Dr Sebastián Murcia-López

Dr Klaus Ellmer

Prof Roel Van De Krol

Prof Joan Ramón

Marc Birot

Marie-Anne Dourges

Odille Babot

Dr Joachim Brötz

Dr Eric Le

Keywords: Craquage de l'eau, anisotropie, BiVO 4, CuFeO 2, hétérostructures, spectroscopie water splitting, anisotropy, BiVO 4, CuFeO 2, heterostructures, photoelectron spectroscopy, interface experiment, photodeposition, band alignment

Titre: Analyse d'interface et développement des hétérostructures de BiVO 4 et CuFeO 2 pour le craquage photochimique de l'eau Résumé: Le craquage photo(électro)chimique (PEC) de l'eau par l'énergie solaire est considéré comme une méthode prometteuse de production renouvelable d'hydrogène. Dans ce travail, des hétérostructures à base de BiVO 4 et CuFeO 2 ont été choisis pour effectuer la réaction d'oxydation et de réduction de l'eau, respectivement. Cependant, les avantages exacts des hétérostructures n'ayant pas encore été complètement élucidés. Ce travail a eu pour objectif d'examiner les propriétés de certaines hétérojonctions à base de BiVO 4 et de CuFeO 2 par des expériences d'interface. Dans ce but, un certain matériau a été pulvérisé sur un substrat de BiVO 4 ou de CuFeO 2 et des mesures de spectroscopie de photoélectrons ont été effectuées à chaque étape du dépôt. Nous avons ainsi pu interpréter l'alignement des bandes entre le substrat et le matériau pulvérisé, et déterminer l'accordabilité du niveau de Fermi pour les absorbeurs étudiés. Par ailleurs, des hétérostructures à base de particules de CuFeO 2 et de BiVO 4 anisotropes sont été élaborées par photodéposition. Les performances de ces poudres dans des expériences de craquage photochimique de l'eau ont ensuite été déterminées.

Motivation

At the moment we are still living in a hydrocarbon based economy, in the sense that we are mostly dependent on coal, natural gas and oil for transportation, food production, electricity and heating. The reliance on fossil fuels, however, poses several problems. Fossil fuels are unevenly distributed over the globe, creating often geopolitic tensions and giving some states, such as the Gulf Arab states or Russia, tremendous power. Fossil fuel burning has also led to a plethora of environmental problems, being regarded as a main cause of climate change due to it being the largest source of greenhouse gas emissions and a major cause of air pollution due to the undesired release of nitrogen oxides, sulphur dioxide, volatile organic compounds and heavy metals. Potentially, these problems can be overcome by shifting from a fossil fuel based economy towards a hydrogen based economy with hydrogen as predominant energy carrier. Actually, approximately 65 Mt hydrogen is produced yearly, 1 for the production of ammonia (49% of global H 2 production), cracking of heavy petroleum derivatives into lighter fractions (37%), for methanol production (8%) and for a few niche uses (6%). 2 These niche uses include energy storage, fuel cell conversion, agent in speciality chemical production and agent in carbon based fuel production.

In a hydrogen based economy these niche uses would grow in importance. There are a few ways of producing hydrogen; some which are existing mature technologies: steam methane reforming (SMR) and coal gasification (96% of total hydrogen production), and water electrolysis (4%); while others are still at the lab scale such as photo(electro)chemical water splitting. Lab scale methods with reduced CO 2 emissions need to be further developed and adaptations to the existing technologies need to be implemented to realise a renewable hydrogen economy with a limited environmental impact. Figure 1.1 shows how a renewable hydrogen economy could look like. 3 The main reason why there is now such a strong discrepancy between the amount of hydrogen produced from water electrolysis and from fossil fuels is due to the extremely low cost to Figure 1.1: Scheme of a future energy infrastructure based on the hydrogen economy idea. Central in this scheme is the production of hydrogen through direct and indirect water splitting, using wind and sun as renewable and fossil-free energy sources. The hydrogen can then be used to produce carbon-, nitrogen-, and hydrogen-based chemical fuels through already existing large-scale infrastructure for methanation and Fischer-Tropsch syntheses or used in the Haber-Bosch process to produce ammonia. 3 produce hydrogen from fossil fuels. For instance, the current levelised cost (including capital and operating costs) of hydrogen production (LCH) through steam methane reforming is estimated at around $1.25 per kg H 2 based on the linear relationship between natural gas prices and hydrogen production cost with natural gas prices lying around $3.50 per GJ. 4 On the other hand, water electrolysis driven by the industrial coal-powered electricity price of $0.05 per kWh -1 has an approximate cost of $4.09 per kg H 2 , which was calculated through an H2A i type analysis. 5 Another study showed a levelised cost of hydrogen produced through grid electrolysis to be $5.5

