preheating water and eighteen for producing steam. Another solar thermal utilization technology is the use of phase change materials in the residential walls. The phase change materials have the following advantages: Firstly, the phase change process is generally isothermal or approximately isothermal process. This feature is useful to maintain the temperature in a small range, and thus people will feel comfortable. Secondly, the phase change materials can have a high latent heat, so a small amount of material can store large amounts of heat. Using phase change materials to build the walls can significantly enhance the regenerative capacity of the wall, increasing the building's thermal inertia index, thus improving the indoor environment of thermal comfort [12]. Take the heating for example, when the indoor temperature is higher than the phase transition temperature, the phase change materials melt and absorb excess heat in the room; at night when the room temperature is lower than the phase transition point, the phase cha nge materials solidify and slowly release the stored heat into the room to increase the indoor temperature. Compared to the classic building materials (such as concrete, brick, etc.), the phase change materials have higher thermal storage density and lighter weight, and can be therefore large-scale used in lightweight buildings to improve its thermal performance and indoor thermal environment [13,14].

Photochemical conversion

Since 1972, Fujishima and Honda first reported the possibility of electrolyzing water into hydrogen using solar illuminated TiO 2 semiconductor electrode. Hydrogen production from water photolysis, as a new way to utilize renewable solar energy, has attracted more and more attentions [15]. Figure 1-9 shows the main processes of photocatalytic water splitting. (3) Construction of surface reaction sites. Ideally, charge is extracted via the water splitting half-reaction at the solid/liquid interface. The extraction process can be slowed or completely inhibited by poor energetic alignment or poor surface kinetics at the photoelectrode or counter electrode surfaces. Moreover, parasitic or corrosion reactions competing with the watersplitting reactions can result in substantial losses. Surface treatments can be employed to kinetically and/or energetically favor water splitting over the parasitic processes, but such treatments could also block sunlight. Surface incorporation of nanoparticle catalysts is one approach. Since photoelectrochemical water splitting is a low-current density process (typically operating below 20 mA• cm -2 ), non-precious-metal catalysts are necessary. In addition, nanostructure of electrode surfaces can increase effective surface area for enhanced charge extraction, although this can also lead to higher surface recombination loss.

The positions of the conduction band and valence band edges of the semiconductor are important as they determine whether spontaneous water splitting will occur and whether the material will be stable in a photoelectrochemical (PEC) cell. ). These active substances can decompose most organic compounds (formula 1.8). In addition, for the most used photocatalyst such as anatase TiO 2 , since the electric potential of its photo-generated hole is greater than 3.0 eV, which is higher than that of chlorine, ozone and permanganate, thus showing strong oxidation. Therefore, such photo-generated holes provide another way to decompose the organic pollutants (formula 1.9) [19].

e -+ O 2 → •O 2 - (1.1) • O 2 -+ H 2 O → •HO 2 + OH - (1.2) 2•HO 2 → H 2 O 2 + O 2 (1.3) e - + • HO 2 + H 2 O → H 2 O 2 + OH - (1.4) e -+ H 2 O 2 → •OH + OH - (1.5) h + + H 2 O → •OH + H + (1.6) h + + OH - → •OH (1.7) Organic pollutants + (• HO, •O 2 - , •HO 2 ) → Products (1.8)
Organic pollutants + h + → Products (1.9)

Photovoltaic conversion

Solar photovoltaic system usually consists of collector panels, controller, inverter, battery and other devices (as shown in Figure 1-12). The collector panels include a number of small series / parallel unities, depending on the requirements for the output current and voltage; the controller is usually used to control the stability of the output parameters, in order to maximize the use efficiency of obtained solar energy; the inverter is used to transform the DC power into AC power for subsequent use; the battery is typically used to store excess power. Solar cells are the core devices of solar photovoltaic system, and already have a development history of nearly two hundred years. The first event to be noted in the history of solar cells was the discovery of selenium made by Berzelius in 1817 [21]. Incidentally, he was also the first to prepare elemental silicon a few years later. However it was over two decades before Alexandre-Edmond Becquerel, a French physicist, observed the photovoltaic effect in 1839. In 1873, Willoughby Smith reported the photoconductivity of selenium [22]. This has inspired W.G.Adams and R.E. Day to make the first selenium cell in 1877. However, the typical cell efficiency was less than 1%, and therefore unrealistic for large-scale power production. In 1904, Hallwachs found a new material of the combination of copper with cuprous oxide demonstrating photosensitive property. The science of photovoltaics also evolved with Milikan's experimental proof of the photovoltaic effect in 1916 and Albert Einstein's Nobel Prize of 1921 for his contribution to explaining this effect [23]. In 1954, Bell labs researchers D.M. Chapin, C.S. Fuller and G.L. Pearson, reported the monocrystalline silicon solar cells with efficiency of 4.5%, which reached 6% a few months later [24]. Such cells were applied to America's second satellite, Vanguard 1 in 1958 [ 25 ]. The most significant market growth for PV has come since 1995, and the world's solar cell production capacity has reached 37.2 GW in 2011 [26].

Solar cells convert solar energy into electrical energy based on the "photovoltaic effect". As shown in Figure 1 [27].The best efficiency of CIGS thin film solar cells reached 20.3% [28]. The range of solar cells continues to expand, with two relatively new entries-dyesensitized solar cells (DSC) and organic photovoltaic cells (OPV)-making rapid progress in laboratory efficiencies. The DSCs and organic solar cells are late starters, but both of these low-cost technologies have reached efficiencies that make them commercially interesting if scale-up and acceptable lifetimes can be demonstrated. Among all of these solar cells, multijunction cells have achieved much higher efficiencies.

The two main reasons are: (1) the highest theoretical efficiencies may be achieved if multiple semiconductor materials (with a range of bandgaps) are chosen to match the spectral distribution of the Sun; (2) the compound semiconductors used in these cells are mostly direct-gap materials and can be grown with near-perfect quality in mass production. The multijunction approach has been described extensively in the literature [29,30]. In pure silicon crystal, the four valence electrons of a silicon atom are tied with strong covalent bonds to four adjacent silicon atoms. Therefore at zero temperature, silicon is a perfect insulator. Doping is used to make silicon more conductive. By doping silicon with a small number of five-valent atoms (phosphorous, arsenic, or antimony), silicon gets an extra electron to share. This is called an n-type semiconductor. In contrast, doping silicon with an atom of three electrons in its valence, the silicon essentially "loses" an electron to share and this creates a "hole". This is called a p-type semiconductor. These two types of silicon are put into contact with each other and a p-n junction is formed at the boundary between them. Once Several issues affect the efficiency of silicon solar cells [34] (as shown in Figure 1-17): (1)

The energy of the photons hitting the solar cell is less than the band gap, so the light cannot be converted into electricity and is lost. Crystalline silicon solar cells, including monocrystalline and polycrystalline cells, are the most widely used of all solar cells due to their relatively high efficiency, high stability and high service life. Figure 1-18 shows the typical production process for crystalline silicon solar cells [35]. The cost of fabricating monocrystalline silicon solar cells is essentially due to the purification process and single crystal growth. To purify silicon, quartzite gravel, high purity silica, is melted and reduced using a carbon bed at a temperature above 1800 °C; this process makes the silicon 98-99% pure and is called metallurgical grade silicon (MGS). MGS is then ground and reacted with hydrochloric acid at 300 °C to make trichlorosilane (TCS). The TCS is heated to 1100 °C in a hydrogen atmosphere to make electronic grade silicon (EGS); this final process makes the silicon 99.9999999% pure. Then a process known as the Czochralski process is used for single crystal silicon production. In this process, a small single crystalline silicon seed is inserted into molten EGS at a temperature above 1700 °C. The seed is continually rotated and slowly drawn out of the melted silicon, which allows the silicon atoms to attach to the seed and arrange in a single crystal lattice; this forms a uniform single crystalline ingot of silicon. Silicon wafers are sliced from the ingot, with a thickness of about 0.2mm. And then these wafers are chemically and mechanically polished. In the production process of solar cells, the first step is to produce a textured structure and then to dope silicon through diffusion, thus forming a p-n junction on the surface of silicon. This step is a key step to produce crystalline silicon solar cells. Next step is the deposition of anti-reflective film to further improve light absorption of cell and finally the screen printing method is used to form silver or aluminum contact on the cell surface. And then the cell is sintered, leading to the formation of alloys between the metal electrodes and silicon to achieve ohmic contact. Compared to monocrystalline silicon solar cells, polycrystalline cells are much less expensive to manufacture but also less efficient. Polycrystalline silicon is made by pouring molten silicon in a mold and allowing it to cool. This process is called casting. The resulting silicon has no overall lattice structure, but the ingot thus produced has large column grains of crystallinity. The ingot is still sliced and treated as the single crystalline silicon ingot except that the bottom and top of the ingot are removed before slicing due to poor crystallinity at the edges. The casting process is much cheaper and simpler than the Czochralski method used for single crystal growth.

The efficiency of crystalline silicon solar cells is related to the solar cell structure [36].

Representative examples of high-efficiency monocrystalline silicon solar cells are the passivated emitter rear localized (PERL) cell, the heterojunction with intrinsic thin layer (HIT) cell, the back contact-back junction (BC-BJ) cell and metal wrap through (MWT) cell (as shown in Figure 12345678910111213141516171819). The PERL cell is a research PV cell with front and rear surface passivation layers, an inverted-pyramid light-trapping surface, a rear localized p+ layer (BSF), a double-layer anti-reflection coating (ARC) and a p-type float zone (FZ) monocrystalline silicon substrate. A PERL cell efficiency of 25.0% was reported by researchers from the University of New South Wales (UNSW) in 2009 [ 36 ]. The HIT cell has a unique heterojunction structure consisting of very thin, amorphous p-and n-doped layers and intrinsic amorphous layers on the front and rear surfaces of a Czochralski-grown (CZ) n-type monocrystalline-silicon substrate. The best efficiency of this cell is 23.0% [37]. The BC-BJ cell has front and rear surface passivation layers, a random-pyramid light-trapping surface, front surface field (FSF), interdigitated n-and p-doped regions on the back surface, n and p contact gridlines on n-and p-doped regions, a single-layer ARC and a CZ n-type single crystalline silicon substrate. The best conversion efficiency reported so far for a large-area Chapter 1

-20 -industrial BC-BJ cell is 23.4% [38]. The MWT cell requires only a relatively small number of through-holes to direct photo-generated electrons to the back surface through the metal electrodes and n-doped emitters, and produce higher collection photocurrents due to the absence of a bus bar (main electrode) on the front surface as in conventional cells. The best efficiency of this cell has reached 18.3% [39]. silicon atoms occurs more frequently in amorphous silicon than in crystal silicon, allowing much more light to be absorbed. Thus, an ultra-thin amorphous silicon film of less than 1µm can be produced and used for power generation. Also, by utilizing metal or plastics as substrate, flexible solar cells can be produced. Amorphous silicon layer uses silane (SiH 4 ) as its source gas and is fabricated using a CVD method [40]. The fundamental photodiode inside an amorphous silicon solar cell has three layers deposited in either the p-i-n or the n-i-p sequence. The three layers are a very thin (typically 20 nm) p-type layer, a much thicker (typically a few hundred nanometer) undoped intrinsic layer, and a very thin n-type layer, as

shown in Figure 1-20. This p-i-n junction corresponds to the p-n junction of a crystalline silicon solar cell. In the process of this junction formation, a number of cells are connected in series on a substrate at one time. The sunlight can enter through the transparent substrate, which is usually glass or a transparent plastic. The insulating substrate needs a conducting layer, which is typically a "transparent conducting oxide" (TCO) such as SnO 2 . The amorphous silicon photodiode layers (p-i-n junction) are then deposited onto the TCO, starting with a p-type layer. Finally, a back contact (containing reflector) is deposited onto the photodiode. The reflector acts as an electrode to the n-type photodiode layer. However, the poor crystallinity and the existence of dangling bonds in amorphous silicon will lead to short carrier diffusion length and instability upon prolonged illumination [42].

Moreover, the optical threshold of amorphous silicon is rather high (1.7-1.8 eV) and the material suffers from light induced degradation [ 43]. Therefore, short life cycle and low efficiency have become the big problems for amorphous silicon solar cells. The conversion efficiency of amorphous silicon solar cells is much lower than that of the crystalline silicon solar cells. The best conversion efficiency of this cell achieves 12.5% in laboratory-scale, while industrial conversion efficiency is not more than 10% [44].

Polymer solar cells

As shown in Figure 1-21, the photovoltaic energy conversion process of polymer solar cells can be divided into four steps: light absorption, exciton diffusion, charge transfer, and charge collection [ 45,[START_REF] Thompson | Polymer-fullerene composite solar cells[END_REF]. The horizontal lines on the right and left of each illustration correspond to the Fermi energies (E F ) of the cathode and anode contacts, respectively. The boxes correspond to the acceptor (left) and donor (right) energy gaps, respectively. Here, LUMO is the lowest unoccupied molecular orbital, and HOMO is the highest occupied molecular orbital of the organic film. Red dots are electrons and red circles are holes, with dashed lines drawn between them to represent excitons. Also, the dip in the energy levels in the vicinity of the exciton qualitatively depicts its binding energy (0.5-1 eV), placing this quasi-particle at an energy somewhat below that of the HOMO-LUMO gap energy. η A is the absorption efficiency of light within the active region of the solar cell, η ED is the exciton diffusion efficiency to a dissociation site, η CC is the charge collection efficiency, and η CT is the charge transfer efficiency. (1) Light absorption (exciton generation): The active polymer layer absorbs incident photons and then the electrons from the polymer highest occupied molecular orbital (HOMO) are excited to the lowest unoccupied molecular orbital (LUMO), forming the excitons. This process occurs in 10 -13 -10 -9 seconds. The active polymer layer is a major factor that determines the efficiency of polymer solar cells. The currently best active layer material can absorb up to 40% of the sunlight, thus wasting the majority of the solar energy. Therefore, in order to improve the energy conversion efficiency of the final device, the first thing need to be done is to make the active polymer layer more absorbents. For an efficient collection of photons, the absorption spectrum of the organic active layer should match the solar emission spectrum and the layer should be sufficiently thick to absorb the entire incident light. A better overlap with the solar emission spectrum is obtained by lowering the band gap of the organic material, but this will ultimately have some negative effects on the open-circuit voltage.

Increasing the layer thickness is positive for light absorption, but will slowdown the charge transport and also increase the internal resistance of the device [START_REF] Heeger | Semiconducting polymers: the third generation[END_REF][START_REF] Pivrikas | A review of charge transport and recombination in polymer/fullerene organic solar cells[END_REF].

Chapter 1 -23 -

(2) Exciton diffusion: Creation of charges is one of the key steps in the conversion of sunlight into electrical energy. However, the exciton is separated into charges (i.e. electrons and holes), only at the interface of electron donor/acceptor material. Therefore, when the exciton is generated far from the electron donor/acceptor interface, the exciton must diffuse to the electron donor/acceptor interface before further charge separation. But exciton diffusion is in competition with other decay processes such as luminescence or non-radiative recombination to the ground state. For an efficient solar cell, all excitons have to reach the electron donor/acceptor interface within exponential lifetime, which determines the distance an exciton is able to cross. This distance is generally limited to 10 nm. Therefore, the morphology of electron donor/acceptor is essential [START_REF] Thompson | Polymer-fullerene composite solar cells[END_REF].

(3) Charge transfer (exciton dissociation): In this step, an electron is transferred from the electron donor to the electron acceptor driven by the energy difference (offset) between the LUMOs of the donor and acceptor. It appears that a minimum energy difference of 0.3 eV is required to induce the exciton splitting and charge dissociation. In the photo-induced electron transfer process, an exciton at the donor/acceptor interface decays by creation of the chargeseparated state consisting of the radical cation of the donor and the radical anion of the acceptor. For an efficient charge generation, it is essential that the charge-separated state is the thermodynamically and kinetically most favourite pathway for the exciton. Therefore, it is important that the energy of the absorbed photon is used for generation of the charge separated state and is not lost via competitive processes like fluorescence or non-radiative decay. In addition, it is of importance that the charge-separated state is stabilized, so that the photo-generated charges can migrate to one of the electrodes [START_REF] Brabec | A low-bandgap semiconducting polymer for photovoltaic devices and infrared emitting diodes[END_REF][START_REF] Brabec | Plastic solar cells[END_REF].

(4) Charge collection: In this step, charges transfer to the respective electrodes, and the collection of charges at the electrodes is regularly accomplished by a transparent conductive oxide (TCO) such as ITO on one side and a metal contact on the other side. The effectiveness of the charge collection is mainly affected by the energy level matching between electron donor material and electron acceptor material, as well as the contact situation [START_REF] Rand | Organic doubleheterostructure photovoltaic cells employing thick tris(acetylacetonato)ruthenium(III) exciton-blocking layers[END_REF][START_REF] Wu | Stable small-molecule organic solar cells with 1,3,5-tris(2-N-phenylbenzimidazolyl) benzene as an organic buffer[END_REF].

In 1986, Tang firstly introduced a double-layer structure of a p-and n-type organic semiconductor with power conversion efficiency of about 1%. Since then, polymer solar cells have undergone a gradual evolution that has led to energy conversion efficiencies of about 10% [31]. To date, two main approaches have been explored in the effort to develop viable devices:

the donor-acceptor bilayer [START_REF] Xue | 4.2% efficient organic photovoltaic cells with low series resistances[END_REF], commonly achieved by vacuum deposition of molecular components, and the so-called bulk heterojunction (BHJ) [START_REF] Halls | Efficient photodiodes from interpenetrating polymer networks[END_REF], which is represented in the ideal case as a bicontinuous composite of donor and acceptor phases, thereby maximizing the all-important interfacial area between the donors and acceptors (as shown in Figure 1 particles link each other and form a network structure by sintering, which not only provides a huge surface area, but also has good inter-particle electron conductivity. Usually, the porous TiO 2 film is supported on the transparent conductive substrate such as FTO glass substrate, and sensitized with a molecular dye chemisorbed on its surface. An electrolytic solution containing iodide and triiodide ions is used as a redox mediator to reduce the oxidized dye molecules. A platinum-coated glass substrate placed parallel to the working electrode with a face-to-face configuration acts as counter electrode. Compared with existing solar cells based on silicone technology, DSSC has big advantages in manufacturing cost and productivity of large devices because high vacuum and high temperature are not necessary for production. In addition, the DSSC will also work in low light conditions (i.e. cloudy skies and indirect sunlight), and it has been suggested that with the use of DSSC, indoors-light from the various lighting sources could even be absorbed [START_REF] O'regan | A low-cost, high-efficiency solar cell based on dye-sensitized colloidal TiO 2 films[END_REF].

The major disadvantage is that the liquid electrolyte used in DSSC is temperature-sensitive.

The electrolyte can freeze at low temperatures cutting power production and causing physical damage. Sealing the panels becomes a difficult task, when the liquid expands at higher temperatures. The use of a liquid electrolyte causes some serious additional problems such as potential instability, limitation of maximum operation temperature, danger of evaporation and extra cost for forming an electrical series connection 57 . Moreover, compared with the first and second generation solar cells based on conventional semiconductor materials with conversion efficiencies of ∼20 -30%, current dye-sensitized solar cells still demonstrate efficiencies relatively low. To date, the world's highest level for dye-sensitized solar cells is about 11.4% [31].

Compound solar cells

Compound solar cells include gallium arsenide (GaAs) thin film solar cell, cadmium telluride (CdTe) thin film solar cells, copper indium gallium selenide (CIGS) thin film solar cells.

GaAs is a direct bandgap semiconductor with a bandgap of 1.45 eV, thus is regarded as an ideal photovoltaic semiconductor materials. The first GaAs solar cells having efficiencies in excess of 6% were already reported in 1956 [START_REF] Seitz | Displacement of atoms duringi[END_REF]. Since then, researchers began to study homojunction GaAs solar cells. However, the cost/efficiency ratio of this cell was high compared with silicon-based solar cells, thereby inhibiting the development of GaAs solar cells. In the early 1970s, multi-junction structure attracted attentions of scientists. Such structured GaAs solar cells consist of several stacked p-n junctions (as shown in Figure 1-24).

Chapter 1 -27 -Each junction is tuned to a separate range of light wavelengths, so that the multiple p-n junctions are able to absorb a wider range of the solar spectrum (as shown in Figure 1-25), allowing for increased conversion efficiency over the solar spectrum. Progress in the development of multi-junction efficiency has shown promising results [START_REF] Yastrebova | High-efficiency multi-junction solar cells: Current status and future potential[END_REF]. From the 16% in the 1980s to the current record of 43.5%, it is expected that within the coming years, multijunction solar cell for commercial use would reach efficiencies of 50% [32]. However, this kinds of solar cells are complex and expensive to fabricate, requiring additional tunnel layers between p-n junctions. The record efficiency for a laboratory CdTe solar cell is 18.7% to date [START_REF] Green | Solar cell efficiency tables (version 41)[END_REF], which is well above current commercial CdTe modules. However, Tellurium is a rare metalloid, while

Cadmium is toxic if breathed in fine dust or fumes. Therefore the application of CdTe solar cell is limited. The structure of CIGS solar cells is shown in Figure 1-27. Soda lime glass is commonly used as a substrate. Furthermore, many companies are also looking at lighter and more flexible substrates such as polyimide or metal foils [START_REF] Dhere | Toward GW/year of CIGS production within the next decade[END_REF]. A molybdenum layer, which serves as the back contact, is deposited on the glass substrate to reflect most unabsorbed light back into the absorber. Then a p-type CIGS absorber layer is grown on the molybdenum layer, followed by the deposition of a thin n-type buffer layer such as CdS. Next, the buffer is overlaid with a thin, intrinsic ZnO layer (i-ZnO) to protect the CdS and the absorber layer from sputtering damage while depositing the ZnO:Al layer. Finally, a thicker Al doped ZnO layer, serving as a transparent window layer, is deposited on the top of the ZnO layer to collect and move electrons out of the cell while absorbing as little light as possible.