per kg H 2 , assuming an electrolysis system efficiency of 61% and a grid electricity cost of $0.07 per kWh -1 . 6 Reducing greenhouse gas emissions is one argument for switching towards a hydrogen economy. However, hydrogen should then be produced in another way as both SMR and electrolysis based on electricity from coal power plants emit more CO 2 (Figure 1.2) than if fossil fuels were to be used directly. A few greenhouse gas emission reduction strategies have been suggested, which, however, all come with a cost compared to the base cost of SMR. Carbon capture and i The H2A (hydrogen analysis) project was established at the initiative of the Department of Energy of the United States and provides modelling tools which enable users to assess the cost of producing and delivering hydrogen. More information can be found on https://www.hydrogen.energy.gov/h2a_analysis.html Figure 1.2: CO 2 emissions for different hydrogen production technologies. 4 sequestration (CCS) has been frequently proposed as a ways to reduce the CO 2 emissions originating from the reforming of fossil fuel feedstocks. Alternatively, methane pyrolysis, whereby hydrocarbons are decomposed at high temperatures in the absence of oxygen, yield lower greenhouse gas emissions, due to carbon being transformed into solid carbon black. Hydrogen can eventually also be produced through biomass gasification, which, when coupled to CCS, potentially yield negative CO 2 emissions, due to CO 2 being fixated by photosynthesis during growing processes. A promising but immature hydrogen production technology involves the direct conversion of nuclear heat through thermochemical cycles, 7 whereby the sulphur-iodine and copper chloride cycles have been identified as being most performant. 8 Although, the direct emissions from water electrolysis are very low, the indirect emissions related to power generation and electrolyser systems have to be considered, whereby renewable (solar, wind, hydro, geothermal) and nuclear power generation yield the lowest CO 2 emissions (Figure 1.2). Recently, Parkinson et al. made a life cycle assessment of all the above listed hydrogen production techniques in their current state, comparing the cost of mitigation and the proportion of decarbonization (Figure 1.3). 9 Herein, the decarbonization of fossil fuel based hydrogen production methods was regarded as being most cost-effective with, however, only middling decarbonization levels. Also, methane pyrolysis links high cost-effectiveness to moderate decarbonization levels, which, unfortunately suffers from cost variability, due to the dependence on the carbon black price. The use of nuclear energy is promising due to the high cost-effectiveness and significant carbon emission reduc-tions, but its implementation could be complicated due to concerns with regards to proliferation and perceptions of safety together with the required high capital investment costs. Additionally, according to the authors, the studies regarding the conversion of nuclear heat through thermochemical cycles exhibited limited depth, transparency and reliability. Renewable water electrolysis also provides hydrogen with strongly reduced CO 2 emissions, which, can, however, not be implemented yet on a large scale due to the high production costs. Reducing the hydrogen production costs of electrolysis requires higher capacity factors, lower electrolyser costs and/or a low cost electricity supply. Figure 1.3: Proportional CO 2 emission reduction versus the percentual cost increase relative to steam methane reforming (SMR). LCCM corresponds to the levelised cost of carbon mitigation and is a measure for the extra cost related to emission reduction from hydrogen production per ton of CO 2 emitted compared to the emissions of hydrogen production through SMR. 9 To assess the economics of solar hydrogen production and to provide an incentive to focus on certain research goals Shaner et al. made an extensive technoeconomic analysis of photovoltaic-electrolytic (PV-E) and photoelectrochemical (PEC) water splitting. 10 Herein, basecase PV-E with an electrolyser efficiency of 61%, a photovoltaic efficiency of 16%, a photovoltaic/electrolyser capacity factor of 0.204 and a photovoltaic area of 0.75 km 2 with current balance of system (BoS), electrolyser and PV costs would yield an LCH of $12.1 per kg H 2 with the overall plant efficiency having the largest impact on the LCH. 10 Because of the high BoS costs, halving the PV module and electrolyser cost would only result in a reduced LCH of $10 per kg H 2 whereas doubling the plant efficiency by increasing the capacity or using more efficient photo-voltaics (PV) could reduce the LCH to $6 per kg H 2 (Figure 1.4a). The authors further specified the effect of the electrolyser and PV module cost on the LCH (Figure 1.4b). When the PV levelised cost of electricity (LCOE) would drop to about $0.02, which can be expected by 2050 in some regions in the EU, a drop in PV capital cost to about $0.50 per watt was calculated, assuming a PV lifetime of 25 years and a capacity factor of 30%. Moreover, by augmenting PV-E with wind energy, an electrolyser capacity up to 75% can be obtained, so that only an electrolyser capital cost reduction to $0.6 per watt is needed to obtain an LCH of $3 per kg H 2 (Figure 1.4b). Another option to reduce the LCH of PV-E is to allow the electrolysis installation to run on grid electricity during low solar power output, a technique which is called grid supplemented photovoltaic electrolysis (GSPV-E). Although, GSPV-E requires a higher BoS cost due to the implementation of additional power electronics such as AC-DC converters, a remarkably lower LCH, $6.1 per kg H 2 , was calculated with respect to PV-E, because of an increased electrolyser capacity of 97%. The combination of PV and/or wind facilities with electrolysers is an indirect way of water splitting. Alternatively, water splitting could be done directly through a process called photoelectrochemical (PEC) water splitting, which basically integrates light absorbers with electrocatalysts into one monolithic device, or through a particle suspension, which is designed for overall water splitting or for hydrogen/oxygen evolution using a redox mediator for charge transfer between particles. Pinaud et al. described how these direct water splitting processes could be integrated into large scale centralised reactors and calculated the LCH of each of these reactors through the H2A model by considering the current lab scale performance of the direct water splitting techniques, and the capital and operating costs of the reactors and auxiliaries. 11 The authors came up with 4 different realistic reactor types (Figure 1.5), making the distinction between particulate and photoelectrochemical water splitting. They calculated that a simple PEC array (type 3 reac-Figure 1.5: Four different direct water splitting reactor types: (a) baggies containing a particle slurry which directly converts sunlight and water into hydrogen and oxygen, (b) an alternating series of two different baggies each containing a particle slurry which is responsible for oxygen or hydrogen evolution, (c) A PEC panel containing several photoactive layers sandwiched between a cathode for hydrogen evolution and an anode for oxygen evolution, (d) a linear PEC panel in combination with a parabolic solar concentrator. 11 tor) operating at a solar-to-hydrogen conversion efficiency (STH) of 10% would have an LCH of $10.40 per kg H 2 , while supplementing the PEC array with a solar concentrator (type 4 reactor) now operating at 15% STH would lower the cost to $4.10 per kg H 2 due to the tenfold increase in solar concentration. Meanwhile, Shaner et al. calculated an LCH of $11.4 per kg H 2 for a type 3 PEC system (15% STH) and an LCH of $9.2 per kg H 2 for a type 4 PEC cell working at an STH of 20%. 10 Type 3 and type 4 PEC reactors have only a slightly lower estimated LCH than what was estimated for PV-E (see above), assuming a 10% STH for both cases. The lower LCH for PEC mainly stems from the approximate 2x lower BoS costs of a PEC device due to a more cost efficient energy transmission and because they have only 1 set of hard BoS costs compared to the two sets of BoS costs of a photovoltaics-electrolyser combination. However, the base case PEC assumptions, an STH of 10% and 10 year lifetime, have not yet been demonstrated on a lab scale, whereas PV panels and electrolysers, which together can operate at the same base assumptions, can already be produced on an industrial scale. 6,12,13 Thus, PV-E is a technology which is much more mature with proven electrolyser and photovoltaics industries. PV-E has the advantage that the systems for light absorption and water electrolysis can be separately optimized, which is less complicated than to design one device which can do both in a stable, safe and efficient way. For instance, the light absorber in PEC devices has to be stable in an electrolyte during operation for a prolonged period of time, the amount of electrocatalysts need to be limited to avoid parasitic light absorption and the extra interfaces need to be optimised for charge extraction without generating surface state induced charge carrier recombination. Therefore, critics assume that PEC devices which are cosmetically similar (type 3 and type 4) to PV-E systems will not be able to compete with PV-E on a mid to long term. 14,15 More so, because advancements in PEC design which originate from electrolyser or PV technology such as buried junctions or usage of earthabundant electrocatalysts would not give PEC systems a competitive edge over PV-E, because

PV-E would also benefit from these advancements. 16,17 Due to the higher simplicity, and thus a lower BoS cost, of particulate based water splitting reactors lower LCHs could potentially be reached. In the technoeconomic analysis of Pinaud et al. an LCH of $1.60 per kg H 2 was calculated for a single bed particle suspension operating at an STH of 10% and an LCH of $3.20 per kg H 2 was calculated for a dual bed particle slurry with an STH of 5%, thus indicating that such reactor designs (Figure 1.5) could be competitive to PV-E water splitting or even steam methane reforming. 11 In the calculations a strong sensitivity of the LCH on the efficiency was established, so that research efforts into particulate based water splitting should focus on improving the efficiency. However, on a lab scale, the state of the art particulate based water splitting system could only reach an STH of around 1% 18,19 and for a real device, a flat panel reactor of 1 m 2 operating under direct sunlight, an even lower efficiency, STH of 0.4%, was measured. 20 Possible technological limitations for type 1 and type 2 PEC reactor types are the safe operation of a single bed particle suspension system due to mixing of O 2 and H 2 and the long distances redox shuttles need to travel between the two different particles in a type 2 reactor, respectively.

All in all, renewable ways of water splitting could be viable in the foreseeable future if further efficiency improvements, capacity increases and cost reductions can be realised in PV and electrolysers, a breakthrough in PEC or particulate water splitting is achieved or a large enough tax on CO 2 emissions is implemented. The strong reliance of the LCH on the efficiency of PV-E, PEC and particulate based water splitting is an invitation to further the applied and fundamental research into these aspiring technologies.

Thesis overview and structure

As outlined above, particulate water splitting has the potential to become a cost competitive and green hydrogen production alternative to steam methane reforming, if the above listed base conditions, 10% STH + 5 years lifetime for Type 1 and 5% STH + 5 years lifetime for Type 2, 11 are fulfilled. Therefore, we opted to prepare new heterostructured particulate powders for water splitting as well as to improve the understanding of the junction properties of the solid state interfaces inherent to heterostructures. The ultimate goal was to come up with a working type 2 particulate water splitting system. Hence, we selected two semiconducting oxides, one which is appropriate for water oxidation, BiVO 4 , 21 and one which should be suitable for water reduction, CuFeO 2 . 22 These oxides have, however, a few shortcomings, such as charge carrier recombination and unfavourable water redox kinetics, which prevent them from reaching their theoretically predicted water splitting capacities. [22][23][24][25] A recognized solution is the creation of heterostructures by attaching additional contact materials onto the main photoabsorber. 26 However, in most cases the precise effects of those heterostructures are not fully understood yet, as many characteristics of the main photoabsorber, including surface passivation, 27 charge carrier separation 28 and surface redox kinetics, 25 can be altered through the incorporation of an additional contact material. Special mention goes to the remarkable finding that for {010}/{011} dual facet exposed BiVO 4 truncated bipyramidal shaped particles, anisotropic charge carrier separation and accumulation takes place. 29 Due to the anisotropic charge accumulation, anisotropic heterostructures, with an oxidation cocatalyst on one facet and a reduction cocatalyst on the other, could be created through facet-selective photoreduction and photo-oxidation reactions. 30 Dual cocatalyst coated BiVO 4 heterostructures demonstrated an up to two orders of magnitude increased water oxidation performance compared to bare {010}/{110} dual facet exposed BiVO 4 particles, due to a simultaneous optimization of charge carrier separation and synergistic surface redox kinetics. 25 Given this landmark finding, we wanted to develop novel anisotropic heterostructures based on anisotropically shaped BiVO 4 and CuFeO 2 . Additionally, due to the current open questions regarding the effects of heterostructures, we analysed the interfacial characteristics of selected BiVO 4 and CuFeO 2 heterostructures through in situ photoelectron spectroscopy and discussed how the results could improve the understanding of oxide heterostructure interfaces with respect to charge carrier separation and interfacial electronic structure.

Chapter 2 gives an introduction to the broad topic of photocatalysis, photochemical water

THESIS OVERVIEW AND STRUCTURE

splitting and interface analysis. The emphasis is laid on fundamental principles and on the state of the art heterostructured particulate water splitting systems. Furthermore, the specific evolution of BiVO 4 and CuFeO 2 based systems is detailed.

Chapter 3 encloses the experimental details of all synthesis and analysis techniques used during this work. Where relevant a short introduction to the used apparatuses is given. Also, all recipes are denoted in detail.

In Chapter 4 the results regarding the development of anisotropic BiVO 4 heterostructures will be given and discussed. First, the upscaling of the {010}/{110} dual facet exposed BiVO 4 powder synthesis is detailed. Afterwards, the preparation of heterostructured BiVO 4 /NiO x OH y , BiVO 4 /CoO x OH y , BiVO 4 /Ag and BiVO 4 /Pt through photodeposition is reviewed and the regioselectivity of the deposits is described through electron microscopy. Subsequently, the dye degradation and water oxidation performance of the heterostructured BiVO 4 powders are discussed.

Chapter 5 presents an overview of the BiVO 4 based interface experiments, carried out through a combination of thin film sputtering and photoelectron spectroscopy. This Chapter is opened with an overview of the surface and bulk characteristics of the two BiVO 4 thin film substrates, which were used to perform the interface experiments, denoted as gradient BiVO 4 and PLD The techniques and knowledge acquired in Chapter 4 and 5 were used to develop and study novel CuFeO 2 heterostructures, which is the topic of Chapter 6. The Chapter starts with a description of the upscaling of the synthesis of 2H-CuFeO 2 hexagonal nanoplatelets, which were used as the anisotropic substrates during the photodeposition experiments. Specifically, Pt, Ag and NiO x OH y were deposited, whereby the regioselectivity of the deposits was examined through electron microscopy. Thereafter, the sacrificial water reduction efficiency of the heterostructures was tested. Finally, interface experiments were performed with CuFeO 2 pellets, prepared by pressing CuFeO 2 powder, as substrates, and ITO and water as contacts.

( Chapter 2

Introduction

The fundamental concepts presented in this chapter have been deduced from the following main references, according to topic: General solid state physics, 1 semiconductor device physics, 2,3 semiconductor physics and interfaces, 4 photocatalysis, 5 photoelectrochemical water splitting 6 and photoelectron spectroscopy. 7 2.1 Photo(electro)chemical water splitting

Fundamentals of semiconductors

Photo(electro)chemical water splitting is a photocatalytic process, wherein semiconductors utilize the energy obtained from light absorption to split water into hydrogen and oxygen. Although, molecular water splitting exists as well and hybrid molecular/inorganic water splitting is on the rise, photo(electro)chemical water splitting in this work refers exclusively to the process performed with inorganic semiconductors. Semiconductors distinguish themselves from conductors and insulators through their intermediate electrical conductivity, for which an explanation can be found in the energy band model. The energy band model can be seen as an extension to the LCAO (linear combination of atomic orbitals) model, whereby molecular orbitals of a molecule are considered to be the result of linear combinations of the atomic orbitals of the atoms constituting the molecule. Semiconductors have, just like many other solid state materials, a crystalline structure, which is built-up through a repetitive three-dimensional translation of unit cells. The unit cells are the smallest grouping of atoms, which constitute the repeating pattern. In contrast to a molecule, which in most cases only consist of a few to a few 100s of atoms, macromolecules not included, a crystalline structure contains around 10 22 unit cells per cm 3 . Now, when such a large number of atomic orbitals overlap, the resulting "molecular" orbitals of similar nature (e.g. combination of 1s atomic orbitals) will be so closely spaced in energy that they are considered to form a continuum of electronic states, also called an energy band (e.g. the 1s atomic orbitals form a 1s band). For semiconductors, a series of bands will be completely filled with electrons, whereby the band containing the electrons with the highest energies is denoted as the valence band (VB), i and the empty band closest in energy to the valence band is considered to be the conduction band (CB). In between those bands there is a "gap", generally denoted as band gap, where no electronic states exist. Band gaps vary in the range of 0.1 eV -4.0 eV, thus in the UV-Vis-NIR (ultraviolet-visible-near infrared) spectral region. Electrical conductivity depends on the number of mobile charge carriers and to be considered mobile, charge carriers should have access to empty states, which are close in energy. Due to semiconductors having a band gap, which do not contain any electronic states, a considerable amount of energy, be it thermal or light, is required to allow the electrons to cross the band gap and reach the conduction band where the electrons have access to empty states. The electrons that are excited to the conduction band leave behind empty states or "holes" in the valence band, which are also considered to be mobile charge carriers, as electrons at the top of the valence band can move into these empty states. In the absence of light, the probability that a state of energy E is filled with an electron, f(E), is governed by the Fermi-Dirac distribution, which takes into account the Pauli exclusion principle which states that two or more identical fermions, electrons being fermions, cannot occupy the same electronic state:

f(E) = 1 1 + e (E-µ)/k B T (2.1)
where k B T is the product of the temperature and the Boltzmann constant, and µ stands for the total chemical potential of the electrons, which in semiconductor physics is also commonly i Within this thesis the denotion of the band energies is in relation to the band position in the used energy diagrams: higher (lower) lying bands have a lower (higher) distance to the vacuum level with increased (reduced) electron energy.

denoted as the Fermi level (E F ). The mobile charge carrier concentration, for electrons n and for holes p, can then be obtained by multiplying the Fermi-Dirac distribution with the density of states D(E). Figure 2.2 shows the Fermi-Dirac distribution, the density of states distribution, and the resulting hole and electron charge carrier concentration for an intrinsic, an n-type and a p-type semiconductor. n-type and p-type semiconductors distinguish themselves from intrinsic semiconductors through their higher electron and hole charge carrier concentration, respectively. A higher charge carrier concentration may be caused by intrinsic defects or due to the presence of impurities. A distinction is made between acceptor and donor impurities which may increase the electron and hole charge carrier concentration, relatively, as these impurities, associated with inter-gap states, tend to ionize and introduce additional charge carriers into the bands. The controlled incorporation of impurities is called doping. In this way, p-type silicon can be obtained by doping with a lower valence element (acceptor), such as boron, and n-type silicon can be obtained by doping with a higher valence element (donor), such as phosphorus.

The distinction between direct and indirect semiconductors is another semiconductor property that has to be considered with regards to photocatalysis (Figure 2.3). For all interactions between elementary particles, including electrons, holes, photons and phonons (lattice vibrations) the requirement of conservation of energy and conservation of crystal momentum (k-vector) ii has to be fulfilled. Now, semiconductors are called direct when their valence band maximum (VBM) and conduction band minimum are associated with the same crystal momentum, as is the case for some III-V semiconductors like GaAs and InP. When the crystal momentum is different, a semiconductor is considered to be indirect, which is the case, for instance, for Ge and Si.