CIGS solar cells have many advantages as follows: (1) Compared to silicon solar cells, CIGS layer can be deposited in a polycrystalline form directly onto molybdenum coated glass sheets, steel bands, especially flexible substrates. Moreover, the CIGS solar cells show a better resistance to heat than silicon solar cells. (2) Compared to CdTe solar cells, CIGS solar cells have much lower level of cadmium, in the form of cadmium sulfide. In some designs, zinc is used instead of cadmium. (3) CIGS has an exceptionally high absorption coefficient of more than 10 5 /cm for 1.5 eV and higher energy photons [START_REF] Stanbery | Copper indium selenides and related materials for photovoltaic devices[END_REF]. CIGS solar cells with record conversion efficiency of 20.3% to date have been reported by Zentrum für Sonnenenergie und Wasserstoff Forschung (ZSW) [31], which is also a record efficiency of thin film solar cells.

However, the efficiency of CIGS solar cells is still lower than crystalline silicon solar cells, for which record efficiency lies at 25% [31]. In addition, the cost of fabricating the CIGS solar cells makes it difficult to be competitive with current grid prices. It may take several more years to solve the manufacturing problems and to bring the production costs in line with the cost of other solar cells.

Figure 1-27 Schematic of CIGS solar cells

In summary, Silicon is the oldest technology and up to now it is also the most diffused in solar cell market (exceeding 80% in 2011 [67]). However, the record efficiency of silicon solar cell is 25.0%, approaching the theoretical efficiency limit of 30% [START_REF] Vasekar | Photovoltaic characterization of Copper-Indium-Gallium Sulfide (CIGS 2 ) solar cells for lower absorber thicknesses[END_REF]. Its main drawback is the low scalability, which results in modules with low efficiency or very high cost. The single junction record efficiency of 29.1% is held by GaAs, however it is mainly used for space application due to the high production cost. Although multi-junction cells achieve the best conversion efficiency of 43.5%, its structures are extremely complex and expensive. The DSC and OPV are late starters, their efficiencies need further development.

While CIGS is one of the most promising thin film technologies due to its high efficiency and relatively low production cost. CIGS holds the thin film efficiency record and it is also suitable for flexible modules production. CIGS solar cells on flexible substrates offer several advantages for their manufacturing as well as applications compared to solar cells on rigid glass substrates. They are lightweight, can be applied on uneven surfaces and can be rolled-up when not in use. Flexible and lightweight CIGS modules open up new application possibilities for terrestrial and space applications. Recently, the world efficiency record has reached 20.3%, thus demonstrating the potential for futur development [31].

The photoelectric materials studied in this thesis have some similitude with the CIGS compounds. This is the reason why special attention has been paid to the state-of-the-art of the CIGS solar cells.

Development of CIGS solar cells

History of CIGS solar cells

In 1953, CuInSe 2 material was synthesized for the first time by Hahn and co-workers [START_REF] Hahn | Untersuchungen über ternäre Chalkogenide. V. Über einige ternäre Chalkogenide mit Chalkopyritstruktur[END_REF].

In 1974, the first CuInSe 2 solar cells with an efficiency of 12% have fabricated by evaporating n-type CdS onto p-type single crystals of CuInSe 2 by Wagner and co-workers [START_REF] Wagner | CuInSe 2 /CdS Heterojunction Photovoltaic Detectors[END_REF].

In 1976, the first thin-film CuInSe 2 /CdS devices were fabricated by Kazmerski and coworkers, using films deposited by evaporating CuInSe 2 powder in excess of Se vapor [START_REF] Kazmerski | Thin-film CuInSe 2 /CdS heterojunction solar cells[END_REF].

In 1981, the first high-efficiency all thin film solar cell based on the system n-ZnCdS/p-CuInSe 2 was prepared by Mickelsen and co-workers from Boeing, realizing a conversion efficiency of about 9.4% and in 1985 they reached the efficiency of 11.4% [START_REF] Mickelsen | Polycrystalline thin-film CuInSe 2 solar cells[END_REF][START_REF] Devaney | Paper presented at 18th IEEE Photovoltaic Specialists Conference[END_REF].

In 1988, by using a different process based on selenization of stacked metal layers, Arco Solar achieved cell efficiencies as high as 14.1% [START_REF] Mitchell | Single and tandem junction CuInSe 2 cell and module technology[END_REF].

Since the late 1980s, in order to make full use of the solar spectrum, the gallium element was introduced into the CuInSe 2 absorber layer (has a bandgap of 1.04 eV) to increase the band gap (an indium to gallium ratio of 20-40% will result in a bandgap of approximately 1.1-1.2 eV) and to get a better match with the solar spectrum, leading to an important increase in photoelectric conversion efficiency [START_REF] Potter | Enhanced photocurrent ZnO/CdS/CuInSe 2 solar cells[END_REF].

In 1994, the so-called three-stage process, introduced by NREL [START_REF] Gabor | Highefficiency CuIn x Ga 1-x Se 2 solar cells made from (In x ,Ga 1-x ) 2 Se 3 precursor films[END_REF], is obtained by starting the deposition with an (In,Ga) x Se y precursor, followed by the co-deposition of Cu and Se until Cu-rich overall composition is reached, and finally the overall Cu concentration is readjusted by subsequent deposition of In, Ga and Se. The obtained conversion efficiency reached 16.4%. This method leads, up to now, to the most efficient solar cells.

From then on, it was possible to fabricate thin film solar cells based on the CdS/CuInGaSe 2 system with an efficiency of 18.8% in 1999 [START_REF] Contreras | Progress toward 20% efficiency in Cu(In,Ga)Se 2 polycrystalline thin-film solar cells[END_REF], of 19.2% in 2003 [START_REF] Ramanathan | Properties of 19.2% efficiency ZnO/CdS/CuInGaSe 2 thin-film solar cells[END_REF] and of 20.3% in 2011 [31]. This last result is also the highest value for energy conversion efficiency in an all thin film photovoltaic device.

Research progress on CIGS absorber layer

Nowadays, the aim of CIGS solar cells is to achieve high conversion efficiency and low production cost. As the core part of CIGS solar cells, the CIGS absorber layer still attracts attention from the scientific community and the most significant progresses are achieved in the two following domains:

(1) Material research on CIGS absorber layer Cu(In,Ga)(Se,S) 2 alloys from the I-III-VI group of the periodic table are used as absorber layers, including a wide range of band gaps. CuInSe 2 has a relatively narrow band gap of 1.04 eV which can be better adjusted to the solar spectrum by replacing part of In with Ga and part of Se with S. The band gap can be adjusted from 1.04 eV (CuInSe 2 ) to 1.53 eV (CuInS 2 )

and to 1.7 eV (CuGaSe 2 ) [START_REF] Wei | Effects of Ga addition to CuInSe 2 on its electronic, structural, and defect properties[END_REF][START_REF] Yakushev | Energy of excitons in CuInS 2 single crystals[END_REF]. Hence one can change the bandgap throughout the absorber material in the solar cell and make a graded CIGS solar cell. This kinds of cell have two main advantages: i) Increasing the bandgap towards the back of the cell (back grading) can create an electrical field which pulls the electrons towards the space charge region so that they can contribute to the light current. ii) Increasing the bandgap towards the front of the absorber layer (front grading) will reduce the recombination losses in the space charge region and thus increase the cell efficiency [START_REF] Topic | Band-gap engineering in CdS/Cu(In,Ga)Se 2 solar cells[END_REF]. A graded CIGS solar cell has achieved an efficiency of 19.9% [START_REF] Repins | 19.9%-efficient ZnO/CdS/CuInGaSe 2 solar cell with 81.2% fill factor[END_REF]. Most recently, the multilayer CIGS solar cells with Cu gradients in the absorber layer have been studied. The results show that with the Cu and Ga distribution stepped simultaneously in the multilayer cells, the improvement of efficiency could be approximately 3.5% due to the enhanced charge carrier collection for longer-wavelength photons, and to the reduced recombination at the heterojunction and back regions of the cell [START_REF] Sharbati | Model for increased efficiency of CIGS solar cells by a stepped distribution of carrier density and Ga in the absorber layer[END_REF].

The use of rare materials such as In and Ga increases the production cost of the entire device and it is therefore necessary to find new alternative materials for In and Ga elements.

Zn and Sn were mostly used to substitute In and Ga, forming Cu 2 ZnSn(Se, S) 4 (CZTS) absorber layer, while the structure of the solar cells was very similar to the CIGS solar cells standard structure having CdS as buffer layer, ZnO as front and Mo as back contact. The band gap of Cu 2 ZnSnS 4 is around 1.5eV, while for Cu 2 ZnSnSe 4 the gap stays around 1 eV [START_REF] Redinger | The Consequences of Kesterite Equilibria for Efficient Solar Cells[END_REF][START_REF] Hibberd | Non-vacuum methods for formation of Cu(In, Ga)(Se, S) 2 thin film photovoltaic absorbers[END_REF][START_REF] Weber | On the Sn loss from thin films of the material system Cu-Zn-Sn-S in high vacuum[END_REF]. The band gap of CZTS is affected by the deposition methods, crystallization procedure as well as presence of secondary phases, such as Zn(Se, S), Cu(Se, S), CuSn(Se, S). The best solar cells show a Zn-rich and Cu-poor composition, which provides a Zn(Se, S) secondary phase. As Zn(Se, S) has a wide band gap and low conductivity, it would not affect the opencircuit voltage, and could lead to the high series resistance observed in all solar cells [START_REF] Katagiri | Development of thin film solar cell based on Cu 2 ZnSnS 4 thin films[END_REF]. The efficiency of CZTS solar cells was obviously increased to 5.45% in 2003 [START_REF] Katagiri | Solar cell without environmental pollution by using CZTS thin film[END_REF] and then to 6.7%

in 2008 [START_REF] Katagiri | Enhanced conversion efficiencies of Cu 2 ZnSnS 4 -based thin film solar cells by using preferential etching technique[END_REF] by optimization of the sulfurization process. In 2010, the efficiency of 9.6% was reached by inserting a selenization step [START_REF] Todorov | High-efficiency solar cell with earth-abundant liquid-processed absorber[END_REF]. In August 2012, IBM announced they have developed a world-record setting CZTS solar cell with conversion efficiency of 11.1% [START_REF] Todorov | Shedding light on new frontiers of solar cell semiconductors[END_REF].

However, the efficiency of CZTS solar cell is still lower than that of CIGS solar cell.

(2) Production process research on CIGS absorber layer

The absorber deposition methods mainly include vacuum-based and non-vacuum based methods [START_REF] Reinhard | Technological status of Cu(In,Ga)(Se,S) 2 -based photovoltaics[END_REF][START_REF] Jackson | New world record efficiency for Cu(In,Ga)Se 2 thin-film solar cells beyond 20%[END_REF]. The most widely used and successful vacuum-based method is the 'coevaporation', in which all the four elements of Cu, In, Ga and Se are simultaneously delivered to the Mo coated substrate, forming the absorber layer in a single step. The Zentrum für Sonnenenergie und Wasserstoff forschung (ZSW, Germany) and the National Renewable Energy Laboratory (NREL,USA) used a high-temperature ( 550 o C) three-stage coevaporation process to fabricate CIGS absorber layer on glass substrates and achieved high efficiencies of about 20% [START_REF] Jackson | New world record efficiency for Cu(In,Ga)Se 2 thin-film solar cells beyond 20%[END_REF][START_REF] Repins | 19.9%-efficient ZnO/CdS/CuInGaSe 2 solar cell with 81.2% fill factor[END_REF], while the Swiss Federal Laboratories for Materials Science and Technology (Empa, Switzerland) obtained CIGS solar cell with 20.4% efficiency through a low-temperature (< 450 o C) multi-stage co-evaporation process to prepare the absorber layer using flexible substrates [START_REF] Green | Solar cell efficiency tables (version42)[END_REF]. An alternative vacuum-based method is the sputtering of metal precursors and subsequent selenization followed by sulfurization. Solar Frontier achieved 19.7% efficiency by using such a method [START_REF] Nakamura | Achievement of 19.7% efficiency with a small-sized Cu(InGa)(SeS) 2 solar cell prepared by sulfurization after selenization process with Zn-based buffer[END_REF]. To date, the record efficiency CIGS solar cells are all processed using absorber layers grown via vacuum-based methods.

However, the high vacuum process used in this method increases the fabrication costs. The co-evaporation generally requires a high degree of control and large area coatings with uniform properties are difficult to achieve. Moreover, some of the processes use toxic gases, such as selenium vapor or an H 2 Se atmosphere, in order to convert the deposited film of the Cu-In-Ga alloy or the oxide thereof into the final solar responsive form [START_REF] Ahn | Nucleation and growth of Cu(In,Ga)Se 2 nanoparticles in low temperature colloidal process[END_REF]. Hence, the cheap and convenient non-vacuum based methods have attracted wide attentions of scientists [START_REF] Yamaguchi | Preparation of Cu(In,Ga)(S,Se) 2 thin films by sequential evaporation and annealing in sulfur atmosphere[END_REF].

The various non-vacuum based approaches to fabricate CIGS absorber layer may be grouped into two main categories: preparing from well-dispersed solutions containing CIGS nanoparticles/precursors and preparing through electrodeposition.

The well-dispersed solutions prepared from CIGS nano particles are usually called CIGS inks. The CIGS nano particles (or quantum dots) are commonly synthesized by using hotinjection or solvo-thermal routes [START_REF] Stolle | Nanocrystal photovoltaics: a review of recent progress[END_REF][START_REF] De Mello Donegá | Physicochemical evaluation of the hot-injection method, a Synthesis Route for Monodisperse Nanocrystals[END_REF], which are generally rinsed and modified to improve the dispersibility in certain nonpolar solvents, thus getting stable CIGS inks. Next, CIGS inks are coated on the substrate using liquid deposition techniques (such as dip-coating, spincoating etc.) under ambient condition and then solidified and crystallized at relatively low temperature, which is suitable for coating on flexible substrates [START_REF] Kessler | Technological aspects of flexible CIGS solar cells and modules[END_REF]. For ink-based methods, highest efficiency of 17.1% was achieved by Nanosolar in 2012 [START_REF] Brown | Device characteristics of a 17.1% efficient solar cell deposited by a non-vacuum printing method on flexible foil[END_REF]. In addition, the CIGS absorber layer can also be prepared by direct coating of a homogeneous solution containing a dissolved metal precursor (Cu, In, Ga organics, salts, chalcogenides etc.) in an organic solvent.

The solvent is a key factor that influences the stability of the solution. Many factors are considered in selecting a suitable solvent, including the boiling point, vapor pressure, polarity, reactivity, surface tension, toxicity and cost [102,[START_REF] Ahn | CuInSe 2 thin-film solar cells with 7.72% efficiency prepared via direct coating of a metal salts/alcohol-based precursor solution[END_REF]. In 2013, IBM achieved a conversion efficiency of 15.2% by this solution-based method using hydrazine solvent [START_REF] Todorov | Solution-processed Cu(In,Ga) (S,Se) 2 absorber yielding a 15.2% efficient solar cell[END_REF].

Electrodeposition is another solution-based method to fabricate CIGS absorber layer, which have advantages of low instrumental and materials cost, high deposition speed, efficient utilization of raw materials and feasibility in making large area films. This deposition technique can be divided into one-step method and two-step method [START_REF] Bhattacharya | CIGS-based solar cells prepared from electrodeposited stacked Cu/In/Ga layers[END_REF][START_REF] Bhattacharya | Electrodeposited two-layer Cu-In-Ga-Se/In-Se thin films[END_REF]. In a one-step deposition, the concentration and pH value of the electrolyte must be adjusted, so the electrode potentials of all the individual elements may come closer to each other. Therefore, one-step electrodeposition of CIGS layers is rather difficult due to the large difference in the values of equilibrium reduction potential for each constituent. Many works have been carried

out to find out the best concentration of electrolyte for obtaining the stoichiometry, as well as the appropriate deposition potential in one-step electrodeposition of CIGS layer [START_REF] Quang | Effect of electrodeposition potential on the composition and morphology of CIGS absorber thin film[END_REF]. While the two-step method can be employed via stacked layer structure deposition of pure elements, binary or ternary compounds, followed by a selenization treatment [START_REF] Friedfeld | Electrodeposition of CuIn x Ga 1-x Se 2 thin films[END_REF][START_REF] Ganchev | Preparation of Cu(In,Ga)Se 2 layers by selenization of electrodeposited Cu-In-Ga precursors[END_REF]. As the asdeposited CIGS layers are poorly crystalline and may contain Cu-Se second phases, the high temperature post-annealing is carried out in a Se-rich environment to avoid micro cracks and defects in the layer. The layer morphology and the resulting conversion efficiency depend on the relative ratio of the component elements in the fabricated layer [START_REF] Hibberd | Non-vacuum methods for formation of Cu(In, Ga)(Se, S) 2 thin film photovoltaic absorbers[END_REF].

4 Photocatalytic materials

State-of-the-art and research directions

From the photocatalytic mechanism analyzed in section 2.2, it is obvious that the efficient generation, separation and migration of photo-generated charge carriers (electrons and holes) is necessary for the semiconductor photocatalytic reaction. Based on this understanding, the current semiconductor photocatalysis research is mainly carried out to address the two following issues:

(1) How to efficiently capture the photo-generated charge carriers? The high specific surface area of photocatalytic material can enhance the adsorption of O 2 and H 2 O, thereby increasing the yield of active groups to obtain a higher photocatalytic activity. Furthermore, the high specific surface area of the photocatalyst is also beneficial for increasing the contact area with the organic contaminants, thus improving the photocatalytic oxidation reaction rate.

Therefore, photocatalytic materials with high specific surface area have become an important research direction.

Chapter 1 -35 -

(2) How to efficiently produce the photo-generated charge carriers and to promote their efficient separation and migration? For example, the most commonly used TiO 2 photocatalyst, has a band gap of about 3.2 eV (for anatase phase), thus deciding that only the UV light with a wavelength less than 387.5 nm, representing only a small part (3-5%) of solar irradiation, can excite electrons in the valence band . Therefore, the enhancement of visible photocatalytic activity has become a hot research topic. In addition, studies have shown that for a pure photocatalyst, the recombination process of photo-generated charge carriers is faster than its separation process, and only the electrons and holes yielded from effective separation of photo-generated carriers will further react and form active substances, leading to the decomposition of organic pollutants. Therefore, inhibiting the recombination of photogenerated electrons and holes becomes an effective way to improve the photocatalytic activity.

Research on porous photocatalytic materials

Currently, in order to increase the specific surface area of the photocatalytic material, the photocatalytic materials are mainly prepared in the forms of nano powder, porous film and a porous bulk material.

(1) Photocatalytic nano powder

Photocatalytic nano powder has a large specific surface area and its preparation method can be divided into gas and liquid methods. The most typical example of gas preparation is the commercial P25 type TiO 2 nano powder produced by Degussa Company through oxidation of TiCl 4 gas [START_REF] Raillard | Photocatalytic oxidation of volatile organic compounds present in airborne environment adjacent to sewage treatment plants[END_REF]. The most studied method is still a liquid phase method. In this method, the particle size of photocatalyst is mainly determined by the two following parameters: one is an aggregation of nano powder in common drying process; the other one is the growth of nanoparticles during the heat treatment.

By means of freeze-drying or supercritical fluid extraction, one can effectively eliminate the water capillary force to overcome the gel collapse during the drying process, thus preventing the aggregation. Stallings et al. [START_REF] Stallings | Synthesis of nanostructured titania powders via hydrolysis of titanium isopropoxide in supercritical carbon dioxide[END_REF] added water and surfactant into the supercritical CO 2 , forming a micro emulsion. Therefore the reaction can be controlled in the droplets to avoid further crystal growth; furthermore, the nanoparticle aggregation in the drying process can be avoided by using supercritical CO 2 . The as-prepared TiO 2 microspheres have a specific surface area of 275~475 m 2 /g.