A process which one tries to avoid in photocatalysis is charge carrier recombination, whereby light-induced holes and electrons annihilate each other with the release of energy in the form of phonons or photons. Recombination is in direct competition with photocatalysis, due to the electrons and holes not being involved in chemical redox reactions; the actual goal of photocatalysis. In direct semiconductors radiative recombination with the release of photons can easily take place, because the electron-hole annihilation preserves the crystal momentum and photons with energies of a few eV have almost zero momentum. In contrast, radiative recombination in indirect semiconductors necessitate the emission or absorption of a phonon to compensate for the momentum difference between the hole and the electron. The indirect photon-phonon recombination process in indirect semiconductors has a lower probability than the direct radiative recombination in direct semiconductors. Therefore, indirect semiconductors are more suitable for photocatalysis.

In this Section, the energy band model concept has only been described briefly. If one would ii The crystal momentum P crystal is defined as P crystal ≡ k with being the reduced Planck constant and k the wave vector. The wave vector is a consequence of Bloch's theorem, wherein the motion of electrons in a crystal lattice are considered to be subjected to an infinite periodic potential. According to Bloch's theorem, the single particle wavefunction (ψ n ) of an electron in a crystalline lattice can be written as: ψ n = e ikx u nk (x) , u nk (x + a) = u nk (x). Hence, the single particle wave function finds its stationary state solutions in terms of a periodic function u(x) multiplied by a plane wave, whereby the wave vector k can be seen as a representation of the stationary state solutions. like to have a more complete derivation of the theory and to gain a better understanding on semiconductor properties, I would recommend the reader to read through "Band Theory and Electronic Properties of Solids" by John Singleton. 9 In addition, readers are invited to read "Physics and Technology of Semiconductor Devices" by Andy Grove 3 or "Physics of Semiconductor Devices" by Sze 2 to learn how semiconductor physics play a role in the design of semiconductorbased devices.

Photocatalysis

As defined by the International Union of Pure and Applied Chemistry (IUPAC), photocatalysis is the initiation or the change in rate of a chemical reaction under the action of infrared, visible or ultraviolet radiation in the presence of a substance, the photocatalyst, which absorbs the light and is involved in the chemical transformation of the reaction partners. 10 According to this definition, a semiconductor based photocatalytic process follows three steps (Figure 2. whereby the diffusion coefficient according to the Einstein relation can be written as:

D = µ. k B T e (2.3)
with e being the elementary charge and µ being the mobility of the charge carriers. Charge carrier mobility is defined as the proportional constant which relates the velocity of a charge carrier (v D ) moving through a particular semiconductor under the influence of an electric field

E: v D = µ.E (2.4)
In a simple model, where charge carriers are considered to be delocalized, i.e. not attached to individual atoms, mobility depends mostly on the scattering time τ sc , which is the average time between scattering events. A scattering event namely negatively impacts the drift velocity, as the acceleration of the charge carrier by the electric field is stopped and the energy and/or direction of the charge carrier is changed. Several bulk scattering events have been identified, including ionized impurity scattering, defect scattering, phonon scattering, inelastic scattering, electron-electron scattering and some other minor scattering events. The exact relation between µ and τ sc , for delocalized electrons, becomes:

µ = e.τ sc m * (2.5)
with m * being the effective mass of the charge carriers. The notion of an effective mass is a repercussion of the above-discussed band theory and originates from the different motion of charge carriers in a periodic potential, as found in crystalline solids, with respect to the motion of an electron in a vacuum. The effective mass is material dependent and describes the mass a charge carrier seems to have in response to a certain force. The effective mass of electrons m * e and holes m * h differs in most cases from each other and is related to the density of states near the conduction band minimum N C and valence band maximum N V , relatively. For an ideal semiconductor with parabolic valence and conduction band structures near the band edges:

N C = 2 ( 2π m * e k B T h 2 ) 3 2 (2.6) N V = 2 ( 2π m * h k B T h 2 ) 3 2 (2.7)
The assumption of a parabolic band structure is valid for many semiconductor valence bands, such as for Ge, Si and GaAs, and for some semiconductor conduction bands, such as for GaAs.

For asymmetric bands, such as the conduction band of Si, the effective mass also depends slightly on temperature. For the semiconductors analysed in this work, BiVO 4 and CuFeO 2 , charge carrier mobility is slightly more complicated as electrons and holes in these semiconductors are assumed to form polarons, which have considerably higher effective masses, as will be explained further in Section 2.3. In summary, the diffusion length depends on the charge carrier lifetime and the charge carrier mobility, whereby mobility is governed by scattering events and the effective masses of the charge carriers.

For indirect semiconductors, which contain bulk defects, charge carrier lifetime and mobility are strongly dependent on the number of defects. These defects can, namely, act as recombination centres and scattering centres, which reduce charge carrier lifetime and charge carrier mobility, respectively. Recombination at bulk defects, which act as deep level traps, whereby the energy is exchanged in the form of lattice vibration is also called Shockley-Read-Hall recombination. 11,12 Bulk defects intrinsic to a certain material cannot be overcome. However, structure related bulk defects such as dislocations, point defects and grain boundaries can be reduced by synthesizing the semiconductor under the "right" conditions. For instance, monocrystalline silicon has a strongly reduced number of defects compared to polycrystalline silicon, which is why monocrystalline silicon photovoltaics generally outperform polycrystalline silicon photovoltaics.

Besides trying to optimize the charge carrier diffusion length, bulk recombination can be prevented through more efficient charge carrier separation. Charge carrier separation can be realised through the introduction of an internal electric field and/or the implementation of charge selective contacts. Herein, interfaces play an important role. How semiconductor surfaces and interfaces influence charge carrier separation will be explained in greater detail in Section 2.2, wherein the physics and effects of semiconductor surfaces and interfaces are described. Another way to reduce bulk recombination is through nanostructuring, because then the distance between the point of light-induced electron-hole generation and the surface is reduced, which especially benefits semiconductors with limited diffusion lengths. Furthermore, nanostructuring leads to a decrease in the bulk-to-surface ratio, which has the advantage that the number of potential surface reaction sites is increased, but which may also increase the number of surface recombination centres. Additionally, light absorption may be negatively impacted, because semiconductor UV-Vis absorption lengths are on the scale of 100 -1000 nm.

Although, intergap states are, just like for bulk recombination, the main cause for surface recombination, their origin in most cases is different and is related to the surface/interface structure of the semiconductor. Surface states are inherent to crystalline structures and originate from the sudden disruption of the crystalline lattice, with atoms at the surface having fewer bonding partners. Hence, some electrons at the surface do not engage in bonding, but form so-called dangling bonds, whereby the electrons are unpaired and reside in electronic energy levels inside the normally forbidden band gap. Strategies to reduce surface recombination are commonly referred to as surface passivation strategies and mainly try to reduce the number of dangling bonds or attempt to reduce the concentration of a particular carrier at the surface.

Instead of participating in the desired redox reactions or recombining at surface states, charge carriers arriving at the surface can also engage in corrosion processes, leading to a gradual dissolution of the semiconductor. In this case, the semiconductor should be coated by a protection layer, which shield the semiconductor surface from the electrolyte. Applying a conformal protection layer is relatively easily done for thin films, but can be cumbersome for particulate based systems.

All in all, photocatalytic systems can be plagued by a series of problems, which all necessitate tailor-made solutions depending on the semiconductor bulk and surface properties, and the desired application. At the moment, the optimization and implementation of photocatalysts is mainly studied with regards to the decomposition of air and water borne substances and photo(electro)chemical water splitting.

Fundamentals of photo(electro)chemical water splitting

Basically, water splitting processes focus on the dissociation of water into hydrogen and oxygen.