Hydrothermal method has become another effective route to prepare nanoparticles because it allows avoiding subsequent high temperature crystallization treatment. Cheng et al. [START_REF] Cheng | Hydrothermal preparation of uniform nanosize rutile and anatase particles[END_REF] diluted TiCl 4 into ice-water to obtain a gel and prepared TiO 2 nano powders with an average particle size of 8~20 nm under hydrothermal conditions. The organic solvent can effectively reduce the interaction between particles due to its weak polar and small surface tension. Thus, Wang et al. [START_REF] Wang | The synthesis of nanocrystalline anatase and rutile titania in mixed organic media[END_REF] used the organic solvents to replace the traditionally used water solvent in the hydrothermal reaction and prepared TiO 2 nano powders with an average particle diameter of only a few nanometers under solvothermal conditions. While Kominami et al. [START_REF] Kominami | Synthesis of thermally stable nanocrystalline anatase by high-temperature hydrolysis of titanium alkoxide with water dissolved in organic solvent from gas phase[END_REF] diluted a titanium alkoxide into an organic solvent, and then injected into the tube, which was subsequently put into the autoclave. Between the tube and the inner wall of the autoclave, there were a certain amount of water or its mixed solvents. Therefore, the formed water vapor could promote the hydrolysis of titanium alkoxide with water dissolved in organic solvent from water vapor, resulting in the formation of crystalline nano TiO 2 in the organic solvent.

(2) Photocatalytic porous film Currently, there are mainly two ways to fabricate photocatalytic porous film: First one is increasing the roughness of the film by controlling the particle size in the film to obtain pores between the particles; Second technique is to use a template to form the film, followed by the removal of the template, leading to a porous film. Doherty et al. [START_REF] Doherty | Preparation and characterization of transparent nanocrystalline TiO 2 films possessing well-defined morphologies[END_REF] prepared a porous TiO 2 film with large specific surface area by controlling the particle size through traditional sol-gel method. O'Regan [START_REF] O' Regan | A low-cost, high efficiency solar cells based on dyesensitized colloidal TiO 2[END_REF] prepared a three-dimensional network of the porous TiO 2 films on conductive glass with roughness factor of 1000 by hydrothermal method film. Zhang et al. [START_REF] Zhang | Low-temperature fabrication of efficient porous titania photoelectrodes by hydrothermal crystallization at the solid/gas interface[END_REF] prepared a mechanically stable mesoporous nanocrystalline titania thick films by hydrothermal treatment of an aqueous mixed paste containing nanocrystalline TiO 2 powder and Ti alkoxide. The Ti monomers were converted into crystalline TiO 2 , which acted as a "glue" to chemically connect the TiO 2 particles.

The template method is most commonly used to fabricate the porous photocatalytic film.

Many works have been reported on the preparation of porous TiO 2 thin film by using block copolymers surfactant, latex spheres, polystyrene particles, polyethylene glycol, colloidal crystals and monomer droplets as templates [START_REF] Bu | Synthesis of TiO 2 porous thin films by polyethylene glycol templating and chemistry of the process[END_REF][START_REF] Du | Synthesis of mesoporous SrCO 3 spheres and hollow CaCO 3 spheres in room-temperature ionic liquid[END_REF]. These templates are eliminated from TiO 2 thin film via an annealing step in a typical procedure, accompanied with the crystallization of amorphous TiO 2 colloidal particles, thereby forming the porous crystalline TiO 2 thin film. In addition, the template can also be removed by etching effect using appropriate solvents. For example, Kotani et al. [START_REF] Kotani | Effects of addition of poly(ethylene glycol) on the formation of anatase nanocrystals in SiO 2 -TiO 2 gel films with Hot Water Treatment[END_REF] put the poly(ethylene glycol) (PEG) in SiO 2 -TiO 2 gel film with hot water treatment, so the PEG could be removed from the film and pores are obtained in the film.

(3) Photocatalytic porous bulk materials At present, there are two types of method to prepare photocatalytic porous bulk materials.

One is to directly load the nano photocatalyst on the porous carrier. Owing to the high adsorption ability of the porous carrier, the mass transfer rate and the photocatalytic efficiency could be increased. The common porous carriers are some minerals such as carbon, zeolite, clay and montmorillonite [ 121 , 122 ], as well as some artificial porous material such as honeycomb ceramics, molecular sieves, porous SiO 2 and Al 2 O 3 etc [START_REF] Bouazza | Enhancement of the photocatalytic activity of pelletized TiO 2 for the oxidation of propene at low concentration[END_REF]. Another technique is to prepare the aerogel materials. Dagan et al. [START_REF] Tomkiewicz | Morphology and photocatalytic activity of TiO 2 Aerogels[END_REF][START_REF] Dagan | Preparation and characterization of TiO 2 aerogels for use as photocatalysts[END_REF] prepared TiO 2 aerogel with specific surface area close to 500 m 2 /g, and both surface and the interior of TiO 2 aerogel could absorb the light due to the relatively high transmittance of the bulk aerogel.

Research on photocatalytic materials with high quantum efficiency

Improving the quantum efficiency of photocatalytic materials was mainly realized by changing the energy band structure, leading to more absorption of visible light or to the more efficient separation and migration of photo-generated carriers. The current improving methods are mainly divided into metal ion doping, nonmetal ion doping, noble metal deposition, organic dye sensitization and use of semiconductor composite.

(

1) Metal ion doping

Metal ion doping is to introduce transition metal or rare earth metal into the TiO 2 lattice structure by physical or chemical method, which leads to the formation of defects or a change in the lattice type. Therefore, doping metal iron in TiO 2 can change the energy band structure (as shown in Figure 1-28), such as a shift of the energy level [START_REF] Yang | Zn-doped CdS nanoarchitectures prepared by hydrothermal synthesis: Mechanism for enhanced photocatalytic activity and stability under visible light[END_REF], narrowing of the band gap [START_REF] Luo | One-pot synthesis of CdS and Ni-doped CdS hollow spheres with enhanced photocatalytic activity and durability[END_REF] and the creation of a doping level [START_REF] Li | Mechanistic study of codoped titania with nonmetal and metal ions: A case of C + Mo codoped TiO 2[END_REF]. The objective is to promote the transfer of photo-generated electrons and holes, to adjust their distribution state and to facilitate photogenerated carrier separation. Xu et al. [START_REF] Xu | The Preparation, Characterization, and their photocatalytic activities of rare-earth-doped TiO 2 Nanoparticles[END_REF] have studied the La 3+ , Ce 3+ , Er 3+ , Pr 3+ , Gd 3+ , Nd 3+ , Sm 3+ and other lanthanide rare earth ions doped TiO 2 , and found that the absorption wavelength of doped TiO 2 was "red shifted", in the order of : Gd 3+ > Nd 3+ > La 3+ > Pr 3+ (Er 3+ )> Ce 3+ > Sm 3+ , Gd 3+ with a higher visible photocatalytic activity. In addition, the photocatalytic activity was also determined by the doping concentration and the optimum doping concentration was about 0.5%. Umebayashi et al. [START_REF] Umebayashi | Band gap narrowing of titanium dioxide by sulfur doping[END_REF] prepared S-doped TiO 2 powder using the oxidation annealing of TiS 2 .

The remaining S could occupy the position of O atom in TiO 2 lattice, forming a Ti-S bond.

The orbital hybridization of S 3p orbital with the O 2p orbital can lead to narrowing of the band gap and to 'red shift' of absorption edge. Yamaki et al. [START_REF] Yamaki | Formation of TiO 2-x F x compounds in fluorineimplanted TiO 2[END_REF] prepared F-doped TiO 2 by ion implantation method and believed that the F atoms replaced the O atoms in the lattice, forming the hole-type defects in the crystal, thus leading to enhanced absorption of visible light. In addition, Yu et al. [START_REF] Yu | Effects of F-doping on the photocatalytic activity and microstructures of nanocrystalline TiO 2 powders[END_REF] considered that the F doping could not only enhance the absorption of visible light, but also improve the crystallinity of anatase TiO 2 , suppress the formation of brookite phase, and prevent the transformation from anatase to rutile phase.

(

3) Noble metal deposition

With the deposition of appropriate amount of noble metal (Pt, Ag, Pd, Au, etc.) on the surface of TiO 2 , the photo-generated electrons will be transferred from TiO 2 (with high Fermi level) to noble metal (with low Fermi level), forming Schottky barrier, which can effectively capture electrons [START_REF] Linsebigler | Photocatalysis on TiO 2 surfaces: Principles, mechanisms, and selected results[END_REF], thereby limiting the recombination of electrons and holes (as shown in Figure 1-29). surface tend to be small. This is because the excessive deposition of noble metal will reduce the contact area between TiO 2 and organic pollutants. It was reported that the photocatalytic activity of Pt/TiO 2 with Pt loading amount of 0.2% was 4 times higher than pure TiO 2 ; while when the Pt loading amount reached 0.5%, the photocatalytic activity of Pt/TiO 2 was lower than pure TiO 2 [START_REF] Zhao | Photodegradation of sulforhodamine-B dye in platinized titania dispersions under visible light irradiation: Influence of platinum as a functional co-catalyst[END_REF].

(4) Organic dye sensitization

The organic dye sensitization is to modify TiO 2 material with organic photoactive substances such as chlorophyll, phthalocyanine, eosin, porphyrin, ruthenium complexes [START_REF] Ohmori | Photovoltaic properties of phthalocyanine based p-n diode evaporated onto titanium dioxide[END_REF][START_REF] Amao | Preparation and properties of dye-sensitized solar cell using chlorophyll derivative immobilized TiO 2 film electrode[END_REF], which can absorb photons and form excited state under the visible light. The photogenerated electrons are then transferred from the excited dye molecules to the TiO 2 , so that the excitation wavelength of the light can be expanded into the visible light range (as shown in Figure 1-30). However, in the photocatalytic reaction, the photosensitized TiO 2 may lead to decomposition of the conjugated structure or chromophore in dye molecule, even highly mineralizing the dye to CO 2 and H 2 O. In addition, most used sensitizers have weak absorption in the near infrared region, which cannot be well matched to the solar spectrum. Therefore, the selection of dye is important to get a better performance. With semiconductor composite, the energy band overlap can occur, leading to the separation and migration of the photo-generated carriers within these semiconductors, thereby effectively inhibiting recombination of photo-generated electrons and holes. Figure 1-32 describes photo excitation process of semiconductor composite by taking TiO 2 -CdS system as an example.

Currently, the compound semiconductors, such as TiO 2 -CdS [START_REF] Hirai | Preparation and photocatalytic properties of composite CdS nanoparticles-titanium dioxide particles[END_REF], TiO 2 -CdSe [START_REF] Ho | Sonochemical synthesis and visible light photocatalytic behavior of CdSe and CdSe/TiO 2 nanoparticles[END_REF], TiO 2 -SnO 2 [START_REF] Liu | An efficient bicomponent TiO 2 /SnO 2 nanofiber photocatalyst fabricated by electrospinning with a side-by-side dual spinneret method[END_REF], TiO 2 -ZnO [START_REF] Marcì | Preparation characterization and photocatalytic activity of polycrystalline ZnO/TiO 2 systems. 2. surface, bulk characterization, and 4nitrophenol photodegradation in liquid-solid regime[END_REF], TiO 2 -WO 3 [149] have shown higher photocatalytic activity than pure TiO 2 . In addition, the combination of TiO 2 with some narrow-bandgap semiconductors (such as CdS, CdSe, WS 2 , etc. [START_REF] Hirai | Preparation and photocatalytic properties of composite CdS nanoparticles-titanium dioxide particles[END_REF][START_REF] Ho | Sonochemical synthesis and visible light photocatalytic behavior of CdSe and CdSe/TiO 2 nanoparticles[END_REF][START_REF] Ho | Preparation and photocatalytic behavior of MoS 2 and WS 2 nanocluster sensitized TiO 2[END_REF]), can lead to a visible photocatalytic activity. The studies have suggested that in order to effectively expand the absorption wavelength range to visible light region, the band gap of the narrow band gap semiconductor must be small enough to produce full absorption of visible light (typically smaller than 1.6 eV);

Meanwhile, the conduction band potential of narrow gap semiconductors should be higher than that of pure TiO 2 to ensure the migration of photo-generated electrons [START_REF] Ho | Preparation and photocatalytic behavior of MoS 2 and WS 2 nanocluster sensitized TiO 2[END_REF].

Bessekhouad et al. [START_REF] Bessekhouad | Bi 2 S 3 /TiO 2 and CdS/TiO 2 heterojunctions as an available configuration for photocatalytic degradation of organic pollutant[END_REF] prepared CdS and TiO 2 mixed precipitation followed by the calcination at 400 o C and obtained CdS/TiO 2 composite. After two hours of visible light illumination, the photocatalytic degradation rate of 4-hydroxybenzoic acid over CdS/TiO 2 was about 60% and this rate was only about 10% with the prue CdS. Ho et al. [START_REF] Ho | Preparation and photocatalytic behavior of MoS 2 and WS 2 nanocluster sensitized TiO 2[END_REF] reported that nanoclusters of MoS 2 and WS 2 were coupled to TiO 2 by an in-situ photoreduction deposition method by taking advantage of the reducing power of the photo-generated electrons from

Multiple oxide semiconductor

The only common semiconductors, with small band gap making possible the full use of sunlight in the visible region (400-700 nm) and at the same time with a conduction band potential higher than that of TiO 2 , are CdS, GaP, GaAsP, and SiC. Therefore, the selection of the narrow band gap semiconductor for combining with TiO 2 is relatively limited. The synthesis of new narrow band gap semiconductors becomes very important and necessary. Of course, these new narrow band gap semiconductors can not only be used to prepare visiblelight compound photocatalysts, but also directly serve as photocatalysts due to their strong visible photocatalytic activity.

Currently, many multiple oxide semiconductors such as TaON (2.5 eV) [START_REF] Yamasita | Recent progress of visiblelight-driven heterogeneous photocatalysts for overall water splitting[END_REF], BiVO 4

(2.4 eV) [START_REF] Harada | Overall water splitting by sonophotocatalytic reaction: the role of powdered photocatalyst and an attempt to decompose water using a visible-light sensitive photocatalyst[END_REF], InVO 4 (2.0 eV) [START_REF] Ye | A novel hydrogen-evolving photocatalyst InVO 4 active under visible light irradiation[END_REF] etc. were prepared and used for decomposing the organic matter or water under visible light irradiation. Furthermore, Ishikawa et al. [START_REF] Ishikawa | Oxysulfide Sm 2 Ti 2 S 2 O 5 as a stable photocatalyst for water oxidation and reduction under visible light irradiation (λ ≤ 650 nm)[END_REF] mixed In summary, the synthesis of new narrow band gap semiconductors has been developing rapidly in recent years. However, the studies are mainly concentrated on the photocatalytic decomposition of water and rarely focused on the visible photocatalytic degradation of organic compounds. In addition, the current new narrow gap semiconductors are mainly prepared by high temperature solid phase method, resulting in very small specific surface area, with therefore limited photocatalytic activity.

Nitride semiconductor

Most recently, a polymeric semiconductor on the basis of a defective graphitic carbon nitride (g-C 3 N 4 ) has stimulated intensive interest in photocatalytic applications (as shown in Figure 1-34) [START_REF] Wang | A metal-free polymeric photocatalyst for hydrogen production from water under visible light[END_REF]. The simple polymer-like semiconductor, made only from carbon, nitrogen and some minor hydrogen content, can function as a metal-free photocatalyst for the extraction of hydrogen from water or for degrading organic pollutants. Among all the allotropes of carbon nitrides, g-C 3 N 4 is regarded as the most stable at ambient conditions and has the smallest direct band gap (2.7 eV) due to the sp 2 hybridization of carbon and nitrogen that forms the π conjugated graphitic planes. Unfortunately, its practical applications are mainly limited by the rapid recombination of photo-generated electron-hole pairs which often results in a low quantum yield and poor photocatalytic activity. To date, a number of methods have been developed to enhance the photocatalytic activity of g-C 3 N 4 , such as doping [START_REF] Zhang | Phosphorus-doped carbon nitride solid: Enhanced electrical conductivity and photocurrent generation[END_REF],

and coupling with grapheme [START_REF] Xiang | Preparation and enhanced visible-light photocatalytic H 2production activity of graphene/C 3 N 4 composites[END_REF] or other semiconductors [START_REF] Kumar | Synthesis of a novel and stable g-C 3 N 4 -Ag 3 PO 4 hybrid nanocomposite photocatalyst and study of the photocatalytic activity under visible light irradiation[END_REF]. Considering the matched band potential between g-C 3 N 4 (the calculated CB and VB edge potentials of g-C 3 N 4 are -1.12 and 1.57 eV) and the most studied photocatalyst (TiO 2 ), it is feasible to combine g-C 3 N 4 with TiO 2 , and the photo-generated electrons can migrate effectively from the CB of g-C 3 N 4 to that of TiO 2 . In this way, a highly efficient photocatalyst with improved quantum efficiency in visible light can be obtained [171]. From Figure 2-1, it can be seen that the annealed samples with the composition 62.5GeS 2 -12.5Sb 2 S 3 -25CsCl gradually lose transmission in the visible region with the extension of annealing time. This is due to the fact that nanocrystals are generated in the glass matrix.

Since the refraction index of these crystals (CsCl) is very different from that of the base glass, the optical losses due to scattering increase with the increasing number or/and size of the crystals [4].

Indeed, when the incident light meets the particles/crystals, the light direction will change dramatically, leading to light scattering or diffusion. Scattering intensity depends on the particle shape, size and concentration. When the size reaches approximately 1/20 of the incident light's wavelength, the scattering will become significant. If the particles are uniformly spherical, models of light scattering can be divided into three domains based on a dimensionless size parameter, α which is defined as

   r 2  (2.1)
where r is the radius of a particle, λ is the wavelength of incident light. Based on the value of α, these domains are (1) α << 1: Rayleigh scattering, (2) α ≈ 1: Mie scattering and (3) α >> 1:

Geometric scattering [5].

In our case, the crystal size, observed with scanning electronic microscope, is typically in the range of 50-100 nm (comparable with our working wavelength). The optical loss is essentially due to Mie scattering which is inversely proportional to the wavelength. semiconductors GeSb [7]: the transmission decreases with increasing wavelength. The absorption here in the middle infrared region is attributed to the presence of free charge carriers, which can be modeled by the theory of Drude-Zener [8]:

N S N m m n m c e FCA 2 2 0 0 2 0 3 2 3 ) * ( 1 4                                   (2.2)
α FCA is the absorption coefficient of free carriers, m* is the effective mass, µ is the mobility, N is the concentration, n is the refractive index, λ is wavelength and S is a constant.

According to the principle of quantum mechanics, there are different values of the exponent of the wavelength, which are determined by the dominant loss mechanism. There are mainly three modes of losses, which are losses by acoustic phonons (λ 1.5 ), losses based on optical phonon (λ 2.5 ) and losses due to impurities (λ 3 or λ 3.5 ). Consequently, the absorption coefficient can be simplified as follows

  3 5 . 2 5 . 1     C B A S FCA    (2.

Glass preparation

The glasses are prepared by the melt-quenching process in vacuum sealed silica tube. The base glasses are synthesized using the polycrystalline germanium (99.999%), antimony (99.999%), selenium (99.999%) and CuI (98%) as starting materials. Because the chalcogenide glass is sensitive to oxygen and water during the melting, the synthesis was done in a silica tube shown in Figure 234. A vacuum trap filled with liquid nitrogen is used to condense eventual volatile products and to improve the vacuum. The vitreous domain of the GeSe 2 -Sb 2 Se 3 -CuI system was determined using a silica tube with 9 mm internal diameter.

The tube was cleaned by hydrofluoric acid for one minute to eliminate dust presented on the wall. It was then thoroughly rinsed with distilled water and dried. Appropriate quantities of each material were weighed according to glass composition and loaded into the silica tube which was then sealed under a typical vacuum of around 10 -5 mbar. The tube containing the mixture was put into a rocking furnace (as shown in Figure 2345). A typical thermal procedure for the synthesis of a chalcogenide glass is shown in Figure 23456. The silica tube containing all the starting materials was heated to 800 o C at the rate of 1.5 o C/min and hold for about 12 h to ensure a thorough reaction and homogenization. Then the furnace was cooled down to 700 o C and the thermal quenching was performed by putting the sealed tube into water at room temperature. The thermal shock in the quenching process generated high mechanical stresses in the glass. In order to reduce these stresses, the tube was annealed at (T g -10) for 3 h. The glass is then cooled slowly and released from the silica tube.

Thermal dynamical consideration of crystallization

Associated with the wide transmission range in the infrared region, chalcogenide glasses have relatively weak chemical bonds, which lead to poor thermomechanical properties [9].

One efficient way to mitigate this weakness is to generate a high number of small crystals. For achieving this objective, it is useful to understand the crystallization mechanism in order to induce controlled crystallization.

The crystallization of a glass is generally governed by a two-step mechanism (Figure 23456789) [10]. One is the nucleation at a temperature near the glass transition temperature. In this temperature region, nuclei are formed in the glass matrix. The other step is the crystal growth around the nuclei. Nucleation is the process of generating a new phase from an old one whose free energy has become higher than that of the emerging new phase. The mechanism of nucleation is the subject of many studies [11][12][13][14][15]. Indeed, germination plays a crucial role in controlling the microscopic properties (size, purity, morphology and crystalline structure) that directly affect the physical and chemical properties of the generated macroscopic crystalline material [16][17][18].