The overall reaction is as follows: and can also be written in terms of two standard reduction reactions: of 0 V and 1.23 V relative to the standard hydrogen electrode (SHE), respectively. Hence, the total cell potential (E 0 ) amounts to -1.23 V and is associated with a strongly positive standard Gibbs free energy according to:

2 H 2 O 2 H 2 + O 2
∆G 0 = -nFE 0 (2.8)
with n being the number of electrons involved in the overall reaction and F being the Faraday constant. Herein, ∆G 0 amounts to 237.13 kJ mol -1 , meaning that the overall water splitting reaction is thermodynamically uphill, thus not spontaneous, for which a minimal potential difference of 1.23 V needs to be supplied. This potential difference can be supplied by an external bias and/or through the absorption of light, which leads to the concept of photo(electro)chemical water splitting.

According to the aforementioned IUPAC definition of photocatalysis, photo(electro)chemical water splitting can be understood as a photocatalytic process, because light allows the initiation of the energetically uphill water splitting reaction in the presence of an appropriate photocatalyst. Therefore, authors also commonly refer to photocatalytic water splitting instead of photo(electro)chemical water splitting, which is used primarily in this work.

To overcome the 1.23 V potential difference with light alone, a semiconductor has to be found with a band gap of at least 1.23 eV and with band positions that straddle the H + /H 2 and O 2 /H 2 O reduction potentials. Figure 2.6 shows which binary semiconductors could be suitable for overall water splitting according to their band gap and band positions. 13 Herein, the band gap requirement of 2.2 eV is more strict than the minimal 1.23 eV to accommodate the need of overpotentials, due to activation barriers, shunt losses and resistances. As can be observed, not many binary semiconductors are suitable for overall water splitting and the ones that are, are normally wide-band gap semiconductors such as, TiO 2 , ZnO and ZnS. Because their band gaps are above 3 eV, wide-band gap semiconductors are only able to absorb UV radiation, which makes them rather inefficient for solar water splitting, due to the small fraction of UV light in the solar spectrum (Figure 2.7). The main reason why mostly only wide-band gap oxides are suitable for water splitting is because of their valence band maxima, which lie in most cases far below the water oxidation potential (see Section 2.2.1.3 for a description on band edge position determination).

The relatively low VBMs of oxides are mainly due to the low lying O2p states which make up most of the oxide valence band maxima. Binary II-VI and III-V semiconductors have generally higher valence band maxima, but they are often highly susceptible to photocorrosion. [14][15][16] Due to the unsuitability of binary semiconductors for overall solar photochemical water split-Figure 2.6: Scheme of binary semiconductor band positions according to their band gaps. The blue zone indicates the band positions and/or band gaps that are unsuitable for overall water splitting. 13 ting, one could select suitable ternary and quaternary semiconductors (Section 2.3) and/or select multiple semiconductors with complementary band positions and/or supply electric energy to achieve enough voltage to realize the water splitting reactions. The strong advantage of using a combination of photoabsorbers is that semiconductors with band gaps below 2.2 eV become useful, meaning that a larger fraction of sunlight may be captured. Figure 2.8 shows a scheme of how two semiconductors could be combined to achieve overall photo(electro)chemical water splitting. Herein, one semiconductor should be an n-type semiconductor with a relatively wide band gap and a VBM suitable for water oxidation which preferably exhibits upwards band bending, while the other semiconductor should be a p-type semiconductor with a more narrow band gap and a conduction band minimum (CBM) suitable for water reduction which preferably exhibits downwards band bending. The two semiconductors should be able to shuttle charge car-Figure 2.7: Solar radiation spectrum at sea level and at the top of the atmosphere. 17 riers through an external circuit, with which the semiconductors should form ohmic contacts, or through a redox mediator, which is often a dissolved compound with multiple valence states such as iron. 18 Essentially, two semiconductors can be combined according to two different geometries: 1) a stack can be formed by placing one semiconductor on top of the other, so that the top semiconductor is irradiated by the full solar spectrum, while the bottom photoabsorber only receives the photons with energies below the band gap of the top photoabsorber or 2) the two light absorbers can be placed "side-by-side", so that they both receive the full solar spectrum, which is, for instance, the only possibility in particulate photochemical water splitting. Pinaud et al. calculated, for both geometries, the maximally attainable solar-to-hydrogen (STH) efficiency for a set of band gaps, taking into account entropic losses and reaction overpotentials (Figure 2.9). 19 For a dual side-by-side geometry, a maximum STH efficiency of merely 15.5% could potentially be reached through the combination of two semiconductors with an identical band gap of 1.59 The external circuit connecting the two semiconductors could also be replaced by a redox mediator to shuttle charges between the two semiconductors. Adapted from [19].

Engineering the band gap of oxide semiconductors is more complicated than it is for III-V semiconductors, because of the higher compositional freedom of III-V semiconductors, and because of the O2p states giving rise to relatively low VBMs. Hence, finding an oxide semiconductor able to reduce water with a band gap below 3.0 eV is not straightforward (Figure 2.6). However, for certain oxides the VBM is not only governed by the O2p states but also by the valence states of other elements, as is the case for bismuth, copper and iron containing oxides (Figure 2.10). 22 These kind of oxides are, thus, promising for (partial) photo(electro)chemical water splitting.

Additionally, higher VBMs can be obtained with oxide semiconductors through the addition of nitrogen, forming oxynitrides. Oxynitride compounds pose, however, new challenges such as a lower stability against photocorrosion. 23 To learn more about oxynitrides I refer the reader to a recently published review by Takata et al. 24 In this work, BiVO 4 and CuFeO 2 were chosen for the photochemical water oxidation and water reduction reaction, respectively (Figure 2.11). Although, BiVO 4 may not have the optimal optical band gap, 2.4 eV, 25 it is at the moment the state of the art oxide photoanode having demonstrated an STH efficiency of 4.9% in combination with a double junction amorphous Si solar cell 26 or having shown a photocurrent of 6.72 mA cm -2 vs 1.23 V RHE when combined with WO 3 into a core-shell nanostructured photoanode. 27 CuFeO 2 , in contrast, has a band gap, 1.5 eV, 28 which is very close to the optimal band gap and has suitable band positions for water reduction, however its efficiencies so far have been rather poor. The BiVO 4 and CuFeO 2 state of For the tandem configuration the top photoabsorber is illuminated by the whole solar spectrum, whereas the bottom photoabsorber only receives the photons transmitted by the top photoabsorber. For the side-by-side configuration both photoabsorbers are illuminated by the complete solar spectrum. 19 Figure 2.10: Band gaps and band positions of some promising oxide semiconductors for photo(electro)chemical water splitting, whereby the contribution of the metal, oxygen and nitrogen atomic orbitals to the valence band and conduction band are indicated. 22 the art is further described in this Chapter in Section 2.3.

On a side note, several efficiency parameters have been used throughout literature to assess photo(electro)chemical systems, whereby the STH efficiency is considered to be the central benchmark. I refer the reader to the article of Qureshi et al. for a complete overview of the efficiency parameters and how to appropriately measure those efficiencies, 29 because efficiency parameters other than the STH efficiency will be mentioned occasionally throughout this work. V vs NHE (pH=0)

Particulate water splitting

As mentioned in the Preface, four types of photo(electro)chemical water splitting reactors can be distinguished (Figure 1.5), whereby type 3 and type 4 are based on PEC stacks and type 1 and type 2 on baggies containing particle slurries. Due to the high BoS costs and the high resemblance to PV-E, questions arise whether PEC stacks will be able to compete with PV-E and SMR in the near to far future. 19,30 In contrast, type 1 and type 2 particulate water splitting systems could be economically viable, with computed levelised hydrogen production costs of $1.60 and $3.20 per kg H 2 , respectively, if the baseline conditions are fulfilled, i.e. 5% STH efficiency + 5 years lifetime and 10% STH efficiency + 5 years lifetime, respectively. 19 Nonetheless, is the research into PEC water splitting ongoing, because the knowledge gained on PEC systems can be used for the development of photovoltaics and/or particulate water splitting systems. Therefore, literature references to photo(electro)chemical systems have been mainly included in this work to highlight certain structural modifications which improved upon the water splitting performance or to provide insight into certain photo(electro)chemical water splitting mechanics.

Particulate water splitting can be realised in the absence or in the presence of a redox mediator, which transfers the electrons from the water reduction photoabsorber to the water oxidation photoabsorber. In the presence of a redox mediator the particulate water splitting system is referred to as a Z-scheme type system (Figure 2.12), which name is derived from the "Z-scheme" natural photosynthesis process in green plants. 18 The redox mediator can be an aqueous redox mediator such as Fe 2+ /Fe 3+ , 31 IO 3 -/I -32 and I 3-/I -33 or a solid state redox mediator like reduced graphene oxide 34 or gold. 35 The redox mediator can also be a conductive film onto which the particulate photoabsorbers have been immobilized. Remarkably, these kind of integrated particulate systems have demonstrated state of the art particulate-based unassisted lab-scale STH efficiencies ranging from 1.0 to 1.2%. 36,37 At a larger scale, a state of the art particulate-based STH efficiency of 0.4% was recently achieved under direct sunlight with a 1 m 2 -sized photocatalytic water splitting panel, consisting of immobilized Al-doped SrTiO 3 /RhCrO x particles. 38 Figure 2.12: Schematic of a Z-scheme type photochemical water splitting system. 18 In the following Section, fundamental semiconductor surface and interface properties will be described as well as their relevance to photo(electro)chemical water splitting.

Semiconductor interfaces

Interfacial band structure

Solid state junctions

As was discussed in Subsection 2.1.2, heterostructures, whereby the photoabsorber is contacted with another material, may help to overcome certain intrinsic semiconductor deficiencies. Besides the chemical and electronic properties of the contact material, the resulting interface structure may play an important role in the overall device performance. A first junction which one may come across in photo(electro)chemical devices are semiconductor/metal junctions, also called Schottky barriers. For instance, platinum has been added occasionally to the main photoabsorber to act as water reduction "cocatalyst". 39,40 The alignment between two solid state compounds is often represented through an energy band diagram, whereby certain energy levels of the compounds are plotted as a function of a spatial dimension using the Fermi level or vacuum level as energy reference. The vacuum level, herein, is defined as the energy level of a free stationary electron outside the material but near the surface. In ductor is proportional to the difference between the semiconductor electron affinity and metal work function. However, for many semiconductors the Mott-Schottky rule is not applicable, as the created interface gives rise to surface states within the band gap, which reduces the expected amount of charge transfer, a phenomenon called Fermi level pinning. These surface states are often called metal-induced gap states due to tails of the metal wave functions spilling into the semiconductor, creating inter-gap states. 41,42 Inter-gap surface states may also arise from dangling bonds and from structural defects at the semiconductor surface. 43 Fermi level pinning may even be caused through bulk semiconductor defects in a process called "self-compensation". 44 The amount of Fermi level pinning at the semiconductor/metal contact has also been found to depend largely on the chemical nature of the semiconductor, whereby covalent semiconductors such as Si have shown strong Fermi level pining, whereas ionic semiconductors such as oxides have generally demonstrated weak Fermi level pinning. 45,46 vac The transfer of mobile charge carriers between semiconductor and metal may induce a space charge region in the semiconductor, due to inefficient screening of the excess charge in the semiconductor near the semiconductor/metal interface. The inefficient screening of the excess charge may lead to a decrease, increase or inversion, of the number of native mobile charge carriers near the semiconductor/metal interface at the semiconductor side, so that the space charge region can be further specified as a depletion, accumulation or inversion region (Figure 2.14). For instance, a depletion region in an n-type semiconductor has a lower electron density near the semiconductor/metal interface compared to the bulk of the semiconductor. A depletion region will, thus, exhibit insulating behaviour. Through the application of an external bias, one can switch between depletion, accumulation and inversion, and thus between the conductive and resistive behaviour of the Schottky barrier, which is, therefore, also called a Schottky diode.

CB VB χ φ SC IP,SC E F E g φ M Before contact A�er contact qV BB Φ B,n SC Φ B,n φ M - = χ SC Barrier height: qV BB φ M φ SC - = Band bending: Pinning Φ B,n Semiconductor Metal Semiconductor Metal Semiconductor Metal Barrier height: Φ B,n - = Φ 0 S . φ M x = 0 x 0 States Surface
The space charge region width can be calculated with Poisson's equation and depends strongly on the doping concentration. Below, the calculation is made for the depletion region width of an n-type semiconductor. In one dimension, Poisson's equation equates to:

d 2 V dx 2 = - ρ ϵϵ 0 (2.9)
with V being the potential, ϵ the dielectric constant, ϵ 0 the vacuum dielectric constant and ρ the charge density, whereby for an n-type depletion region:

ρ = N d e (2.10)
with N d the ionized donor density. The potential V can be replaced by the electron energy φ:

d 2 φ dx 2 = N d e 2
ϵϵ 0 (2.11) In the absence of Fermi level pinning, following boundary conditions can be defined:

φ = ( dφ dx ) = 0 at x = 0 (2.12) φ = φ M -φ SC at x = x 0 (2.13)
whereby x = 0 corresponds to the beginning of the depletion region and x 0 is the position of the semiconductor/metal contact. By integrating Poisson's equation according to the boundary conditions, following solution can be obtained:

φ M -φ SC = N d x 2 0 e 2 2ϵϵ 0 (2.14)
The depletion region width x 0 , thus, scales with N nm and 100 nm are typical. In photocatalysis the depletion region width plays an important role, as it is the parameter which state how far from the interface a photon-induced charge carrier is under the influence of the interface-induced internal electric field, which helps to separate the charge carriers. In a band diagram the internal electric field is commonly visualised through band bending, which results from the electron electrochemical potential changing with respect to the band edges due to the presence of the internal electric field.

Deple�on Accumula�on Inversion

CB When an n-type and p-type semiconductor are brought into contact, a p-n junction may be established, whereby a depletion region forms at the interface between the two semiconductors.

VB E F + + + CB VB E F + + + + + + + + CB VB E F ---
The band diagram of a p-n junction is easily visualised for homojunctions, whereby the p-type and n-type semiconductor are made from the same material. For instance, Figure 2.16 shows how a homo p-n junction is formed when p-type and n-type Si are brought into contact. Because of the different Fermi levels, charge transfer takes place between p-type and n-type Si, forming a depletion region near the interface at both sides of the interface. The width of the space charge region (w p -w n ) can again be calculated through the integration of the one dimensional Poisson's equation (Equation 2.9) with the following boundary conditions:

( dφ dx ) = 0, φ = φ SC,n at x = -w n (2.15) ( dφ dx ) = 0, φ = φ SC,p at x = w p (2.16) φ = φ 0 ; dφ dx and φ are continuous at x = 0 (2.17) w p * N A = w n * N D Charge neutrality (2.18)
with w n ,w p corresponding to the start of the depletion region in the n-type and p-type semiconductor, respectively, N d ,N a corresponding to the ionized donor and acceptor concentration, respectively, and x = 0 corresponding to the position of the semiconductor/semiconductor contact. After integration of Poisson's equation with the above defined boundary conditions:

φ SC,p -φ SC,n = e 2 N d w 2 n [(N d /N a ) + 1] 2ϵϵ 0 (2.19)
Hence, the depletion region width for a semiconductor/semiconductor junction is again dependent on the doping concentration in the semiconductors. Due to the depletion region, the p-n junction acts like a diode. Under forward bias, the depletion region width decreases, thus allowing an electric current to pass, whereas for a reverse bias the depletion region width increases, making the area around the p-n junction interface even more insulating. The creation of p-n junctions is a common strategy in the design of photovoltaics to increase charge carrier separation, because photons absorbed within the depletion region give rise to electron-hole pairs which are more efficiently separated due to the internal electric field.