There are two basic types of nucleation during bulk crystallization [19,20]:

(1) Homogeneous nucleation: nucleation occurring arbitrarily throughout the system with no preferential sites for nucleus formation.

(2) Heterogeneous nucleation: formation of the nuclei of a new phase on the surface of an already existing phase boundary.

The rate of nucleation has been considered from the point of view of the classical model for steady-state thermal nucleation. The model based on quantitative determination of the nucleation rate as a criterion of the glass formation was developed by Turnbull [21]. He expressed the nucleation rate using an expression for the steady-state nucleation rate in terms of the classical theory of homogeneous nucleation. The free energy of phase created between the crystalline phase and the supercooled liquid is expressed by the sum of two expressions, which is the function of volume and the surface area [21,22]:

surface volume l tota G G G      (2.4)
In fact, the formation of spherical nucleus thermodynamically affects the free energy of the system in two ways [16]: on one hand, it induces the minimization of the free energy per unit of volume transformed; on the other hand, the spontaneous creation of heterogeneous nucleation in an initially homogeneous system requires the formation of high energy interfaces. The total free energy of the spherical regions is

   2 3 4 3 4 r G r G volume total     (2.5)
where r is the radius of regions, σ is the interfacial energy per surface area unit.

Figure 2-10 shows the evolution of the volume free energy, interfacial energy and overall free energy of the system according to the radius of nucleus [23]. When the temperature decreases below the melting point Tm, the driving force becomes positive. The new phase regions formed with small r are again unstable until they attain a certain critical size r*, at which further growth becomes spontaneous. After having reached the maximum value ΔG* when r = r*, its overall free energy decrease in the system.

The nuclear size r* with maximal free energy corresponding to ΔG* can be found from

condition   0   dr G d total (2.6) So the value of critical radius r* is volume G r     2 * (2.7)
Then the value of critical free energy

Crystal growth

After the formation of stable nuclei in the parent phase, crystallization proceeds by growth of the new phase. The factors controlling crystal growth processes are (1) the rate of heat transfer at the phase boundary, (2) the rate of diffusion of structural units, (3) the rate of the reaction in which the structural units are incorporated into the crystal lattice. In the typical glass-forming systems, the heat of crystallization is mostly small, so that the crystal growth is mostly controlled by molecular transport or by reactions at interface between crystal and the melt. Theoretical description of the rate of crystal growth is based on more or less idealized models that assume a certain character for interface and for sites on the interface where atoms (molecules) are incorporated or removed. These models can be demonstrated for simple example of the growth of new phase from parent phase, where the composition does not change [10].

For the GeSe 2 -Sb 2 Se 3 -CuI glasses studied in this work, we have observed that the nucleation region is quite close to the crystal growth region with overlaps. Therefore, a simplified preparation method is firstly adopted here and the glass ceramics are prepared by directly annealing the glass at crystallization onset temperature T x . A detailed study of the influence of the nucleation and crystal growth on the properties of the glass ceramic will be conducted in the next chapter.

Presentation of characterization techniques

Photoelectrochemical (PEC) measurement

This measurement method is based on a PEC cell with generally an appropriate liquid electrolyte. The principle of the measurement is the following: a certain bias is applied on the semiconductor which is illuminated with chopped light. The light will generate free charge carriers on the surface of the semiconductor in contact with the electrolyte and will consequently increase the current. The applied bias is generally scanned and the appearance or absence of photocurrent at positive or negative bias can be used for the determination of the type of the semiconductor [24,25]. The PEC measurements can either use a two-electrode method or a three-electrode method.

Two-electrode method

The two-electrode method (Figure 2-11) uses the semiconductor as working electrode and a metal as counter electrode. It measures the relationship between reverse bias and photocurrent in this bias [26]. However, in the absence of reference electrode, the real bias applied on the semiconductor is unknown. Besides, we cannot record the data until the system become stabilize. For these reasons, this method is generally not used. The advantage of the three-electrode method is that there is no current flow throw the reference electrode to ensure the stability of the electrodes. Therefore, the measured potential is the applied bias, relative to the reference potential of Ag/AgCl, on the sample. The measurement consists of recoding the current as function of the applied bias under chopped illumination to distinguish the dark and light current.

By putting a chopper between lamp and sample, the dark current and light current can be recorded to calculate the photo current under bias potential. Furthermore, if we can measure the number of photons absorbed by the semiconductor, the qua ntum efficiency of semiconductor can be calculated as follows [27] (2.9)

Other characterizations of glasses and glass ceramics

As described before, glasses used in this study have been prepared by the melt-quenching of the glass contained in a vacuum-sealed silica tube. The characteristic temperatures of the glasses are obtained by thermal analysis which was performed with a differential scanning Optical transmission was measured using a polished disc of 9 mm in diameter and typically 1.5 mm in thickness. The measurements in the visible and near-infrared region were done with a double-beam spectrometer with transmission reproducibility of about ±0.5%. In the mid and far infrared region, they were performed by a Fourier transform spectrometer with a transmission precision of ±2%.

X-ray diffraction data were collected by PANalytical's X-ray diffractonmeter with a step width of 0.02 o and a counting duration of 13 s/step. A polished disc was generally used for this measurement. In this paragraph, we will perform a systematic screening of glass and glass ceramic compositions in relation with the photo-electro-chemical properties. Optimization and detailed structure studies will be performed in the following chapters. In this series of glass ceramic, the main crystalline phase is Cu 2 GeSe 3 , which has relative higher conductivity of about 5 Ω -1 • cm -1 at 0 o C [30].This conductive cubic phase is likely not interconnected inside of the glass ceramics and therefore, all the glass ceramics in this series still show high resistance. Even though the signal is weak, it was still possible to measure the type of the semiconductor. It can be seen from These two samples show p-type semiconductor nature (Table 2-3) in agreement with the fact that selenium is a p-type semiconductor [31,32]. Moreover, the photocurrent increases with the Cu 2 GeSe 3 content. However, the photocurrent intensity is weak for all the samples of this More detailed study is necessary for understanding the reason. It is true that the crystallization of a phase in a quite complex amorphous environment with low mobility of atoms can be very different from the crystallization of a melt with the same composition. Compared to samples of other series, the samples in this series show much higher conductivity with the same content of CuI or Cu 2 GeSe 3 (as shown in Table 234). It indicates that there must be other important parameters, beside of material composition, which affects greatly the conductivity of the glass ceramic. This will be discussed below. Combining the photocurrent and conductivity results of the samples, it can be found that the high photocurrent is based on the high conductivity. Moreover, it can also be seen from Table 234that the conductivity type of the samples in this series is changed from p-type to n-type with the increase of CuI content. It must be mentioned here that the conducting type of the glass ceramics is determined macroscopically by using the PN tester described earlier in this chapter. It is absolutely not in contradiction with the possibility of having simultaneously ptype and n-type semiconducting phases in the same composite materials, as suggested by the [START_REF] Stallings | Synthesis of nanostructured titania powders via hydrolysis of titanium isopropoxide in supercritical carbon dioxide[END_REF] is increased from 0.46, 0.44, 1.35 to 2.11with the decrease of Ge content and the increase of Sb content. Based on the PEC analysis of all these samples, it can be found that the photocurrent intensity first increased and then decreased with the increase of the XRD intensity ratio of Sb 2 Se 3 (221) to Cu 2 GeSe 3 [START_REF] Stallings | Synthesis of nanostructured titania powders via hydrolysis of titanium isopropoxide in supercritical carbon dioxide[END_REF]. It means that increasing the Sb 2 Se 3 content in glass ceramic is favorable for the generation of photocurrent. Sb 2 Se 3 is a direct band gap semiconductor with an energy gap of 1.2 eV [33]. It has been demonstrated that the maximum efficiency of around 30% for single junction solar cells is predicted for materials with band gaps in the range of 1.2-1.4 eV under standard AM1.5 conditions [34,35]. Therefore, Sb 2 Se 3 can be a suitable absorber material for photovoltaic application. During the PEC measurement, when the glass ceramic working electrode is exposed to the simulated sunlight, Sb 2 Se 3 in the glass ceramic can efficiently absorb photons, leading to the formation of electron-hole pairs. This is however not sufficient for generating high photocurrent intensity as demonstrated i n the case of 32GeSe 2 -48Sb 2 Se 3 -20CuI glass ceramic containing Sb 2 Se 3 crystals but with weak photocurrent. Actually, the photocurrent is generated by the migration of photo-generated charge carriers, which is determined by the conductivity of glass ceramic. The four-point probe resistivity test shows that the conductivity of 48GeSe 2 -32Sb 2 Se 3 -20CuI, 40GeSe 2 -40Sb 2 Se 3 -20CuI and 32GeSe 2 -48Sb 2 Se 3 -20CuI is 3.27×10 -7 Ω -1 • cm -1 , 3.13×10 -2 Ω -1 • cm -1 and 1.18×10 -6 Ω -1 • cm -1 respectively, while the conductivities of 64GeSe 2 -16Sb 2 Se 3 -20CuI and 56GeSe 2 -24Sb 2 Se 3 -20CuI cannot be measured here (<10 -8 Ω -1 •cm -1 ). Considering the associated PEC results, it is clear that the high conductivity can promote the migration of photo-generated carriers, leading to the enhancement of photocurrent.

The conductivity of the glass ceramic obviously depends on the material composition and structure. For example, 56GeSe 2 -24Sb 2 Se 3 -20CuI and 40GeSe 2 -40Sb 2 Se 3 -20CuI all contain two crystalline phases of Cu 2 GeSe 3 and Sb 2 Se 3 . Cu 2 GeSe 3 exhibits a metallic behavior at room temperature and its conductivity is about 5 Ω -1 • cm -1 at 0 o C, whereas the conductivity of Sb 2 Se 3 is about ~10 -6 Ω -1 • cm -1 [36]. However, the conductivity of 56GeSe 2 -24Sb 2 Se 3 -20CuI cannot be measured and is much lower than that of 40GeSe

Introduction

The properties of glass ceramics are controlled by the composition and microstructure, whereas the microstructure is often responsible for the mechanical and especially many functional properties, such as optical transmission, photoelectricity, photocatalysis etc [1][2][3].

In the chapter 2, the optimum composition of glass ceramic has been determined after a systematic screening. The highest photocurrent has been obtained with the partially crystallized 40GeSe 2 -40Sb 2 Se 3 -20CuI glass. The main objective of this chapter is to optimize the crystallization process of this selected glass in relation particularly to the microstructure and the photoelectric properties of the obtained glass ceramics.

The crystallization process can determine the microstructure of a glass ceramic. There are two main approaches to get crystals in the base glass [4,5]. The first one, commonly used for preparing glass ceramics, is to crystallize the base glass by a two-stage heat treatment. The first stage is a low-temperature heat treatment at the temperature with a high nucleation rate, leading to the formation of a high density of nuclei in the base glass [6]. A high density of nuclei is favorable for forming a microstructure consisting of a large number of small crystals.

The second stage is a higher temperature heat treatment, resulting in the growth of nuclei at a reasonable rate. The second crystallization technique is a single-stage heat treatment. This method is based on the fact that there is generally temperature overlap of the nucleation and crystal growth processes. Heat treatment at one carefully selected temperature may allow the generation of nuclei and crystal growth occur simultaneously [7].

The nucleation and crystal growth phenomena are sensitive to the composition of base glass which can be thermal-dynamically more or less stable [8]. For the glass composition of 40GeSe 2 -40Sb 2 Se 3 -20CuI, it has found that the necessary overlap occurs between the nucleation and growth rate curves, as demonstrated in chapter 2. Therefore, the two types of crystallization process can be used to prepare glass ceramics from this base glass.

In this chapter, the influence of different crystallization processes on the microstructure and properties of as-prepared glass ceramics was investigated. A reproducible process will be proposed to regulate the structure of such glass ceramic material in order to obtain some specific properties.

2 Glass ceramics prepared by a single-stage heat treatment 2.1 Influence of annealing temperature on the microstructure and properties  (degree) However, with further increasing annealing temperature from 260 o C to 300 o C, the photocurrent and the conductivity significantly decrease. As discussed in chapter 2, the properties of a glass ceramic are highly dependent on its microstructure. The fast growth of crystals at an excessively high temperature is obviously not favorable for building conductive channels inside the glass ceramic. The probable mechanism will be discussed later with the extension of annealing time, indicating that the amount of crystals is increased in the glass ceramics. Moreover, the width of the diffraction peaks gradually decreases with increasing annealing time due to the growth of crystal grains [9]. Consequently, the different conductive regions can be interconnected, leading to the formation of the so-called 'conductive channels' in the glass ceramic. Actually, when the annealing time is increased to 3h and above, the glass ceramics show similar morphology and have conductivity with the same order of magnitude (10 -2 Ω -1 • cm -1 , as shown in Figure 34567).

This morphology increases the conductivity of the glass ceramic, and the sample annealed for 9h obtains much higher conductivity of 2.44×10 -2 Ω -1 • cm -1 . When the annealing time is increased to 3h, the photocurrent generated from this sample is about -25.70 μA at a bias voltage of -0.5 V. The photocurrent is changed with further increase

1E-8 1E-7 1E-6 1E-5 1E-4 1E-3 0,01 0,1 Conductivity ( -1 cm -1 )
Annealing time (h)

1 3 5 7 9
of annealing time, but on the whole the samples show good photocurrents from -18.23 μA to -27.65μA. The highest photocurrent is generated from the sample annealed for 9h. It has been seen that the crystallization process can have a decisive influence on the properties of the glass ceramics. Therefore, we have decided to perform a thermal-dynamic study of the base glass in order to determine the optimum thermal treatment for nucleation and for crystal growth.

Determination of the nucleation rate

The optimum nucleation temperature is determined for the base glass of 40GeSe 2 -40Sb 2 Se 3 -20CuI by using the following method. The base glass is firstly isothermally nucleated in the 200-250 o C range and then a normal DSC analysis is performed with the nucleated samples. The nucleation rate is considered as proportional to surface of the crystallization peak. Once the optimum temperature is fixed, the optimum nucleation time can be determined by using the same procedure. The base glass is nucleated at the optimum temperature for different times (10-240 min), followed by a standard DSC analysis. o C, there is no significant shift in the position of the large endothermic peak as the concentration of nuclei may reach saturation [10]. However, the small endothermic peak disappears, indicating that partial crystallization occurs during nucleation process [11].

Combining the results of XRD patterns, shown in Figure 3-11, it can be deduced that the small endothermic peak is associated to the crystallization of the Sb 2 Se 3 , whereas the large endothermic peak corresponds to Cu 2 GeSe 3 , another crystalline phase of 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramic.

Integral area (J/g) Time (min) crystal size is determined by using the Scherrer equation. The results show that the crystal size changes from 31 nm to 41 nm with the increase of crystallization temperature. It can be clearly seen that the rod-like clustered structure is formed in the glass ceramics crystallized at 280 o C, and some Cu 2 GeSe 3 particles with size of about 100 nm are observed on the clustered structure. However, for the glass ceramics crystallized at 320 o C, the rod-like clustered structure disappears and is replaced by the large particles with smooth surface. This means that at this relatively high crystallization temperature, crystal growth of Sb 2 Se 3 is promoted, leading to bigger but fewer crystals. At the same time, the crystal growth of Cu 2 GeSe 3 , with higher crystallization temperature, seems not affected significantly. These conductive crystals are located between the Sb 2 Se 3 crystals and not on these crystals. This structure will have important consequence on the electric conductivity of the glass ceramic. The photoelectric properties of the glass ceramics shows a similar variation with the conductivity as shown in Figure 3-20. The photocurrent is generated by the transfer of photogenerated carriers, which is determined by the conductivity of glass ceramic. Therefore, there is no photocurrent generation from the sample with poor conductivity. For the sample prepared at 280 o C, the highest photocurrent is achieved, so the optimum crystallization temperature can be determined as 280 o C. However, this photocurrent intensity is much lower than the highest photocurrent intensity of the sample prepared by a single-step heat treatment process discussed above. This is not surprising as we know now that the conductive channels are formed when the two crystals (Sb 2 Se 3 and Cu 2 GeSe 3 ) grow together with relatively low rate. This is the case with the one-step heat treatment at relatively low temperatures. can be seen from Figure 3-22 that, the photocurrent intensity is increased when increasing crystallization time from 0.5h to 1 h. However, with further extension of crystallization time, the photocurrent intensity is decreased although all the samples have the same conductivity.

The difference of photocurrent intensity cannot be attributed here to the difference of Intensity (a.u.)

conductivity. It is probably associated to the separation efficiency of photo-generated charge carriers due to the increase of particle size [12].

Discussion

Based on these results, it can be concluded that the single-stage heat treatment is more efficient than the two-stage heat treatment for the preparation of glass ceramics with enhanced photoelectric properties. For the base glass with the composition of 40GeSe 2 -40Sb 2 Se 3 -20CuI, the crystallization temperature of Sb 2 Se 3 is lower than that of Cu 2 GeSe 3 . According to the DSC analysis, overlap occurs between the crystallization of these two crystals. When the base glass is crystallized at relative low temperature (eg. 260 o C) by using the process of singlestage heat treatment, the Sb 2 Se 3 will crystallize first, thereby providing the growth sites for Cu 2 GeSe 3 crystals. More importantly, the crystallization temperature used for fabricating the glass ceramic is much lower than the crystallization peak temperature of Cu 2 GeSe 3 on the DSC curve, leading to low growth rate of Cu 2 GeSe 3 crystals. In this case, small Cu 2 GeSe 3 crystal grains will grow on the surface of Sb 2 Se 3 crystals and connect to each other, resulting in the formation of the characteristic micron-scale crystalline domains with conductive channels. As shown in Figure 3456, with the increase of crystallization time, these functional domains are interconnected in the whole glass ceramic. The conductivity and the photocurrent intensity are therefore increased.

On the contrary, when crystallizing the base glass through a two-stage heat treatment process, the nucleation stage will significantly decrease the crystallization temperature (as shown in Figure 3-10 and Figure 34567891011121314). For obtaining a high conductivity, it is essential to use a relatively low crystallization temperature (eg. 280 °C) for having comparable crystal growth for Sb 2 Se 3 and Cu 2 GeSe 3 . In this case, the Cu 2 GeSe 3 can follow the growth of Sb 2 Se 3 to form conductive channels.

It has also been observed that even with comparable electrical conductivity, the best glass ceramic obtained with one-step heat treatment shows much higher photocurrent compared to the best glass ceramic obtained with a two-step thermal treatment (nucleation followed by crystal growth). The possible reason can be the following: with the one-step crystallization process at low temperature (eg. 260 °C), the crystal growth rate is low for both Sb 2 Se 3 and Cu 2 GeSe 3 . These two crystals grow together, forming conductive channels and heterojunctions. The presence of these junctions will be evidenced in the following chapter. These junctions will of course promote charge separation, leading to high photocurrent. With the glass ceramic obtained by using a two-step heat treatment, the pre-nucleation obviously promotes the crystal growth of Sb 2 Se 3 . Even with a crystal growth at relatively low temperature (eg. 280 °C), it is more difficult for Cu 2 GeSe 3 to follow the growth rate of Sb 2 Se 3 and to cover the Sb 2 Se 3 . Conductive channels can be formed with lower density and consequently less heterojunctions, explaining the lower photocurrent intensity.

Conclusion

In this chapter, the influence of different crystallization processes on the microstructure and photoelectrical properties of the glass ceramics is investigated. The results show that the crystallization process has a critical influence on the microstructure and the photoelectrical properties of the glass ceramics. The ideal crystallization scenario seems to the following: the Sb 2 Se 3 rods firstly grow in the matrix, followed by the in-situ growth of Cu 2 GeSe nanocrystals on the Sb 2 Se 3 . Compared with the two-step heat treatment process, the singlestep process at a low temperature is a more efficient strategy to obtain the rod-like clustered composite structure. It is believed that with crystallization at low temperature, the crystal growth rate is low and comparable for both Sb 2 Se 3 and Cu 2 GeSe 3 . This will facilitate the growth of the conductive and cubic Cu 2 GeSe 3 on the needle-like Sb 2 Se 3 , leading to interconnected conductive channels and heterojunctions which are of course favorable for photo-generated carrier separation. This will be further discussed and evidenced in the following chapter.

Introduction

Charge photogeneration, dissociation and transport to external collectors are the key steps for efficient conversion of solar energy via photovoltaic effect or via photocatalytic production of chemical fuels. One single material cannot be optimum for all the three tasks.

Organic solar cells have attracted much attention because of the low cost, flexibility and ease of processing. The introduction of internal donor-acceptor heterojunctions was a key breakthrough for improving the efficiency of polymer photovoltaic solar cells [1,2]. The principle was to attribute one specific function (charge photogeneration or charge diffusion) to different molecules in order to overcome the mismatch between the exciton diffusion length (around 10 nm) and the optical absorption depth (more than 100 nm). In this way, each molecule can be optimized for a dedicated function [3,4]. With the creation of an interpenetrating network, the carrier collection efficiency is improved by two orders of magnitude in comparison with bi-layer flat heterojunction devices [5][6][7]. Recently, semiconductor photocatalysts have attracted considerable attention for potential application in photocatalytic degradation of organic pollutents or water splitting [8]. Heterojunction structure has been extensively studied because of its efficiency for promoting separation of photogenerated charges [9][10][11][12][13].

Based on the studies described in previous chapters, we have tried to develop a new strategy for preparing an efficient interpenetrating inorganic heterojunction structure with two relatively narrow bandgap semiconductors, Sb 2 Se 3 and Cu 2 GeSe 3 . These two crystals have complementary direct bandgap with, 1.0-1.2 eV for Sb 2 Se 3 [9,14] and 0.78 eV for Cu 2 GeSe 3 [15], allowing efficient harvesting of solar energy from UV up to 1600 nm. Sb 2 Se 3 has a high absorption coefficient of 10 5 cm -1 and is considered as a promising semiconductor for building heterojunctions to potentially achieve hi gh energy conversion efficiency [16]. It has recently attracted more and more attention [17,18], but its low conductivity of 10 -6 Ω -1 • cm -1 [19] is certainly a big issue for charge transport which can be ensured by the Cu 2 GeSe 3 phase with a conductivity of 10 1 Ω -1 • cm -1 [20]. The structure of Sb 2 Se 3 is orthorhombic with cell parameters a=1.113 nm, b=1.142 nm and c=0.385 nm. It grows typically as rod-like crystals which can be easily interconnected [21]. This may be an important feature for building up an interconnected network. Whereas the structure of Cu 2 GeSe 3 seems more diversified as cubic (a=0.555 nm), tetragonal (a=0.560 nm, b=0.548 nm) and orthorhombic (a=1.186 nm, b=0.396 nm and c=0.546 nm) phases are both reported [22,23]. With our experimental conditions, only cubic phase was precipitated.

In this chapter, a detailed study on the microstructure and properties of the photoelectric glass ceramic will be carried out and the mechanism of charge photogeneration, separation and transport will also be discussed. All the studies will be performed on the 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramic, obtained by annealing the precursor glass at 284 °C during 1 hour, except elsewhere specified.

2 Microstructure of the 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramics

TEM and EDS analysis

SEM analysis

However, the conductivity of the glass ceramics is determined not only by the formation of the conductive microdomain but also by the interconnection of these conductive microdomains at macroscopic scale. Actually, there exist three typical composite morphologies (as shown in Figure 4-3). In the first case (Figure 4-3(a)), the base glass is crystallized incompletely, and the conductive microdomain is therefore separated by the nonconductive glass phase, leading to the poor conductivity of the glass ceramic. In the second case (Figure 4

AFM analysis

The atomic force microscope (AFM) is one kind of scanning probe microscopes (SPM).

Unlike traditional microscopes, the AFM does not rely on electromagnetic radiation, such as photon or electron beams, to create an image. An AFM is a mechanical imaging instrument that measures the three dimensional topography and physical properties of a surface [25].

Basic principles and presentation of the equipment

As shown in Figure 4-4, the AFM consists of a cantilever with a sharp probe at its end that is used to scan the specimen surface. Whe n the tip of the probe is brought close to a sample surface, forces between the tip and the sample lead to a deflection of the cantilever, according to Hooke's law [26]. Typically, the deflection is measured using a laser spot reflected from the top surface of the cantilever into an array of photodiodes. If the tip is scanned at a constant height, a risk would exist that the tip collides with the surface, causing damage. Hence, in most cases, a feedback mechanism is employed to adjust the tip-to-sample distance to maintain a constant force between the tip and sample.

The AFM can be operated in number of modes depending on the application. In our experiments, semi-contact mode (tapping mode) was chosen to scan the surface potential.

The surface potential mapping is carried out by scanning Kelvin probe microscope (SKM) which is integrated in an AFM system NTEGRA Prima (NT-MDT Company). Pt/Ir-coated silicon probes NSG01/Pt (NT-MDT Company) are used for these measurements. Otherwise topography features will contribute to the surface potential image [27]. A 15 nm lift height was used in our experiment to prevent topographic cross talk. The scan frequency was 1 Hz, i.e. the acquisition time of a 256 x 256 pixel image was 8 min.

AFM surface potential

The influence of composite morphology on the conductivity of the glass ceramics can be further revealed by the AFM surface potential measurement. In our previous studies, we have found that the surface potential of the sample is decreased after crystallization, and the asprepared glass ceramic achieves a relatively high conductivity at the same time (as shown in Figure 45). Therefore, it can be deduced that the conductive region has lower surface potential than the non (or poorly)-conductive region in the same image for this system. For the fabrication of an efficient photoelectric material, there are three critical factors that need to be considered: (1) the material should efficiently respond to the whole solar spectrum as wide as possible; (2) the absorbed photons should be efficiently converted into the free charge carriers that can be efficiently separated; (3) the free charge carriers should be efficiently transported to the injecting contacts by the built-in electric field with high carrier mobility [28,29]. Therefore, optical absorption property, photo-generated carrier lifetime, carrier mobility and conductivity are believed to be several main semiconductor properties, which determine whether such material can be a potential candidate for photoelectric applications [30]. The optical absorption behavior of the glass ceramic is mainly determined by its composition and structure. There are mainly two crystalline phases, Sb 2 Se 3 and Cu 2 GeSe 3 , in the glass ceramic. The Sb 2 Se 3 has a direct band gap of 1.2 eV, corresponding absorption edge wavelength is 1033 nm [31]. Whereas the direct band gap of Cu 2 GeSe 3 is about 0.78 eV, leading to the absorption of infrared light with wavelength up to 1590 nm [32,33]. It has been Absorbance (a.u.)

Optical absorption

Wavelength (nm)

confirmed that the doping of I 2-ions in the lattice of Sb 2 Se 3 in the chapter 2. I-doping is expected to reduce the energy gap of Sb 2 Se 3 due to the formation of impurity states in the band gap [34,35]. Furthermore, the Cu 2 GeSe 3 phase can also be non stoichiometric, leading to a decrease of the bandgap width. For example, copper vacancies, with Ge on Cu sites, and selenium interstitials are expected to be the dominant defect species, which will affect the band structure [15]. As an acceptor defect in Cu 2 GeSe 3 , copper vacancy may originate a shallow level in the band gap, leading to the acceptor level-to-valence band transition [15,36]. Besides of impurity and defect levels, the possible band overlapping in Sb 2 Se 3 -Cu 2 GeSe 3 composite system may also lead to the red shift of the optical absorption spectrum. The above factors together determine such wide absorption property of the glass ceramic.

Current-voltage characteristics

Current-voltage (I-V) measurements are important for examining the electrical properties of semiconducting materials. To measure these characteristics, two 1.8-mm metallic circle contacts were deposited onto the same face of each sample by electron beam evaporation.

These contacts should have a separation of 2.5 mm. Silver contact was used for n-type samples while gold contact was used for p-type samples so that ohmic contacts could be obtained (as shown in Figure 456789). Then 2-point method was then carried out in the dark to get the I-V characteristics.

Charge carrier lifetime

In semiconductor physics, charge carrier's lifetime is defined as the average time it takes to recombine with an opposite charge carrier. It is one of the most important parameters for the characterization of semiconductor used in the preparation of power electronic devices and photovoltaic solar cells [37].

Chapter 4

-132 -

Theoretic considerations

There are various recombination mechanisms even within a uniformly doped piece of semiconductor. The main types of recombination are intrinsic or radiative recombination, Shockley-Read-Hall (SRH) recombination and Auger recombination. SRH recombination is determined by the amount of impurities and defects in the semiconductor. For a semiconductor with indirect band-gap, SRH recombination is more important. Radiative (also known as band-to-band) recombination occurs when an electron falls from its state in the conduction band into the empty state in the valence band, by combining with the hole, resulting in light emission. Auger recombination is a process in which an electron and a hole recombine as a band-to-band transition, but the resulting energy is given off to another electron in the conduction band instead of just emitting a photon. The involvement of a third particle affects the recombination rate which makes it different from radiative recombination [34].

The lifetime of carriers in a bulk material  b is composed of a SRH lifetime  SRH , radiative lifetime  rad and Auger lifetime  A with the relation [38,39]:

    A rad SRH b 1 1 1 1    (4.1)
Nowadays, various techniques have been proposed to measure minority carrier lifetime, such as PCD (photo conductivity decay), QSSPC (quasi-steady-state photoconductance) and SCCD/OCVD (short-circuit current/open-circuit voltage decay), etc [30]. In this work, the carrier lifetime is determined by the microwave photoconductivity decay (µ-PCD) method. In the µ-PCD method, the laser injection will produce electron-hole pairs, leading to the increase of conductivity. When the external laser injection is removed, the conductivity presents exponential decay over time, which indirectly reflects the extinction trend of minority carriers.

The detection of µ-PCD method uses the principle that the variation of microwave is proportional to that of electrical conductivity. Compared with traditional PCD, microwave detection can not only guarantee the small deviation of minority lifetime, but also avoid the difficult preparation of ohmic electrodes [40].

Equipment

Lifetime measurements are carried out using a commercially available µ -PCD measurement system from Semilab (WT-2000) (as shown in Figure 4 The existence of nano-junctions is of course extremely favorable for charge separation and this is probably one of the reasons why the composite materials show an extremely long lifetime of minority charge carriers of about 16 µs (as shown in Figure 4-13), compared for example to similar and well studied chalcopyrite CIGS with a lifetime typically shorter than 0.1 µs [42]. The inhomogeneity of lifetime is associated with the composite character of the material with randomly distributed conductive channels which will be illustrated below.

Hall effect

When a current goes through a conductor subjected to a perpendicular external magnetic field, a voltage difference will appear, transverse to the electric current and to the magnetic field. This phenomenon is called Hall effect, which was discovered in 1879 by Edwin Herbert Hall while he was working on his doctoral degree at Johns Hopkins University in Baltimore, Maryland [43]. Hall effect can provide some essential parameters of a semiconductor such as conductivity type, carrier concentration, carrier mobility etc. It is now one of the most important methods for characterizing semiconductor materials.

Theoretical considerations

When a magnetic field is applied to a semiconductor with a direction which is not parallel to the moving direction of the charges, these charges will experience a force, called the Lorentz force. It will force the moving charge carriers deviate from its original straight path.

Consequently, positive and negative charges will be accumulated on both sides of the material. This leads to the creation of an electric perpendicular to the direction of the current. This electric field will eventually balance the Lorentz force to establish equilibrium. Therefore, a steady electrical potential difference, the Hall voltage, is created.

Since the glass ceramic disks are all circular, four symmetrical electrodes need to be uniformly deposited around the sample. The thickness of sample is about 1/3 of its diameter. where I is the current across the plate length, B is the magnetic field, d is the thickness of the disc, e is the elementary charge of electron and n is the charge carrier density. The Hall coefficient R H is defined as the ratio of the induced electric field to the product of the current density and applied magnetic field:

ne IB d V B j E H x y H R 1     (4.3) Chapter 4 -136 -
where j x is the current density of the carrier electrons, and E y is induced electric field.

Measurement

In our experiment, Accent HL5500PC Hall Effect measurement system as shown in Figure 4-15 is used to characterize our samples. The HL5500PC is a turn-key, high performance Hall System for the measurement of resistivity, carrier concentration and mobility in semiconductor. It can be seen that the glass ceramic exhibits relatively high Hall mobility in the range of 10 2 cm 2 • V -1 • S -1 , which is in the same order of magnitude with that of CIGS [44,45]. High charge carrier mobility means a strong conductive ability. This is in agreement with the conductivity of about ~10 -2 Ω -1 • cm -1 , measured with the four-point method, for this glass ceramic. Based on the analysis above, it can be found that there are two key factors affecting the photoelectric conduction of the glass ceramic, which finally determines the photoelectric property. One is the formation of a semiconductor heterojunction between Sb 2 Se 3 and Cu 2 GeSe 3 , which can lead to a more efficient inter-particle carrier separation. Another one is the formation of a conductive channel in the glass ceramic, resulting in the obvious enhancement of conductivity, which can promote the mobility of photo-generated carriers.

The possible mechanism of charge photogeneration, separation and transport can be proposed in Figure 4 • OH etc, which can decompose most organic pollutants [5,6]. Therefore, efficient generation and separation of photo-generated carriers is the key for efficient semiconductor photocatalytic reaction [7][8][9].

In order to increase the amount of photo-generated charge carriers, the development of narrow band-gap semiconductors, for having a high absorption of photons in the visible region, has attracted much attention [10][11][12].New narrow band-gap semiconductors has been rapidly developed in recent years, including particularly sulfide semiconductors [13],

multiple oxide semiconductors [14,15] and nitride semiconductors [16]. However, with the decrease of band-gap width, the photo-generated electrons and holes recombine more easily.

Since only the free photo-generated carriers can further react and form active radicals, leading to the decomposition of organic pollutants, the more easily-occurred recombination in some narrow band-gap semiconductors will decrease their photocatalytic activity although the y can absorb more photons. Therefore, a number of methods have been developed to inhibit the recombination of photo-generated electron-hole pairs, such as doping [17,18] and coupled semiconductor [19]. Recently, fabricating nano-heterojunction through in-situ growth route has become a rising method to develop highly efficient semiconductor composite. The nanoheterojunction can have more efficient interparticle electron transfer between the two components due to the size effect of particles and the intimate contact between heterophase nanoparticles [19,20]. Consequently, nano-scale semiconductor heterojunction with narrow band gap can not only extend the absorption wavelength range to visible region, but also inhibit the recombination of photo-generated electron-hole pairs. This kind of junctions has become a promising strategy for improving the visible-light photocatalytic activity.

In the previous chapters, we have presented a novel glass ceramic with interesting photoelectric activity. This glass ceramic exhibits a relative intense absorption from UV up to the infrared region (240-2000 nm). More importantly, there forms nano-heterojunctions between Sb 2 Se 3 and Cu 2 GeSe 3 crystals, leading to an efficient separation of photo-generated electrons and holes. Such unique features make it interesting as a photocatalyst. Therefore, in this chapter, the 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramic is used as a photocatalyst for decomposing organic pollutants.

Evaluation method of photocatalytic activity

The photocatalytic activity of the selected glass ceramics is evaluated by decomposing methyl orange (MO), taken as an example, in aqueous solution under visible light illumination.

Figure 5-1 shows the molecular structure of MO, the UV-vis absorption spectrum and the color of MO aqueous solution., It can be seen that the MO aqueous solution has a characteristic large absorption peaked at 464 nm, attributed to the -N=N-chromophore [21].

This chromophore can be oxidized during photocatalytic reaction, leading to the discoloration of the solution. Therefore, the absorbance of the solution can be used to monitor the photocatalytic activity and a low residual absorbance indicates of course a high reaction activity of the photocatalyst [22,23]. where A 0 and A t are absorbance at the maximum absorption wavelength of the solution, respectively at time 0 and time t.

The photocatalytic activity can also be monitored by a normalized concentration of MO.

According to the Lambert-Beer law [24], the normalized concentration of the solution equals to the normalized maximum absorbance, i.e., C t /C 0 = A t /A 0 . Here C t and C 0 are the concentration of the solutions respectively at time t and time 0. Therefore, the normalized concentration can be directly determined by using a UV-vis spectrophotometer.

Photocatalytic activities of different glass ceramics

A simple method for comparing samples with different photocatalytic activities is to evaluate the ptotocatalytic decomposition ratio of a selected pollutant as function of time.

Since the difference of photocatalytic activities between various photocatalysts is not always obvious or tends to be particularly small at the initial or final stage of the photocatalytic reaction, it is better to choose a detection time in the middle stage of the photocatalytic reaction. The choice of an optimum detection time is mainly dependent on the particle size of the photocatalyst (for powder photocatalyst), photocatalyst concentration, pollutant concentration and other ptotocatalytic reaction conditions, such as the wavelength and the illumination intensity of light source, and the uniformity of the suspension, etc [25][26][27].

Particle size is an important parameter influencing strongly the photocatalytic activity of powder photocatalyst. Smaller particle size will induce higher photocatalytic activity due to its higher specific surface [28]. In all our experiments, the glass ceramic is manually grinded into powders and then the powders are successively sieved through 30 μm and 20 μm sieves to obtain powders with two different size characteristic, <20 μm or between 20 μm ~30 μm.

Whereas the powder with particle size >30 μm will be further grinded. The photocatalytic activities of these two powders are evaluated by decomposing MO. However, for the glass ceramic powders with smaller particle size (<20 μm), a part of the powders greatly smaller than 20 μm, remains floating on the surface of the MO solution even after ultrasonic dispersion for 10 min. The other part of the powders is in suspension in the MO solution. This phenomenon is mainly due to the hydrophobic feature of the powders and the small particle size. The floating powders will decrease the light absorption of the glass ceramic powders in the solution, leading to lower photocatalytic activity (as shown in Figure 5-3). Therefore, we choose the glass ceramic powders with particle size between 20 μm and 30 μm for the photocatalytic reaction. rate. There is no doubt that the photocatalytic reaction is accelerated by increasing photocatalyst dosage at the same MO concentration. When fixing the dosage of photocatalyst (such as 10 mg or 20 mg), a lower MO concentration has a higher decomposition rate.

However, for the high photocatalyst dosage of 40 mg, although the decomposition ratio at low MO concentration (reaction system g) achieves a high P value of 90.8%, this rate is lower than that at high MO concentration (reaction system h). The likely reason is that the decomposition of MO in the reaction system g achieves high level before the detection time of 1.5 h. Moreover, it also can be seen that excessively high initial concentration of MO leads to the formation of sawtooth profiles on the absorption curve (as shown in Figure 5-4(c)). In this case, a large amount of MO dye and intermediates will absorb the light, so the light intensity on the photocatalyst is weakened, leading to the decrease of photocatalytic reaction rate [29,30]. Consequently, a suitable photocatalytic reaction conditions can be determined as a reaction system with a photocatalyst dosage of 40 mg and a MO concentration of 40 mg/L. Here, the photocatalytic activity of two series glass ceramics is evaluated. Figure 5-5(a) shows the UV-vis absorption spectra of MO solutions before and after photocatalytic reaction with different 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramics, prepared by a single-stage thermal annealing at 260 o C for different times. The photocatalytic activity of the glass ceramic is enhanced with an increase of annealing time from 1 h to 3 h. However, when further extending the annealing time above 7 h, the photocatalytic activity of glass ceramic is gradually decreased. The photocatalytic activity of a material is known to be dependent on the photo-generation and separation of charge carriers [7]. The comparison of structure, conductivity and photocatalytic decomposition rate of MO between different glass ceramics is shown in Figure 567. The analysis of these data leads to the following remarks:

(1) The formation of nano-scale heterojunctions and conductive microdomains leads to high photocatalytic activity. It is well known that some materials with high conductivity and carrier mobility such as a graphene can efficiently promote the separation of photo-generated electron-hole pairs when coupled with semiconductors [31]. The above results indicated that highly conductive Cu 2 GeSe 3 nanocrystals grow on the surface of Sb 2 Se 3 rods, forming nanoscale heterojunctions and conductive microdomains in the glass ceramic. This structure can lead to a more efficient inter-particle carrier separation, thereby increasing the photocatalytic activity of these glass ceramics. Therefore, for the glass ceramic prepared at 260 o C for 1h (as shown in Figure 567), there is too much amorphous phase in the glass ceramic, considerably slowing down the transport of photo-generated charge carriers. The photocatalytic activity is enhanced when the base glass is annealed for 3 h leading to a more complete crystallization.

(2) Higher conductivity of glass ceramic does not lead necessarily to higher photocatalytic activity. The conductivity of a glass ceramic depends on the interconnection of the conductive microdomains at macroscopic scale in the glass ceramic. In order to fabricate this structure, it is necessary to extend annealing time or appropriately increase crystallization temperature (as shown in Figure 567). However, the increase of annealing time or crystallization temperature will increase the size of crystal grains. In general, a larger crystal grain may increase the migration distance of charge carriers, the electrons and holes photo-generated in the grain will take a longer time to reach the surface, leading to a higher recombination probability.

Therefore, it can be seen form Figure 5-7 that the photocatalytic activity of glass ceramic is decreased with the increase of annealing time. Moreover, the photocatalytic activities of the glass ceramics prepared by the single-stage heat treatment are overall superior to that of the glass ceramics prepared by the two-stage heat treatment. Especially for the glass ceramic prepared at 260 o C for 1h, although it still contains a certain amount of glass phase, its photocatalytic activity (with decomposition rate of 37.2%) is comparable to that of the fullycrystallized glass ceramic (with decomposition rate of 36.0%), prepared at 290 o C with the two-stage heat treatment. This result further confirms that an excessively large size of crystal grains will lead to poor photocatalytic activity probably due to a higher recombination probability of photo-generated carriers.