For a heterojunction between two different semiconductors, the bands do not form a continuity at the interface, due to the difference in conduction and/or valence band positions. In the absence of Fermi level pinning, the Anderson model states that:

∆E CB = χ (SC2) -χ (SC1) (2.20) ∆E VB = I p (SC2) -I p (SC1) (2.21)
with I p (SC1), I p (SC2) the ionization potentials of semiconductor 1 and 2, respectively, and χ (SC1), χ (SC2) the electron affinities of semiconductor 1 and 2, respectively. However, Fermi level pinning may also take place between two semiconductors, with altered band discontinuities as a consequence. Ideally, the lattice and thermal expansion coefficient mismatch between the two contact layers should be as small as possible to avoid the creation of interface defects which may induce inter-gap states. and recognized "cocatalysts" such as NiO and CoO x was studied to determine whether the interface properties, and then in particular band bending, could be responsible for the previously observed improved efficiencies.

Heterostructures, in this work, have been mostly investigated through interface experiments, which involved photoelectron spectroscopy and thin film sputtering. The experimental procedure of an interface experiment is fully described in Section 3.4.4.1 of Chapter 3. In short, a certain substrate was introduced into an integrated thin film deposition and surface analysis system, which operates entirely under ultra-high vacuum. After removing the extrinsic impurities from the substrate surface through an adequate cleaning procedure, a thin film contact was stepwise sputtered on top of the substrate, whereby photoelectron spectra were measured in between each deposition step. The interpretation of the solid state contact could then be obtained through the analysis of the temporal evolution of the core level spectra, valence band spectra and ultraviolet photoelectron spectra measured after each deposition step. Due to the overlay of the contact layer valence band spectra with the substrate valence band spectra, the substrate valence band maximum, which gives information on the band bending and creation of a space charge region, has to be monitored through the binding energy (BE) changes in the core level spectra of the substrate, a method which was first described by Kraut and Waldrop. 48 In our group the Kraut method has been used extensively for the interface study of various semiconductor heterojunctions. and anatase/ITO to the band alignment of rutile/RuO 2 and rutile/ITO. In addition to the study of oxide interfaces, the Kraut method has been used in our group to determine the band alignment of particular II-VI semiconductors. [52][53][54] All in all, our studies have shown that the Kraut method can be used to study the band alignment for a wide variety of materials and draw conclusions on the junction properties of the interfaces. Properties such as band discontinuities, band bending and Fermi level tunability (variation E F -E VBM ) can be determined, as well as the relative alignment between materials, by assuming transitivity of the alignments of those materials with RuO 2 and ITO. 55,56

Semiconductor-liquid junction

The interface between semiconductor and electrolyte is arguably the most important interface in photo(electro)chemical water splitting, as the transfer of the charge carriers from the photoabsorber to the adsorbed reactants is a crucial step and is what distinguishes photo(electro)chemical water splitting from photovoltaic driven electrolysis. For charge carrier transfer to take place the energetic state of a reactant should lie within the band gap of the semiconductor. As already outlined above, the valence band maximum should be more positive than the O 2 /H 2 O (1. 23 V) standard reduction potential and the conduction band minimum should be more negative than the H + /H 2 (0.0 V) standard reduction potential to allow charge transfer in the water oxidation reaction and/or water reduction reaction, respectively. Generally, when a semiconductor is placed in an electrolyte, charge transfer will take place until the Fermi levels of semiconductor and solution are equilibrated (Figure 2.17).

Figure 2.17: Band diagram of semiconductor (a) before and (b) after contact with electrolyte containing a reversible one electron redox couple under open circuit and in the absence of light. V H is the potential drop of the Helmholtz layer; V B is the band bending; E F is the Fermi level of the system at equilibrium; ∆E F is the difference between the Fermi level and the conduction band edge; E 0 F,redox is the standard Fermi level of the redox couple; E unoc and E oc , are the energies of the unoccupied states and the occupied states of the redox couple, respectively; and λ is the reorganization energy. 57 The Fermi level of a solution can easily be understood for solutes undergoing a one electron redox process, whereby the solutes are defined as electron acceptors (A) or electron donors (D), depending on whether they accept an electron from or donate an electron to the semiconductor. For electron donors and electron acceptors the energy states of importance in charge transfer are the highest occupied energy state E occ and the lowest unoccupied energy state E unocc , respectively. When an electron is transferred from or to the semiconductor, the charge of the solute changes, giving rise to A -and D + , respectively. The change in charge makes that the solvation shell around the solute re-orients, which is associated with a change in free energy of the solute, referred to as reorganization energy λ. In addition, thermal vibrations lead to a fluctuation of the solvation shell, so that the energy state of the solute with a particular charge is not associated with one specific energy level E occ or E unocc , but rather with a thermal distribution of states with E occ or E unocc being the maxima of those distributions. Now, the Fermi level E F of the redox couple (A/A -or D/D + ) is defined as the energy level lying midway of E occ and E unocc , and which can be expressed as follows:

E F,A/A -= E A/A -= E 0 A/A -+ RTln ( a A - a A ) (2.22) E F,D/D + = E D/D + = E 0 D/D + + RTln ( a D + a D ) (2.23)
For a multi-electron process, the situation can be somewhat more complicated, because more than two states could be present in the electrolyte. When more that two states are present, the standard redox potentials of the one-electron steps of the overall reaction govern the charge carrier transfer from semiconductor to electrolyte. Now, the Fermi level of the solution may not be equal to the standard redox potential of the overall reaction due to the overall standard redox potential not necessarily lying midway between the maxima of the E occ and E unocc distributions.

Charge transfer from the semiconductor to the electrolyte leads to a space charge region in the semiconductor, just like for the semiconductor/metal and semiconductor/semiconductor junctions. The band bending due to the semiconductor/liquid is proportional to the difference in semiconductor and solution Fermi level, when no Fermi level pinning takes place. However, in solution due to the charged semiconductor surface, counterions are attracted towards the surface, which results in the formation of a so-called Helmholtz double layer. 58 The Helmholtz layer has a thickness on the order of one angstrom, and leads to an additional potential drop inside the space charge region of the semiconductor. Therefore, to determine the exact band bending value, the potential drop associated with the creation of the Helmholtz layer should be known.

The flat band potential V fb is a measurable physical quantity, which has been introduced to connect the semiconductor energy levels to those of the electrolyte. In essence, the flat band potential is the electrode potential for which the potential drop in the space charge region becomes zero. Nozik et al. expressed V fb , with the Normal Hydrogen Electrode (NHE) as reference electrode, as follows: 59 V fb (NHE) = χ SC + ∆E F + V H + E 0 (2.24) whereby ∆E F stands for the energy difference between the Fermi level and the majority carrier band edge, E 0 represents a scale factor between the absolute vacuum level and the standard reference electrode potential (for NHE this amounts to -4.44 eV), V H being the potential drop across the Helmholtz layer and χ SC being the semiconductor electron affinity. V H is independent of the charge carrier transfer between semiconductor and electrolyte, and depends solely on the adsorption/desorption equilibrium of electrolyte ions on the semiconductor surface. 60 The potential drop across the Helmholtz layer is mainly governed by the so-called potential determining ions. For an oxide semiconductor in an aqueous electrolyte, the hydroxide and hydrogen ions are the only potential determining ions, 61 which means that V H is completely governed by the pH of the solution. In the absence of a Helmholtz layer V H becomes zero and is referred to as the point of zero charge (pH ZPC ). At pH ZPC the flat band potential is equal to the semiconductor Fermi level, and is the only relevant flat band potential. Under the flat band potential at pH ZPC the band positions at the semiconductor/electrolyte interface are the same as the band positions in the bulk of the semiconductor, so that the semiconductor can be positioned on the vacuum scale when combined with the knowledge on E g and ∆E F . This has also been illustrated in Figure