(3) Crystal defect is another factor which can influence the photocatalytic activity of glass ceramic. It can be observed that the photocatalytic activity is slowly increased when changing crystallization temperature from 280 o C to 320 o C for the glass ceramics prepared by the twostage heat treatment. The most likely reason is that the rod-like Sb 2 Se 3 grow up more smoothly with less defects at the relatively high temperature, thanks to the higher mobility of the atoms leading to less stress in the crystals. Since crystal defects may become a recombination center of photo-generated electron-hole pairs [32], a high degree of crystallinity can increase photocatalytic activity due to lower defect level in the crystals. 4 Photocatalytic reaction process of MO over the glass ceramic 464 nm is associated with the azo bond (-N=N-) [21] and is asymmetric. The broad absorption peak can be deconvoluted into two peaks centred at 464nm and 443 nm (as shown in Figure 56789). This peak centred at 443 nm is attributed to a reverse photo-isomerization of the azobenzene compound, namely, an azobenzene molecule will change from thermodynamically stable trans structure (E-type) to cis structure (Z-type) under excitation of visible light (as shown in Figure 5-10) [33,34]. The absorption peak centred at 276 nm is due to benzene charge-transfer absorption [35]. With the progress of the photocatalytic reaction, the absorption peaks centred at 464 nm, 443 nm and 276 nm decrease gradually, and a new absorption peak ranged from 228 nm to 250 nm appears simultaneously, indicating that some intermediates may be generated. This new absorption peak can be deconvoluted into two peaks located at 228 nm (226 nm) and 250 nm, which probably correspond to 4-hydroxybenzenesulfonic acid sodium and 4dimethylaminophenol respectively according to the Table 5-2. The N-demethylation of MO will occur due to the attack of hydroxyl radicals on the N,Ndimethyl groups, generating sulfanilic sodium and N,N-dimethylamino aniline, leading to the consequent decolorization [40,41]. However, these products cannot be detected in the UV-vis absorption spectra ranged from 210 nm -600 nm as its maximum absorption wavelength is shorter than 200 nm [41,42]. With the N-demethylation of MO, the hydroxylation with deamination of the former two products occurs simultaneously, leading to the formation of 4-hydroxybenzenesulfonic acid sodium and 4-dimethylaminophenol [41], which have characteristic UV-vis absorption peaks at 228 nm and 251 nm respectively according to Table 5-2. These results are in good agreement with the deconvolution of the absorption peak.

Therefore, the amount of these two intermediates increases with the photocatalytic reaction time, which is illustrated by the increase of the corresponding absorption peaks (as shown in leading to the shift of the total absorption peak to shorter wavelength region. Therefore, the likely photocatalytic degradation pathway for MO in this reaction system can be proposed as shown in Figure 5-11. The glass ceramic photocatalyst exhibits a good oxidative deamination ability and its photocatalytic reaction process is different from that of most other photocatalysts reported before [36,44]. Actually, this process is similar to the electrochemical oxidation process [41]. first-order rate constant, to describe these two reactions for comparison. It is clear that the glass ceramic show a more complex photocatalytic mechanism, as described above.

It can be seen from Figure 5-12(a) and (b) that the glass ceramic exhibits an excellent photocatalytic degradation rate of 0.02108 h -1 , which is about 3.5 times higher than that of P25 TiO 2 nano powders (as shown in equation 5.2 and 5.3). It is interesting to note that a different photocatalytic reaction system with a photocatalyst dosage of 100mg and a MO concentration of 10 mg/L is adopted for the P25 TiO 2 nano powders due to its low visible photocatalytic activity. Therefore, under the same photocatalytic reaction conditions, the glass ceramic achieves a higher photocatalytic degradation rate of 0.03135 h -1 (as shown in Figure 5-12c and equation 5.4), which is about 5.3 times higher than that of P25 TiO 2 nano powders.

Moreover, considering that the particle size of the glass ceramic powders between 20 μm-30 μm is much larger than P25 TiO 2 nano powders (about 30 nm), the photocatalytic activity of the glass ceramic might be greatly improved just by decreasing the particle size. The fitting degree for the glass ceramic is quite low (R=0.87175 and 0.91559 respectively), indicating that the photocatalytic reaction over the glass ceramic is more complicated than P25 TiO 2 nano powders and there may exist some unknown time-dependent factors affecting the photocatalytic reaction. The formation of nano-scale heterojunctions and conductive microdomains can induce an efficient separation of photo-generated carriers, leading to a high photocatalytic activity.

Therefore, it is unnecessary to form conductive channels in the whole glass ceramic for the photocatalytic application. Moreover, in order to form conductive channels, the necessary increase of annealing time or/and temperature may decrease the photocatalytic activity due to its relatively large crystal grain size. This photocatalyst has a good oxidative deamination ability and high photocatalytic activity, demonstrating its potential as an efficient photocatalyst. Consequently, the photo-generated charge carriers firstly can be efficiently separated by the p-n junctions, and then transported though the conductive channels in the glass ceramic, leading to exceptionally long lifetime of charge carriers (around 16 µs) and high photocurrent (at least 100 times higher than any of Sb 2 Se 3 and Cu 2 GeSe 3 individually). Moreover, this glass ceramic also exhibited a high photocatalytic activity and the demonstration was made through the photocatalytic decomposition of methyl orange selected as an example. Therefore, this glass ceramic can be considered as an efficient photocatalyst, even still much more work need to be done.

To the best of our knowledge, this is the first report on interpenetrating heterojunctions in glass ceramic systems or even in inorganic materials. Similar microstructures are even relatively new for organic materials and have received much attention. This kind of bulk heterojunctions has led to a key breakthrough for improving the efficiency of polymer photovoltaic solar cells.

This pioneer work could be interesting for developing a totally new family of glass ceramic with highly efficient photoelectrical and photocatalytic activities. The final objective is of course to contribute to the design of totally new materials for converting efficiently the solar energy into electrical or chemical energy. Of course, up to now, the exact structure-property relationship, the charge photo-generation mechanism as well as certain properties are still to be understood. We believe that this work is only the beginning of an important research topic.

Future studies could be carried out in the directions:

(1) Developing new material systems with still more efficient semiconducting phases. The two innovative features provided by this unique structure are conductive channels and randomly distributed nano p-n junctions along the conductive channels. The fabrication of conductive channels opens the possibility of using materials with low electrical conductivity but excellent properties (high absorption, efficient charge photogeneration…) for photovoltaic application. Moreover, a multiple-heterojunction system, which may contain several semiconductors having complementary bandgap for better covering the entire solar spectrum, could also be fabricated along a conductive channel to increase the conversion efficiency of the whole solar spectrum.

(2) Developing new material systems to achieve low cost. The glass ceramic is prepared by crystallizing the base glass with the composition of 40GeSe 2 -40Sb 2 Se 3 -20CuI. Therefore, even if copper is totally with germanium to form Cu 2 GeSe 3 , 75% of germanium stays in the glass matrix. Based on our current understanding, the residual glassy phase does not play an important role for obtaining the interesting photoelectric properties of the glass ceramic.

Consequently, the content of germanium could probably be decreased to achieve low cost. Mots-clés: verres de chalcogénures ; vitrocéramiques; hétérojonction; effet photovoltaïque; photocatalyse

Figure 1 .

 1 Figure 1. Courbe de transmission du verre (a) et de la vitrocéramique (b) de la même composition 60GeSe 2 -30Sb 2 Se 3 -10CuI

Figure 2 .

 2 Figure 2. Domaine vitreux dans le système GeSe 2 -Sb 2 Se 3 -CuI et les séries de verres étudiés

Figure 3 .

 3 Figure 3. Méthode à trois électrodes pour mesures de photocourant

Figure 4 .

 4 Figure 4. Diffraction aux rayons X des vitrocéramiques dans la série 4 ((50-x/2)GeSe 2 -(50-x/2)Sb 2 Se 3 -xCuI)

Figure 5 .

 5 Figure 5. Photocourant des deux vitrocéramiques dans la série 4

  35GeSe 2 -35Sb 2 Se 3 -30CuI optimale de nucléation (figure6) et ensuite à faire croitre les nucléis à une température supérieure pendant un temps relativement court (1 heure environ).

Figure 6 .

 6 Figure 6. Optimisation de la température et du temps de nucléation pour 40GeSe 2 -40Sb 2 Se 3 -20CuI

Figure 7 .Figure 8 .

 78 Figure 7. Conductivité électrique de la vitrocéramique 40GeSe 2 -40Sb 2 Se 3 -20CuI glass en fonction du temps de recuit utilisé pour la fabrication

Figure 9 .

 9 Figure 9. Images HRTEM de la vitrocéramique 40GeSe 2 -40Sb 2 Se 3 -20CuI

Figure 10 .

 10 Figure 10. Image en potentiel surfacique de deux vitrocéramiques représentative avec différente conductivité (a) <10 -8 Ω -1 •cm -1 ; (b) 2.1×10 -2 Ω -1 •cm -1

Figure 11 .

 11 Figure 11. Distribution des éléments dans la vitrocéramique

Figure 12 .

 12 Figure 12. Diffraction aux rayons X de vitrocéramiques avec un taux différent en CuI

  sont bien ohmiques. Avec ces mesures, il est ainsi clair que la vitrocéramique contient des jonctions p-n. En effet, nous avons observé beaucoup de jonctions créées par des contacts entre des Sb 2 Se 3 dopé (type) et des Cu 2 GeSe 3 (type p). La vitrocéramique massive contient donc un grand nombre des jonctions p-n orientées aléatoirement. Cette structure peut être simplifiée électroniquement comme l'illustre la figure 13 (b), permettant d'expliquer le comportement I -V de la vitrocéramique.

Figure 13 .

 13 Figure 13. Coubre I-V d'une vitrocéramique et simplification des jonctions p-n

Figure 14 .

 14 Figure 14. Distribution spatiale de la durée de vie des porteurs de charge minoritaire sur la surface d'une vitrocéramique

Figure 15 .

 15 Figure 15. Photocourant des Sb 2 Se 3 , Cu 2 GeSe 3 et de la vitrocéramique contenant ces deux cristaux

Figure 16 .

 16 Figure 16. Mécanisme de la photo-génération, de la séparation et du transport des charges dans les vitrocéramiques contenant des hétérojonctions

Figure 17 .Figure 18 .

 1718 Figure 17. Montage de la photocatalyse

Figure 19 .

 19 Figure 19. Cinétique de la photo-décomposition du méthyle orange avec la vitrocéramique de chalcogénure et un produit commercial, TiO 2 P25 comme catalyseurs

Figure 1 - 1

 11 Figure 1-1 World Marketed distribution of energy use by fuel type, -2100 (EJ/a)

Figure 1 - 2

 12 Figure 1-2 The Earth's energy budget from the sun [2]

Figure 1 - 3

 13 Figure 1-3 Solar thermal applications

Figure 1 - 4

 14 Figure 1-4 Solar thermal utilization forms in the world[4]

Figure 1 - 5

 15 Figure 1-5 Typical concentrating technologies[5]

Figure 1 - 6 Figure 1 - 7 Figure 1 - 8

 161718 Figure1-6 Central receiver power tower system[9]

Figure 1 - 9 ( 1 )( 2 )

 1912 Figure 1-9 Main processes of photocatalytic water splitting[16]

  oxynitride and oxysulphide materials may be more suitable for spontaneous water splitting than metal oxides, as their valence band edges are closer to the O 2 /H 2 O potential, and those with a suitable band-gap around 2 eV are able to have band edges that overlap the water splitting potentials.

Figure 1 -

 1 Figure 1-10 Principle of water splitting using semiconductor photocatalysts

Figure 1 - 2 -

 12 Figure 1-11 typical setup for PEC water splitting In addition, the photo-generated electrons can also react with the adsorbed O 2 on the semiconductor surface and form various active substances, such as •O 2 -, • HO 2 , H 2 O 2 , • OH, etc (formula 1.1-1.5). Meanwhile, the photo-generated holes will react with the adsorbed H 2 Oor OH -ions on the semiconductor surface and generate hydroxyl radicals • OH (formula 1.6, 1.7). These active substances can decompose most organic compounds (formula 1.8). In

Figure 1 -

 1 Figure 1-12 Solar energy photoelectric conversion system[20]

  -13, when the semiconductor p-n junction is exposed to the sunlight, the electrons in the valence band are excited to the conduction band, resulting in generation of holes in the valence band and electrons in the conduction band, thus forming an electron-hole pair (left part in Figure1-13). The formed electron-hole pairs are separated by the p-n junction electric field and a current proportional to the number of electron-hole pairs created flows through an external circuit (right part in Figure1-13). If the photon energy is less than E g , it doesn't carry enough energy to break a covalent bond and free an electron for conduction. Since electrons cannot occupy the forbidden states between the valence and conduction bands, a photon with energy less than E g cannot be absorbed and will pass through the sample. Only photons which energy is greater than the semiconductor band gap (E g ) can be absorbed, and release excess energy in the form of heat.

Figure 1 -

 1 Figure 1-13 Working principle of semiconductor based solar cell

Figure 1 -

 1 Figure 1-14 Efficiencies of best research solar cells[31]

  the solar cell is exposed to sunlight, an electron-hole pair is generated, and the cell starts converting solar energy into electricity by photovoltaic effect of silicon p-n junction (as shown in Figure1-15, Figure1-16).

Figure 1 -Figure 1 -

 11 Figure 1-15 Schematic presentation of silicon doping[32]

( 2 )

 2 The energy of the incoming photons is greater than the band gap, so the excess energy is lost as heat. (3) The Fermi levels of both n-type and ptype silicon are always inside the band gap of silicon so the open-circuit voltage is smaller than the band gap. (4) The series and shunt resistors generated from the contacts and lattice defects in the silicon consume some of the electricity that is generated. (5) Electron-hole pair recombination instead of conversion to electricity.

Figure 1 -

 1 Figure 1-17 Illustration of energy loss in a p-n junction silicon solar cell: (1) unabsorbed light, (2) excess energy loss as heat, (3) Fermi level mismatch with band gap, (4) transfer loss, and (5) electron-hole recombination

Figure 1 -

 1 Figure 1-18 Production process for crystalline silicon solar cells (© 2010 Sharp)

Figure 1 -

 1 Figure 1-19 Schematics of various crystalline solar cell structures (a) PERL cell; (b) HIT cell; (c) BC-BJ cell; (d) MWT cell

Figure 1 -

 1 Figure 1-20 A typical structure of amorphous silicon solar cell[41]

Figure 1 -

 1 Figure 1-21 Mechanism for polymer solar cells[45]

  -22).Polymer solar cells have certain disadvantages including their low efficiency and short lifetime. Nevertheless, their numerous benefits such as low-cost, lightweight, and flexibility, can justify the current international investment and research in developing new polymeric materials, new combinations and structures to enhance efficiency and to achieve low-cost and large-scale production within the next years.

Figure 1 -

 1 Figure 1-22 Two kinds of structure of polymer solar cells

Figure 1 -( 1 )( 2 )( 4 )

 1124 Figure 1-23 Schematic diagram of dye-sensitized solar cells (DSSC) [55]

Figure 1 -Figure 1 -

 11 Figure 1-24 Schematic of multi-junction GaAs solar cells

Figure 1 -

 1 Figure 1-26 Schematic of CdTe solar cells (Zn 2 SnO 4 as a buffer layer)

Figure 1 -

 1 Figure 1-28 Energy levels of impurity metal ions in TiO 2[START_REF] Choi | The role of metal ion dopants in quantumsized TiO 2 : correlation between photoreactivity and charge carrier recombination dynamics[END_REF]

Figure 1 -

 1 Figure 1-29 Schematic diagram of photoexcitation in metal-modified semiconductor

Figure 1 -

 1 Figure 1-30 Schematic diagram of dye sensitization mechanism of TiO 2[START_REF] Wu | Photoassisted degradation of dye pollutants. V. self-photosensitized oxidative transformation of rhodamine B under visible light irradiation in aqueous TiO2 dispersions[END_REF]

Figure 1 -

 1 Figure 1-31 Scheme of band edge position of several semiconductors (pH=1)[START_REF] Hagfeldt | Light-induced redox reactions in nanocrystalline systems[END_REF]

  TiO 2 particles. Quantum size effects altered the energy levels of the conduction and valence band edges in the coupled semiconductor systems, which favored the interparticle electron transfer (as shown in Figure 1-33). In addition, the coupled systems were believed to act in a cooperative manner by increasing the degree of charge carrier separation, which effectively reduces recombination.

Figure 1 -

 1 Figure 1-33 Energy levels of the conduction and valence band edges versus normal hydrogen electrode (NHE) for pure TiO 2 , WS 2 and MoS 2 with various sizes at pH 7 [150]

Figure 1 -

 1 Figure 1-34 (a) Crystal structure of g-C 3 N 4 and (b) relationship of the photocatalytic hydrogen production activity and the wavelength

Figure 2 - 1

 21 Figure 2-1 Photograph and transmission curves of 62.5GeS 2 -12.5Sb 2 S 3 -25CsCl annealed at 290 o C for different times (a) photograph; (b) visible and near-infrared transmission; (c) infrared transmission

Figure 2 - 2

 22 Figure 2-2 Tenacity (K c ) and critical load for crack initiation (P c ) of 62.5GeS 2 -12.5Sb 2 S 3 -25CsCl annealed at 290 o C for different times

  3) A, B and C are corresponding parameters. The strong transmission decrease in the middle infrared region suggests a relatively high concentration of free charge carriers. Therefore, we have decided to study in detail, the semiconducting properties of the GeSe 2 -Sb 2 Se 3 -CuI system.

Figure 2 - 3

 23 Figure 2-3 Transmission curve of (a) glass 60GeSe 2 -30Sb 2 Se 3 -10CuI and (b) glass ceramic obtained after annealing at 260 o C for 1 h

Figure 2 - 4

 24 Figure 2-4 Setup used for preparation of chalcogenide glasses

Figure 2 - 6

 26 Figure 2-6 Thermal procedure for chalcogenide glass synthesis

Figure 2 - 7

 27 Figure 2-7 Photograph of glass rod and discs

Figure 2 - 8

 28 Figure 2-8 Glass-forming region in the ternary diagram of GeSe 2 -Sb 2 Se 3 -CuI

Figure 2 - 9

 29 Figure 2-9 Schematic diagram of the crystallization process

Figure 2 -

 2 Figure 2-10 Evolution of free energy according to the radius of nucleus when the temperature is below the melting point T < T m

Figure 2 -

 2 Figure 2-11 Schematic presentation of two-electrode method

Figure 2 -

 2 Figure 2-12 Photograph of semiconductor working electrode

  photo

2 - 14 .

 214 calorimeter (DSC) by using a sample weight of 8 mg and a heating rate of 10 o C/min. The temperature reproducibility is typically ±1 o C. From the thermal analysis, the glass transition temperature T g and crystallization onset temperature T x can be graphically obtained as shown in Figure The difference ΔT = T x -T g is a good indication of the stability of a glass against the crystallization.

Figure 2 -

 2 Figure 2-14 Typical DSC curve, showing the characteristic temperatures for a glass

Figure 2 -

 2 Figure 2-15 Photograph of a Semilab PN tester

Figure 2 -

 2 Figure 2-16 Schematic presentation of four-point probe method for resistivity measurement

4. 1

 1 Series 1: (1-x)(80GeSe 2 -20Sb 2 Se 3 )-100xCuI, x=0, 0.1, 0.2 This series of samples is prepared by introducing various amount of CuI into the base glass of 80GeSe 2 -20Sb 2 Se 3 . The DSC curve of this base glass is shown in Figure 2-17, and exhibitsits glass transition at ~308 o C (T g ). It can also be seen that the glass is stable against crystallization, as there is not any exothermic peak, which can be attributed to the crystallization phenomenon, up to 450 o C. However, when CuI is introduced into the base glass, the glass transition temperature decreases and three crystallization peaks appear and are shifted towards lower temperatures with the increasing CuI content[3]. From the XRD pattern shown in Figure2-18, it can be determined that the exothermic peaks at 309 o C and 278 o C are attributed to the crystallization of Cu 2 GeSe 3 , which is facilitated with increasing CuI content. In addition, the Sb 2 Se 3 crystallizes in both 72GeSe 2 -18Sb 2 Se 3 -10CuI and 64GeSe 2 -16Sb 2 Se 3 -20CuI samples when thermally annealed at higher temperature. The peaks at 393 o C and 390 o C correspond to the crystallization of Sb 2 Se 3 . The last exothermic peak respectively at 410 o C and 379 o C for 72GeSe 2 -18Sb 2 Se 3 -10CuI and 64GeSe 2 -16Sb 2 Se 3 -20CuI samples is attributed respectively to the crystallization of GeSe 2 and maybe Sb.

Figure 2 - 1 Figure 2 -Figure 2 -

 2122 Figure 2-17 DSC curves of various glasses of Series 1

Figure 2 -

 2 Figure 2-19 XRD patterns of different glass ceramics

  photocurrent signals with very weak intensity. The photocurrent is here defined as the difference between the current under illumination and the current in the dark at a certain bias voltage. The photocurrent is highly dependent of composition and it disappears with the composition containing 20% of CuI (as shown in Figure2-20(c)). Furthermore, semiconductor type of the glass ceramics is measured and the results are listed in Table2-1.Both glass ceramics having photocurrent are p-type semiconductors, and it is difficult to determine the semiconductor type of all glass ceramics with weak or nonexistent photocurrent.

Figure 2 -

 2 Figure 2-20 Current-voltage characteristics of different glass ceramics in Series 1 under chopped visible light illumination

Figure 2 -

 2 Figure 2-22 shows a typical dark and light current-voltage characteristic of the glass ceramics from the Series 2, indicating that no photocurrent is generated. The most likely reason for this result is that all the glass ceramics have poor conductivity (as shown in Table 2-2).

Figure 2 -

 2 Figure 2-21 XRD patterns of different glass ceramics in Series 2

Figure 2 - 3 Figure 2 -

 232 Figure 2-23 XRD patterns of different glass ceramics in Series 3

  42GeSe 2 -28Sb 2 Se 3 -30CuI 4.4 Series 4: (1-x)(50GeSe 2 -50Sb 2 Se 3 )-100xCuI, x = 0, 0.1, 0.2, 0.3 This series of samples is prepared by introducing various amount of CuI into the base glass of 50GeSe 2 -50Sb 2 Se 3 . The glass ceramics are prepared by annealing the corresponding glass under the conditions detailed in Table 2-4. Figure 2-25 shows the XRD patterns of different glass ceramics. All observed diffraction peaks of the base glass belong to Sb 2 Se 3 , but the relative intensity of each peak is different from the standard PDF card (No. 01-072-1184).

Figure 2 -

 2 Figure 2-25 XRD patterns of different glass ceramics in Series 4

Figure 2 -

 2 Figure 2-26 Current-voltage characteristics of different glass ceramics in Series 4 under chopped visible light illumination

Figure 2 -

 2 Figure 2-26(c).

Figure 2 -

 2 Figure 2-29 shows the XRD patterns of different glass ceramics with the fixed CuI content of 20 mol%. It can be seen that the XRD intensity ratio of Sb 2 Se 3 (221) to Cu 2 GeSe 3 (111) is

Figure 2 -

 2 Figure 2-29 XRD patterns of different glass ceramics with 20 mol% CuI

2 -40Sb 2

 22 Se 3 -20CuI although the main crystalline phase of 56GeSe 2 -24Sb 2 Se 3 -20CuI is the electrically conductive Cu 2 GeSe 3 . The observation of 56GeSe 2 -24Sb 2 Se 3 -20CuI under scanning electron microscope shows that the Cu 2 GeSe 3 crystals are fully dispersed in the glass matrix and surrounded by the amorphous phase (as shown in Figure 2-30(a)). Therefore, the insulating glass phase between Cu 2 GeSe 3 crystals leads to the poor conductivity of the glass ceramic.

Figure 2 -

 2 Figure 2-30 SEM images of the different glass ceramics: (a) 56GeSe 2 -24Sb 2 Se 3 -20CuI; (b) 40GeSe 2 -40Sb 2 Se 3 -20CuI

Figure 2 -

 2 Figure 2-31 HRTEM images of as-prepared glass ceramic of 40GeSe 2 -40Sb 2 Se 3 -20CuI

Figure 2 -

 2 Figure 2-32 XRD patterns of different glass ceramics with various CuI content: (a) 2θ=10°~90°; (b) 2θ=30°~32°

Figure 2 -

 2 Figure 2-33 STEM images and EDS elemental maps of the glass ceramic

Figure 3 -

 3 Figure 3-1 shows the XRD patterns of the prepared glass ceramics of 40GeSe 2 -40Sb 2 Se 3 -20CuI with different annealing temperatures. In order to promote the crystallization at lower temperature, the annealing time is extended from 1 h (annealing duration used in the systematic screening described in the previous chapter) to 5 h. It can be seen that no crystal is formed when annealing the sample at 230 o C. With the increase of the annealing temperature, the intensity of diffraction peaks gradually increases, indicating that high temperature favors the crystal growth. However, the annealing temperature has no effect on the crystallized phases and all the samples annealed above 230 o C contain the main crystalline phase of Sb 2 Se 3 and a second phase of Cu 2 GeSe 3 .

Figure 3 - 1

 31 Figure 3-1 XRD patterns of the glass ceramics with different annealing temperature

Figure 3 -

 3 Figure 3-2 shows SEM images of the glass ceramics prepared at different annealing temperatures. The SEM image of the sample prepared at 230 o C displays a typical morphology of glass, which is consistent with the XRD result. When annealing the base glass at 260 o C for 5h, a large amount of crystals are formed. The particles grow up when further increasing the annealing temperature to 300 o C, leading to bigger crystals. The conductivity of the glass ceramics are shown in Figure 3-3. With the increase of annealing temperature, the conductivity first increases and then decreases. It seems that an excessive annealing temperature is not favorable for the formation of the conductive channels mentioned in the chapter 2.

Figure 3 - 2 Figure 3 - 3

 3233 Figure 3-2 SEM images of the glass ceramics obtained with different annealing temperature: (a) 230 o C; (b) 260 o C;(c) 300 o C

Figure 3 - 4

 34 Figure 3-4 Current-voltage characteristics of the glass ceramics with different annealing temperatures under chopped visible light illumination

Figure 3 - 5 Figure 3 - 6

 3536 Figure 3-5 XRD patterns of the glass ceramics prepared at 260 o C for different annealing times (The diffraction intensities are normalized)

Figure 3 - 6 Figure 3 - 7 Figure 3 - 8

 363738 Figure 3-6 SEM images of the glass ceramics with different annealing times: (a) 1h; (b) 9h

Figure 3 - 8

 38 Figure 3-8 Current-voltage characteristics of the glass ceramics with different annealing times under chopped visible light illumination

Figure 3 -

 3 Figure 3-9 gives the definitions of the parameters used for analyzing the DSC results,including DSC peak temperature (T p ), DSC peak height (δT) p , and DSC peak area (integral area). In this study, (δT) p is considered as proportional to the nucleation rate, whereas the crystallization peak area proportional to the density of nuclei formed during the nucleation stage[10].

Figure 3 - 9 Figure 3 -

 393 Figure 3-9 Different parameters used for analyzing the DSC measurement Temperature

Figure 3 -Figure 3 -

 33 Figure 3-10 DSC curves for the 40GeSe 2 -40Sb 2 Se 3 -20CuI samples nucleated at different temperatures for 1h

Figure 3 -Figure 3 -

 33 Figure 3-12 Plot of the (δT) p versus the nucleation temperature

Figure 3

 3

Figure 3 -Figure 3 -

 33 Figure 3-14 DSC curves for the 40GeSe 2 -40Sb 2 Se 3 -20CuI samples nucleated at 220 o C for different times

Figure 3 -

 3 Figure 3-16 shows the XRD patterns of different glass ceramics which were prepared at different temperatures. The X-ray diffraction peaks of the samples can be assigned to Sb 2 Se 3 and Cu 2 GeSe 3 , and the intensity of peaks all increase with the crystallization temperature, indicating that the amount and size of crystals are increased in the glass ceramics. The Sb 2 Se 3

Figure 3 -

 3 Figure 3-16 XRD patterns of the glass ceramics prepared by crystallizing the prenucleated glasses at different temperatures for 1 h

  Figure 3-14, it can be seen that the crystallization temperatures used here are higher than the crystallization temperature determined by the DSC analysis, leading to the significant growth of Sb 2 Se 3 and Cu 2 GeSe 3 .

Figure 3 -Figure 3 -

 33 Figure 3-17 SEM images of the cross section of the glass ceramics prepared by crystallizing the pre-nucleated (220 °C for 90 min) glasses at different temperatures for 1 h: (a) 270 o C; (b) 280 o C; (c) 290 o C; (d) 300 o C; (e) 310 o C; (f) 320 o C

Figure 3 -Figure 3 -

 33 Figure 3-18 High multiple SEM images of the glass ceramics prepared by crystallizing the pre-nucleated glass at different temperature for 1 h: (a) 280 o C; (b) 320 o C

Figure 3 -

 3 Figure 3-19 Conductivities of the glass ceramics prepared by crystallizing the prenucleated glass at different temperatures for 1 h

Figure 3 -Figure 3 -

 33 Figure 3-20 Current-voltage characteristics of the glass ceramics with different crystallization temperature under chopped visible light illumination

  the increase of crystal content and size. The Sb 2 Se 3 crystal size is determined by using the Scherrer equation and changes from 31nm to 35 nm with the increase of crystallization time.

Figure 3 -

 3 Figure 3-21 XRD patterns of the glass ceramics prepared by crystallizing the prenucleated glass at 280 o C for different time

Figure 3 -

 3 Figure 3-22 Current-voltage characteristics of the glass ceramics with different crystallization time under chopped visible light illumination

Figure 4 -

 4 Figure 4-1 shows the nanostructure and composition mapping of the glass ceramic. It can be seen from Figure 4-1(b) that the glass ceramic exhibits the characteristic rod-like clustered structure, which is in agreement with the SEM images discussed in the previous chapter. At the edge of the rod, nano p-n junctions between n-type Sb 2 Se 3 and p-type Cu 2 GeSe 3 crystals are formed [15, 24], as shown in Figure 4-1(a).Figure 4-1(c) shows the cross section of the

Figure 4 -

 4 1(c) shows the cross section of the rod-like clustered structure. It can be clearly seen from the EDS elemental mappings that antimony element is mainly enriched in the "light color" region, corresponding to Sb 2 Se 3 phase, whereas copper and germanium elements mainly coexist in the "deep color" region, corresponding to Cu 2 GeSe 3 . The similar phenomenon is also observed in the EDS elemental maps of vertical section (as shown in Figure4-1(d)). The results indicate that the Cu 2 GeSe 3 nanocrystals grow on the surface of Sb 2 Se 3 rods, and mainly connect to each other along the rod orientation, leading to the formation of such crystalline domains. Each domain is composed of the rod-like Sb 2 Se 3 crystals and the interconnected Cu 2 GeSe 3 nanocrystals exist between the rods (as shown in Figure4-2). The oriented and relatively conductive Cu 2 GeSe 3 aggregate provides a conductive microdomain for the glass ceramic.

Figure 4 - 1

 41 Figure 4-1 HRTEM and STEM images and EDS elemental maps of the glass ceramic

  -3(b)), the Sb 2 Se 3 crystals grow up excessively and interconnect, resulting in the separation of the conductive microdomains as well as the corresponding poor conductivity. In the last case (Figure 4-3(c)), the conductive microdomains can interconnect well and form conductive region. It is easy to image that the conductive channels inside the glass ceramic are formed by a non (or poorly)-conductive Sb 2 Se 3 network, covered by a conductive Cu 2 GeSe 2 layer. Consequently, the glass ceramics with the third kind of morphology achieve relatively high conductivity.

Figure 4 - 3

 43 Figure 4-3 SEM images of typical morphology of the glass ceramics with different conductivities (a) 3.2×10 -8 Ω -1 • cm -1 ; (b) <10 -8 Ω -1 • cm -1 ; (c) 2.1×10 -2 Ω -1 •cm -1

Figure 4 - 4

 44 Figure 4-4 Schematic presentation of atomic force microscope (AFM)

Figure 4 - 5 Figure 4 -

 454 Figure 4-5 Histogram of AFM surface potential for the base glass before and after ceramisation

Figure 4 - 6 1 3

 461 Figure 4-6 Typical AFM surface potential images of the glass ceramics with different conductivity: (a) <10 -8 Ω -1 • cm -1 ; (b) 2.1×10 -2 Ω -1 • cm -1

Figure 4 -

 4 Figure4-7 shows the absorption spectrum of the glass ceramic, obtained by measuring the diffuse reflectance. It can be seen that the glass ceramic exhibits a wide absorption from UV up to the infrared region (1800 nm). This absorption specturm is well matched with the solar spectrum. . Furthermore, the optical band gap widths is determined by plotting (αhν) 1/2 versus photon-energy (as shown in Figure4-8). The glass ceramic displays two obvious absorption edges, corresponding to the optical band gap widths of 1.0 eV and 0.24 eV respectively.

Figure 4 - 7 Figure 4 - 8

 4748 Figure 4-7 Absorption spectrum of the glass ceramic

Figure 4 - 9 4 - 131 -

 494131 Figure 4-9 Photograph of a sample with two contacts for I-V measurements

Figure 4 -

 4 Figure 4-10(b)).

Figure 4 -

 4 Figure 4-10 I-V curve of the glass ceramic

  -11). The WT-2000 is a powerful tabletop measurement platform usable for different semiconductor materials. The system includes a XY measurement stage with the possibility of lifetime mapping over a surface.

Figure 4 -

 4 Figure 4-11 µ-PCD measurement system (WT-2000)

Figure 4 -

 4 Figure 4-12 Lifetime map and distribution

Figure 4 -

 4 Figure 4-13 Spatial distribution of the minority charge carrier lifetime over the surface of the glass ceramic disk

Figure 4 -

 4 Figure 4-14 Schematic of a Hall effect measurement

Figure 4 -

 4 Figure 4-15 Hall effect measurement system (HL5500PC)

Figure 4 -

 4 Figure 4-16 shows Hall mobility of the glass ceramic versus magnetic flux density in different configurations (positive, negative and averaged). It can be seen that the glass

Figure 4 -

 4 Figure 4-16 Hall mobility of the glass ceramic versus magnetic flux density

Figure 4 -

 4 Figure 4-17 PEC measurement of Sb 2 Se 3 , Cu 2 GeSe 3 and the glass ceramic under a chopped light source

- 18 .Figure 4 -

 184 Figure 4-18 Mechanisms of charge photo-generation, dissociation and transport in the photoelectric glass ceramic

Figure 5 - 1

 51 Figure 5-1 Molecular structure of MO and a typical UV-vis absorption spectrum and color of MO aqueous solution

Figure 5 - 2

 52 Figure 5-2 Schematic presentation of the photocatalytic reactor

Figure 5 -Figure 5 -

 55 Figure 5-3 UV-vis absorption spectra of MO solutions before and after photocatalytic reaction with different particle sizes of glass ceramic powders for the same duration

Figure 5 -

 5 Figure 5-4 UV-vis absorption spectra of MO solutions before and after photocatalytic reaction under different conditions detailed in table 5-1

Figure 5 -

 5 photocatalytic activity, leading to the most obvious color-fading of the MO solution. With an increase of the annealing temperature, the photocatalytic activity of the glass ceramics first decreases and then increases (as shown in Figure5-6).

Figure 5 -

 5 Figure 5-5 (a) UV-vis absorption spectra and (b) corresponding color of MO solutions before and after photocatalytic reaction over different 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramics prepared by a single-stage heat treatment at 260 o C for different times

Figure 5 - 7

 57 Figure 5-7 SEM images, conductivities and photocatalytic decomposition rates of MO for different glass ceramics prepared (a) by a single-stage heat treatment at 260 o C for different times; (b) by a two-stage heat treatment with different temperatures.

Figure 5 -

 5 Figure 5-8 illustrates the UV-vis absorption spectra of MO solutions photocatalytic degraded by the glass ceramic for different times. It can be seen that the absorption peak of the initial MO solution shifts from its characteristic value of 464 nm to 470 nm, which may be due to a detection error generated by the relatively high concentration of MO solution, leading to a very high absorption and a very low signal level. The absorption peek of MO solution at

Figure 5 - 8 Figure 5 - 9 Figure 5 -

 58595 Figure 5-8 UV-vis absorption spectra of MO solutions before and after photocatalytic reaction for different times

Figure 5 - 9 )

 59 . After 6 hours of photocatalytic reaction, the absorption intensity at 228 nm increases more obviously. This is likely because a hydroquinone (the most hydroxylated direct product) is produced from the degradation of 4-hydroxybenzenesulfonic acid sodium and 4-dimethylaminophenol. Hydroquinone has a maximum absorption peak of 223 nm[43],

Figure 5 -

 5 Figure 5-11 Main pathways proposed for MO photodegraded by the glass ceramic under visible light.

Figure 5 -

 5 Figure 5-12 shows the plots of ln(C 0 /C t ) versus photocatalytic reaction time for the glass ceramic photocatalyst and the commercially available P25 TiO 2 nano powders. We take the pseudo-first-order model, expressed by the equation of ln(C 0 /C t )= kt, where k is the apparent

Figure 5 -

 5 Figure 5-12 Plots of ln(C 0 /C t ) versus photocatalytic reaction time for different photocatalytic reaction systems

  photogeneration, separation and transport. The Cu 2 GeSe 3 nanocrystals grew on the surface of iodine-doped Sb 2 Se 3 rods, leading to the formation of p-n junctions. At the same time, the relatively conductive Cu 2 GeSe 3 nanocrystals connected to each other along the doped Sb 2 Se 3 rod orientation, thereby forming a conductive microdomain. The interconnected conductive microdomains eventually provided macroscopic conductive channels for the glass ceramic.

RESUME( 1 )( 2 )( 3 )

 123 Une nouvelle famille de vitrocéramiques, avec une microstructure inédite, a été fabriquée par une cristallisation contrôlée des verres dans le système GeSe 2 -Sb 2 Se 3 -CuI. L'influence de la composition et du processus de cristallisation des verres de base, sur la microstructure et sur l'intensité du photocourant des vitrocéramiques a été étudiée. Une composition optimisée, le 40GeSe 2 -40Sb 2 Se 3 -20CuI, a été particulièrement étudiée avec des résultats suivants : Après une étude systématique , il a été constaté que cette composition donne la plus forte intensité de photo-courant parmi tous les verres étudiés dans ce système pseudo-ternaire GeSe 2 -Sb 2 Se 3 -Cul. Il a été également démontré que le photo-courant généré par différentes vitrocéramiques est non seulement déterminé par la composition, mais aussi par la microstructure composite de la vitrocéramique, qui est déterminée par le processus de céramisation. Ce processus de céramisation a ensuite été optimisé. Par rapport au procédé de traitement thermique en deux étapes, le procédé en une seule étape à basse température est une stratégie plus appropriée pour obtenir une microstructure efficace, favorisant la séparation des charges, construisant des canaux conducteurs et donnant une intensité de photo-courant élevée dans la vitrocéramique. La microstructure composite inédite, discutée ci-dessus est composée de micro-domaines conducteurs interconnectés, formées par des cristaux Sb 2 Se 3 faiblement conducteur en forme de tiges, couverts par des nano-cristaux de Cu 2 GeSe 3 beaucoup plus conducteurs. Le procédé le plus probable de la photo-génération efficace des charges est le suivant: les photons sont efficacement et essentiellement absorbés par Sb 2 Se 3 ainsi que par Cu 2 GeSe 3 . Les hétérojonctions formées par les Sb 2 Se 3 du type n et les Cu 2 GeSe 3 du type p, favorisent la séparation de charges, tandis que les Cu 2 GeSe 3 interconnectées et conductrices fournissent des canaux conducteurs et jouent ainsi le rôle de collecteur efficace de charges. Il en résulte ainsi une très longue durée de vie des porteurs de charge (environ 16 µ s) et un fort photo-courant (au moins 100 fois plus élevé que Sb 2 Se 3 ou Cu 2 GeSe 3 séparément). La formation de nano-hétérojonctions entre les cristaux Sb 2 Se 3 et Cu 2 GeSe 3 dans un seul microdomaine peut conduire à une séparation efficace des électrons et des trous photo-générés. Par conséquent, pour application photo-catalytique, il n'est pas nécessaire de former des canaux conducteurs (conducteurs interconnectés des micro-domaines) dans l'ensemble de la vitrocéramique.De plus, la formation de ces canaux conducteurs, nécessiterait une augmentation de la durée ou/et la température de recuit, pouvant conduire à une diminution de l'activité photo-catalytique à cause de la taille relativement grande des grains cristallins. Les vitrocéramiques optimisées montrent une bonne capacité de désamination oxydative et une forte activité photo-catalytique en général, démontrant ainsi son potentiel en tant que photo-catalyseur efficace.

  Much higher hydrogen production efficiency can be achieved if spontaneous water splitting occurs, without the need to apply an electrical bias. A semiconductor capable of spontaneous water splitting has a bandgap ≥ 2 eV[17] with conduction band energy more negative than that of the H + /H 2 redox potential, and valence band energy more positive than that of the O 2 /H 2 O redox potential as shown in Figure1-10[ 18 ]. Under these conditions, charge transfer between the

semiconductor and the electrolyte is energetically favorable. Most metal oxides have a valence band edge at a similar energy, well below the O 2 /H 2 O potential. Therefore, to have a conduction band edge above the H + /H 2 potential, the material will have a wide band-gap (over 3 eV). Narrower bandgap materials with band edges that overlap the water splitting redox potentials (such as CdS and CdSe) are unstable. Maeda and Domen

[18]

suggest that

  Serpone et al.[START_REF] Serpone | Spectroscopic, photoconductivity, and photocatalytic studies of TiO 2 colloids: Naked and with the lattice doped with Cr 3+ , Fe 3+ , and V 5+ cations[END_REF] used transition metal ions M (M is Cr 3+ , Fe 3+ , V 5+ ) as dopant-ions to prepare M/TiO 2 photocatalytic materials, forming doping level within the band gap of TiO 2 , thus extending the absorption wavelength of TiO 2 to the visible range of 410-620 nm (3.0 eV-2.0 eV). Yamashita et al.[START_REF] Yamashita | Degradation of propanol diluted in water under visible light irradiation using metal ion-implanted titanium dioxide photocatalysts[END_REF] doped TiO 2 with V 5+ , Mn 3+ , Fe 3+ , Co 2+ , Ni 2+ , Cr 3+ and other transition metal ions, effectively expanding the absorption wavelength of photocatalytic material to the visible light. Choi et al.[START_REF] Choi | The role of metal ion dopants in quantumsized TiO 2 : correlation between photoreactivity and charge carrier recombination dynamics[END_REF] studied 21 kinds of metal ions doped TiO 2 and found that 0.1-0.5% of Fe 3+ , Mo 5+ , Ru 3+ , Os 4+ , Re 5+ , V 5+ and Rh 3+ could greatly improve the photocatalytic activity of TiO 2 . In contrast, doping Co 3+ or Al 3+ reduced the photocatalytic activity of TiO 2 likely because they became the electron and hole recombination centers.Nonmetal ion doping is to introduce ions such as N, C, S and F into the oxygen lattice. The orbital hybridization of the p orbital of these ions with the O 2p orbital can lead to the shift of the valence band and to the decrease of band gap, so that the material can absorb the visible light.As early as 1986, Sato et al.[START_REF] Sato | Photocatalytic activity of NO x -doped TiO 2 in the visible light region[END_REF] have discovered that N-doped TiO 2 exhibited photocatalytic activity under the visible light illumination. However, this result failed to attract attentions. Only in 2001, Asahi et al.[START_REF] Asahi | Visible-light photocatalysis in nitrogen-doped titanium oxides[END_REF] reported N-doped TiO 2 photocatalytic materials on 'Science' journal and the nonmetal ion doped TiO 2 became a hot research topic.

	(2) Nonmetal ion doping

Khan et al.

[START_REF] Khan | Efficient photochemical water splitting by a chemically modified n-TiO 2[END_REF]

prepared C-doped TiO 2 thin films by burning metal titanium sheet. The absorption edge of obtained TiO 2 film red shifted to 535 nm, thus achieving photoelectric conversion efficiency of 8.35% when electrolyzing water under visible light irradiation.

Table 1 -

 1 Sm 2 S 2 , Sm 2 O 3 and TiO 2 together, and then performed a calcination at 1000 o C under vacuum, obtaining the quaternary oxide semiconductor of Sm 2 Ti 2 S 2 O 5 with narrow band gap of 2.1 eV. Yamasita et al.[START_REF] Yamasita | Recent progress of visiblelight-driven heterogeneous photocatalysts for overall water splitting[END_REF] calcinated MTa 2 O 7 (M = Ca, Sr, Ba) in NH 3 gas at 850 o C and obtained the quaternary narrow band gap semiconductor of MTaO 2 N (M = Ca, Sr, Ba, corresponding to the band gap of 2.5 eV, 2.1 eV, 2.0 eV). In recent years, Zou and co-workers are very active in this research field. They prepared many series of multiple metal oxides, such as pyrochlore configuration compound of Bi 2 MNbO 7 (M=Al, Ga, In or rare earth element)[START_REF] Zou | Photocatalytic properties and electronic structure of a novel series of solid photocatalysts, Bi 2 RNbO 7 (R = Y, Rare Earth)[END_REF], wolframite configuration compound of InMO 4 (M=Nb, Ta, V)[START_REF] Zou | Photophysical and photocatalytic properties of InMO 4 (M = Nb 5+ , Ta 5+ ) under visible light irradiation[END_REF], perovskite configuration compound of MIn 0.5 Nb 0.5 O 3 and MCo 1/ 3 Nb 2/3 O 3 (M=Ca, Sr, Ba)[START_REF] Yin | Photophysical and photocatalytic properties of MIn 0.5 Nb 0.5 O 3 (M = Ca, Sr, and Ba)[END_REF]. Table1-1 shows a part of Zou's research on the preparation and properties of new narrow band gap semiconductors. It can be seen from the table that the prepared new narrow band gap semiconductors are ternary or multiple oxides, the band gaps are about 2.0-3.0 eV (620 nm to 410 nm). This absorption range of 410-620 nm covers most visible light.

	No 1	Semiconductors In 1-x Ni x TaO 4 (x=0-0.2) [160]	Band gap (eV) or Wavelength 2.3	Activity characterizations Decomposing water for H 2	Irradiation conditions λ>420 nm	Efficiency Quantum 0.66 %	yield
	2	InMO 4 (M=Nb, Ta, V) [155]	Nb: 2.5 Ta: 2.6 V: 2.0	Decomposing water for H 2	λ>420 nm irradiation for 50 h	Nb: 4.0μmol/(g h) Ta: 3.5μmol/(g h) V: 7.0μmol/(g h)
	3	MCo 1/3 Nb 2/3 O 3 (M=Ca, Ba)[161]	Sr,	Ca: 2.80 Sr: 2.46 Ba: 2.46	Decomposing methyl alcohol for H 2	λ>420 nm irradiation for ~13 h	Ca: 0.44μmol/h Sr: 0.17μmol/h Ba: 1.40μmol/h
	4 5 6	MIn 2 O 4 (M=Ca, Sr, Ba) [162] M 2.5 VMoO 8 (M=Mg, Zn) [163] Bi 2 MNbO 7 (M=Al, Ga, In) [164,165]	Maximum response wavelengh 480nm Maximum response wavelengh 530nm Al: 2.9 Ga:2.75 In: 2.7	Decomposing 15.3mg/L methylthion chloride Decomposing AgNO 3 solution of for O 2 Decomposing methyl alcohol for H 2	λ>420 nm irradiation for 2 h λ>420 nm irradiation for ~24 h irradiation for 10 h	Ca: 60 % Sr: 42 % Ba: 30 % Mg: ~3μmol/h Zn: ~6μmol/h Al: 0.71mmol/(g h) Ga: 0.3 mmol/(g h) In: 0.18 mmol/(g h)
	7	M 3 V 2 O 8 (M=Mg, Ni, Zn) [166]	Mg: 3.02 Ni: 2.25 Zn: 2.92	Decomposing AgNO 3 solution for O 2	λ>420 nm irradiation for 5 h	Quantum yield Mg: 0.23 % Ni: 0 % Zn: 0.64 %

1 The photocatalytic activities of various narrow band gap semiconductors under visible light

Table 2 -

 2 1. The diffraction peaks of the glass ceramic obtained from the base glass without CuI consists of two phases, GeSe 2 and Sb 2 Se 3 . than that of Sb 2 Se 3 although the initial amount of Ge is much higher than Sb. Indeed, GeSe 2 is glass former and tends to stay in amorphous state in the 80GeSe 2 -20Sb 2 Se 3 sample.When introducing CuI into the base glass, the obtained glass easily leads to the generation of Cu 2 GeSe 3 , resulting in the decrease of GeSe 2 . Moreover, the formation of Cu 2 GeSe 3 makes Sb amount exceed the stoichiometric ratio, according to the Sb:Se atomic ratio of 2:3 in Sb 2 Se 3 compound. The excessive Sb element will exist in the form of elemental Sb after further increase of the CuI content, and the phase of GeSe 2 disappears at the same time.

			80GeSe 2 -20Sb 2 Se 3
	Heat flow (a.u.)		72GeSe 2 -18Sb 2 Se 3 -10CuI		309	o C			393 o C 410 o C
			64GeSe 2 -16Sb 2 Se 3 -20CuI							390 o C
									278 o C				379	o C
		100				200			300				400
								Temperature (	o C)
								(a)										(b)
	Intensity (a.u.)						384 296 o C 1h o C 1h	Intensity (a.u.)								268 369 o C 1h o C 1h
	10	20	30	40	50	60	70	80	90	10	20	30	40	50	60	70	80	90
				(degree)								(degree)

As GeSe 2 is more difficult to crystallize

[28, 29]

, the diffraction peak intensity of GeSe 2 is lower

Table 2 -

 2 1 Preparation conditions, crystallized phases and photoelectric properties of different glass ceramics in Series 1 This series of samples is prepared by introducing various amount of CuI into the base glass of 70GeSe 2 -30Sb 2 Se 3 . Figure 2-21 shows the XRD patterns of different glass ceramics prepared with conditions listed in Table 2-2. The diffraction peaks of the glass ceramic obtained from base glass without CuI consists of GeSe 2 and Sb 2 Se 3 with possibly some selenium. When the introduction of CuI into the base glass, the crystallized phases of the glass ceramics are considerably modified, and the main crystalline phase is Cu 2 GeSe 3 , with a very small amount of Sb 2 Se 3 . When increasing CuI content up to 30 mol%, the intensity of the diffraction peaks attributed to Sb 2 Se 3 decreases obviously and a new crystalline phase of SbSeI appears due to the presence of a large number of iodine.

	Samples	T x ( o C)	Annealing conditions	Crystal phase	Photocurrent I p at -0.5 V (µA)	Semiconductor type	Conductivity (Ω -1 • cm -1 )
	80GeSe 2 -20Sb 2 Se	-	450 o C 5h	GeSe 2 + Sb 2 Se 3	-0.05	P	<10 -8
	72GeSe 2 -18Sb 2 Se 3 -10CuI	296 296 o C 1h Cu 2 GeSe 3 384 384 o C 1h GeSe 2 + Cu 2 GeSe 3 Sb 2 Se 3 +	0 -0.005	-P	<10 -8 <10 -8
	64GeSe 2 -16Sb 2 Se 3 -20CuI	268 268 o C 1h Cu 2 GeSe 3 369 369 o C 1h Cu 2 GeSe 3 Sb +Sb 2 Se 3 +	0 0	--	<10 -8 <10 -8
	4.2 Series 2: (1-x)(70GeSe 2 -30Sb 2 Se 3 )-100xCuI, x = 0, 0.1, 0.2, 0.3

Table 2

 2 The second crystal phase is Sb 2 Se 3 . With the increase of CuI content, Cu 2 GeSe 3 begins to appear in the sample. When further increasing CuI content from 20 mol% to 30 mol%, the excessive iodine leads to the formation of SbSeI.

	-2 that the 63GeSe 2 -27Sb 2 Se 3 -10CuI

CuI content (10 mol%), the main crystallized phase seems to be close to elemental Se with however slightly shifted diffraction angle. This is difficult to understand, as selenium is a good glass-former. More detailed studies are necessary to elucidate the mystery of the

  The results may be attributed to the poor conductivity of the glass ceramics (as shown in Table2-3).

	Cu 2 GeSe 3 SbSeI Sb 2 Se 3 42GeSe 2 -28Sb 2 Se 3 -30CuI 48GeSe 2 -32Sb 2 Se 3 -20CuI 54GeSe 2 -36Sb 2 Se 3 -10CuI 60GeSe 2 -40Sb 2 Se 3 series. Figure 2-24 Current-voltage characteristics of different glass ceramics in Series 3 under Intensity (a.u.) -0,6 -0,4 -0,2 0,0 0,2 0,4 0,6 -0,02 0,00 0,02 0,04 0,06 Current (µA) Potential applied (V)
			chopped visible light illumination	Se
	10	20	30	40	50	60	70	80	90
	Table 2-3 Preparation conditions, crystallized phases and photoelectric properties of  (degree) different glass ceramics in Series 3
	Samples		Annealing condition	Crystal phase		Photocurrent I p at -0.5 V (µA)	Type	Conductivity (Ω -1 • cm -1 )
	60GeSe 2 -40Sb 2 Se 3	404 o C 1h	Sb 2 Se 3 +Se	0	P		<10 -8
	54GeSe 2 -36Sb 2 Se 3 -10CuI	348 o C 1h	Sb 2 Se 3 +Se	0	P		8.4 × 10 -8
	48GeSe 2 -32Sb 2 Se 3 -20CuI	317 o C 1h	Cu 2 GeSe 3 + Sb 2 Se 3 +Se	-0.2	P		3.3 × 10 -7
	42GeSe 2 -28Sb 2 Se 3 -30CuI	303 o C 1h	Cu 2 GeSe 3 + SbSeI	-1.73	P		<10 -8

  When introducing CuI into the base glass, the obtained glass ceramics, after annealing with the conditions listed in Table2-4, contain three crystallized phases: Cu 2 GeSe 3 , Sb 2 Se 3 and something close to elemental selenium. With the increase of CuI content, the intensity of the diffraction peaks attributed to elemental Se decreases obviously and the Cu 2 GeSe 3 becomes more important. When further increasing CuI content to 30 mol%, the intensity of the diffraction peaks of Sb 2 Se 3 decreases significantly and a new crystalline phase of SbSeI appears due to the presence of a large number of Iodine.Figure 2-26 shows the dark and light current-voltage characteristics of the glass ceramics.The glass ceramics of 50GeSe 2 -50Sb 2 Se 3 and 45GeSe 2 -45Sb 2 Se 3 -10CuI cannot generate significant photocurrent. When the content of Cu 2 GeSe 3 is increased, the photocurrent is generated and reaches the highest value. For the glass ceramic of 35GeSe 2 -35Sb 2 Se 3 -30CuI, the content of Sb 2 Se 3 is decreased since part of antimony exists in the form of SbSeI. The photocurrent generated by this sample is comparable with that of the 40GeSe 2 -40Sb 2 Se 3 -20CuI.

							50GeSe 2 -50Sb 2 Se 3
							45GeSe 2 -45Sb 2 Se 3 -10CuI
							40GeSe 2 -40Sb 2 Se 3 -20CuI
	Intensity (a.u.)						Sb 2 Se 3 35GeSe 2 -35Sb 2 Se 3 -30CuI
							Cu 2 GeSe 3
							Se
							SbSeI
	10	20	30	40	50	60	70	80	90
					 (degree)

Table 2 -

 2 4 Preparation conditions, crystallized phases and photoelectric properties of different glass ceramics in Series 4 The last series of samples is prepared by introducing various amount of CuI into the base glass of 40GeSe 2 -60Sb 2 Se 3 . After different heat treatments detailed in Table 2-5, the XRD patterns of the obtained glass ceramics are shown in Figure 2-27. It can be seen that the base glass of 40GeSe 2 -60Sb 2 Se 3 is stable. It is difficult to attribute the diffraction peak around 21° to elemental selenium. Beside the fact that selenium is a good glass former, in case of crystallization of a large amount of selenium, the remaining composition would be highly instable against crystallization. The addition of CuI can obviously promote the crystallization in the base glass, leading to the formation of Sb 2 Se 3 . With higher CuI content, Cu 2 GeSe 3 appears in the crystallized phases of the glass ceramics. When the CuI content is 30 mol%, the formation of SbSeI crystals is observed.Figure 2-28 shows the dark and light current-voltage characteristics of the glass ceramics.The glass ceramic of 40GeSe 2 -60Sb 2 Se 3 cannot generate photocurrent due to its poor crystallinity and low conductivity (as shown in Table2-5). The photocurrent cannot be detected as well in the glass ceramic of 36GeSe 2 -54Sb 2 Se 3 -10CuI although it contains Sb 2 Se 3 crystals. With the increase of CuI content, the photocurrent is generated in the samples and its intensity first increases and then decreases. Meanwhile, the conductivity type of the samples in this series is changed from p-type to n-type when introducing CuI in the base glass (as shown in Table2-5).

	Samples	Annealing condition	Crystal phase	Photocurrent I p at -0.5 V (µA)	Type	Conductivity (Ω -1 • cm -1 )
	50GeSe 2 -50Sb 2 Se 3 353 o C 1h	Sb 2 Se 3	0	P	<10 -8
	45GeSe 2 -45Sb 2 Se 3 -10CuI	324 o C 1h	Cu 2 GeSe 3 +Sb 2 Se 3 + Se	0	P	1.8 × 10 -4
	40GeSe 2 -40Sb 2 Se 3 -20CuI	285 o C 1h Cu 2 GeSe 3 +Sb 2 Se 3	-28.4	N	4.4 × 10 -2
	35GeSe 2 -35Sb 2 Se 3 -30CuI	242 o C 1h	Cu 2 GeSe 3 +Sb 2 Se 3 + SbSeI	-23.2	N	5.0× 10 -3
	4.5 Series 5: (1-x)(40GeSe 2 -60Sb 2 Se 3 )-100xCuI, x = 0, 0.1, 0.2, 0.3

5 Conclusion

 5 To summarize, we have demonstrated the feasibility of interpenetrating heterojunction network in a purely inorganic composite material, by controlling the crystallization of a GeSe 2 -Ga 2 Se 3 -CuI glass. The network was formed by low conductive rod-like Sb 2 Se 3 crystals, covered by highly conductive Cu 2 GeSe 3 crystals. The Sb 2 Se 3 and Cu 2 GeSe 3 are efficient for charge photogeneration, the p-n junction formed by these two crystals greatly promotes the charge separation and the Cu 2 GeSe 3 plays the role of efficient charge collector. This structure results in high photocurrent (at least 100 times higher than any of these two crystal individually) and exceptionally long lifetime of charge carriers (around 16 µs). In this way, much more materials, such as Sb 2 Se 3 with many excellent features but low electrical conductivity can now be selected for photovoltaic application.

	1 Introduction
	Semiconductor photocatalysis technology is a new type of green technologies for
	efficiently exploiting the solar energy with double environmental benefits: organic wastes can
	be transformed into useful energies (electricity or chemical fuels) [1-4]. Such technology is
	based on the use of the photo-generated charge carriers from semiconductors. The free photo-
	generated electrons or holes can respectively react with the adsorbed O 2 or H 2 O on the
	semiconductor surface, leading to the formation of various active substances, such as •O 2	-
	, • HO 2 , H 2 O 2 ,
	Chapter 5
	Photocatalytic properties of the 40GeSe 2 -40Sb 2 Se 3 -20CuI
	glass ceramic

Table 5 -

 5 2 Single-ring aromatic intermediates of photodegraded MO azo[36][37][38][39][40]

	Molecular formula	Maximum absorption wavelength (nm)	Molecular formula	Maximum absorption wavelength (nm)
	 H 2 N		N   CH 3 CH 3	200	O 2 N			N   CH 3 CH 3	387
	 HO		N   CH 3 CH 3	251	 H 2 N		SO 3 Na	<200
	 HO		SO 3 Na	228	H 2 N			NO 2	381
	 O 2 N				NO 2	255	 HO		NO 2	318
										
	 HO			OH	223			 		243
										
	H 2 N					NH 2	243		N   CH 3 CH 3	200
				NH 2	235		SO 3 Na	264
				OH		211	201

Chapter 3

Optimization of the crystallization process of the 40GeSe 2 -40Sb 2 Se 3 -20CuI glass ceramic Moreover, looking for new low-cost and highly conductive alternative materials to fabricate a conductive channel may be another potential research direction of the glass ceramic.

(3) Optimizing the glass ceramic for photovoltaic application. In this thesis, we just presented a new glass ceramic with unique microstrcture as well as corresponding basic photoelectric properties. But how to use it for photovoltaic application is unclear, and other properties of this glass ceramic also need to be further investigated from the viewpoint of photovoltaic application. Moreover, the preparation of a glass-ceramic film for a potential application in photoelectric film devices is also a promising work, which is ongoing in our laboratory.

(4) Optimizing the glass ceramic for photocatalytic application. The glass ceramic can be used as a new photocatalyst due to its high photocatalytic activity illustrated by the good oxidative deamination ability. However, for its application, there is still many works to do.

Firstly, as a highly efficient photocatalyst, the particle size of the glass ceramic powders need to be further decreased to achieve high specific surface area. Secondly, considering the recovery of photocatalyst, loading technology of the glass ceramic powders, for example by preparing glass-ceramic coating, is certainly to be developed. Finally, the design of photocatalytic reactor based on the glass-ceramic photocatalyst is also an important work, especially for the design of gas-phase photocatalytic reactor considering the hydrophobic feature of the glass ceramic. Surface modification of the glass ceramic may be also necessary.

VU : VU :

Le Directeur de Thèse Le Responsable de l'École Doctorale

Abstract

A totally new family of glass ceramics with a unique microstructure was fabricated by controlling the crystallization of the GeSe 2 -Sb 2 Se 3 -CuI glass system. The influences of the material composition and the crystallizing process of the precursor glasses on the microstructure and photocurrent of the prepared glass ceramics were investigated. An optimized composition, 40GeSe 2 -40Sb 2 Se 3 -20CuI, was particularly studied with the following significant results:

(1) After a systematic study, it was found that this particular composition shows the highest photocurrent density among all studied glasses in the pseudo-ternary GeSe 2 -Sb 2 Se 3 -CuI system. It is also demonstrated that the photocurrent generated by different glass ceramics is not only determined by the composition, but also by the composite microstructure of the glass ceramic, which is determined by the ceramisation process. This process was then carefully studied. Compared with the two-step heat treatment process, the single-step process at a low temperature is a more efficient strategy to build up an efficient composite microstructure, which promotes charge carrier separation and provides a conductive channel, leading to a high photocurrent intensity in the glass ceramic.

( (3) The formation of nano-heterojunctions between Sb 2 Se 3 and Cu 2 GeSe 3 crystals within a single conductive microdomain can fully lead to an efficient separation of photo-generated electrons and holes. Therefore, for the photocatalytic application, it is unnecessary to form conductive channels (interconnected conductive microdomains) in the whole glass ceramic. Moreover, in order to form conductive channels, the necessary increase of annealing time or/and temperature may decrease the photocatalytic activity due to its relatively large crystal grain size. The optimized glass ceramic exhibits a good oxidative deamination ability and high photocatalytic activity, demonstrating its potential as an efficient photocatalyst.