Band edge position determination

There are various photo(electro)chemical techniques available to determine the flat band potential, whereby the Mott-Schottky method is used the most. Central in the Mott-Schottky method is the Mott-Schottky equation which relate the capacitance to the applied voltage across the semiconductor/electrolyte interface. The Mott-Schottky equation can be derived from Poisson's equation (Equation 2.9) using Gauss' law to relate the electric field across the semiconductor/electrolyte interface to the charge contained within that region, and the Boltzmann distribution to describe the electron distribution within the space charge region. For a full derivation I refer the reader to the publication of Geldermann et al. 62 For an n-type semiconductor the Mott-Schottky equation becomes:

1 C 2 = 2 ϵϵ 0 A 2 eN D ( V -V fb - k B T e ) (2.25)
whereby C is the differential capacity ∂Q ∂V , A is the interfacial area, N d is the donor density, V is the applied voltage, k B stands for Boltzmann's constant, T is the absolute temperature, ϵ is the dielectric constant of the semiconductor, ϵ 0 is the vacuum permittivity and e is the elementary charge. Plotting 1 C 2 relative to the applied voltage V should give a linear correlation, whereby V fb can be found from the intercept on the x axis. Additionally, N d can be determined from the slope, if ϵ and A are known. Another commonly used photo(electrochemical)method to determine the flat band potential is by determining the anodic photocurrent onset potential from a photocurrent-potential plot. 63 The Mott-Schottky method was quite popular 30 -40 years ago, having been used to determine the semiconductor band positions of many semiconductors. 61,[63][64][START_REF] Tokunaga | Selective preparation of monoclinic and tetragonal BiVO 4 with scheelite structure and their photocatalytic properties[END_REF][START_REF] Irwin | p-Type semiconducting nickel oxide as an efficiency-enhancing anode interfacial layer in polymer bulk-heterojunction solar cells[END_REF] Today, many of those semiconductor band positions are being published relative to each other on the standard hydrogen electrode (SHE), NHE, reversible hydrogen electrode (RHE) or absolute vacuum scale. 13,22,[START_REF] Barreca | Composition and Microstructure of Cobalt Oxide Thin Films Obtained from a Novel Cobalt(II) Precursor by Chemical Vapor Deposition[END_REF][START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF][START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF] However, the notion on the stringent experimental conditions to perform exact Mott-Schottky measurements may be lost with the passing of time. I, therefore, refer the reader to the letter of Cardon et al., who listed up the semiconductor and electrolyte conditions which should be met to perform reliable Mott-Schottky measurements. [START_REF] Lee | Mechanism for the increase of indium-tin-oxide work function by O 2 inductively coupled plasma treatment[END_REF] For instance, conditions such as the complete absence of surface states and the allowed presence of only one type of localised electronic defect are already complicated enough to be met by any semiconductor. In the end, only by measuring on thoroughly cleaned specific crystal facets of a precisely doped semiconductor single crystal with a certain surface termination, valid and reproducible Mott-Schottky measurements can be obtained.

Another problem arises from listing up just the band positions on the RHE scale or at pH 0 on the NHE scale, due to the loss of a valuable piece of information, namely the pH ZPC . This means that for oxide semiconductors whose band positions show a Nernstian behaviour with pH, i.e. band position variation according to 0.059 V/pH, 61,[START_REF] Tokunaga | Selective preparation of monoclinic and tetragonal BiVO 4 with scheelite structure and their photocatalytic properties[END_REF]71 a pH-dependent discrepancy exist between bulk and surface band edge positions. Nonetheless, is the pH ZPC very surface specific so that strong variations between differently prepared semiconductors of the same crystalline structure may occur. Determining pH ZPC is, therefore, anyway recommended for every synthesized system to be able to completely evaluate the interfacial energy diagram.

Besides experimental methods, band edge positions have also been tried to be calculated. However, most Density Functional Theory (DFT) methods are inappropriate for measuring the band edge positions, requiring the introduction of a semi-empirical Coulomb parameter U (DFT+U). 72,[START_REF] Liechtenstein | Density-functional theory and strong interactions: Orbital ordering in Mott-Hubbard insulators[END_REF] A more simpler calculation, which was introduced by Butler and Ginley, involves the calculation of the electronegativity of the semiconductor, which is considered to be the geometric mean of the electronegativities of the constituent atoms. For binary oxide semiconductors M a X b the equation becomes: 61

E C = -(χ a M χ b X ) 1 a+b - 1 2 E g + 0.059(pH ZPC -pH) + E 0 (2.26)
whereby E c stands for the conduction band position, E 0 is the scaling factor between electrode potential and absolute vacuum level χ M , χ X are the absolute electronegativities of M and X, and 0.059(pH ZPC -pH) reflects the pH dependence of the oxide band positions. Butler and Ginley found that the band positions for many semiconductors calculated through the electronegativity correlate well with those determined through experimentally determined flat band potential measurements (Figure 2.19). 61 In this work, the theory behind band edge determination given within this Section will be used to evaluate whether the information on the independently determined band edge positions of two semiconductors can be used to evaluate the relative band alignment between those two semiconductors. Specifically, in Section 5.3.3 in Chapter 5 the solid state band alignment schemes which are frequently depicted in literature, 22,[START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF] and which were obtained through the aboveexplained experimental or theoretical methods of individual band position determination, will be compared to the solid state band alignment estimated from interface experiments with RuO 2 and ITO.

Interfacial charge transfer kinetics

Now that the energetic conditions of photo(electro)chemical water splitting have been described, the interfacial charge transfer kinetics will be introduced briefly in this Subsection. 74,75 Most of the aspects with regards to the water splitting reaction kinetics described below were derived from electrocatalytic water splitting, but are valid for photo(electro)chemical water splitting as well. The overall water splitting reaction should be split up into the water oxidation reaction and the water reduction reaction, which in electrocatalysis are often denoted as the oxygen evolution reaction (OER) and hydrogen evolution reaction (HER), due to the different mecha- overall hydrogen evolution reaction. 78,79 If hydrogen would bind too strongly or too weakly to the surface, the overall reaction rate would be limited by, respectively, the adsorption step (Volmer step) or the desorption step (Tafel/Heyrovsky). Therefore, ∆G H should preferably be close to zero. Indeed, a plot of the experimentally determined exchange current densities of a series of HER electrocatalysts as a function of ∆G H , determined through DFT calculations, showed a so called volcano activity relationship with ∆G H ≈ 0 corresponding to the maximum of the volcano (Figure 2.20b). [80][81][82] However, ∆G H being close to zero is not the only condition for high HER performance, the kinetic barriers for the elementary steps need to be accounted as well.

Although, MoS 2 has a ∆G H comparable to the noble metals, lower HER activities have been found for MoS 2 , which are supposedly caused by higher kinetic barriers. 83 A few authors have tried to model the kinetic barriers for the proton transfer reactions, [82][83][84][85][86] There are two different OER mechanisms, one which involves four proton-electron transfers, denoted as four-electron pathway, [88][89][90][91] and another which involves two times a two-protonelectron transfer with hydrogen peroxide as intermediate, denoted as the two-electron-twoelectron pathway. 92 Depending on the oxygen dissociation barrier on the electrocatalyst surface a four-electron pathway can either be associative (Figure 2 electrocatalyst should be optimised for all elementary steps, meaning that the binding energies of all three intermediates OOH*, OH*, O* on the catalyst surface should be optimal to achieve the highest OER rate. However, the adsorption energies of these intermediates have been found

to be strongly correlated and cannot be easily decoupled due to scaling relations. 91 The most persistent scaling relation is the one between the adsorption free energies of OOH* and OH*, which amounts to 3.2 ± 0.2 eV for a large set of catalyst surfaces. 75,90,93 The ideal adsorption free energy difference between OOH* and OH* is 2.46 eV so that an overpotential of 0.4 ± 0.1 eV ((3.2 ± 0.2 eV -2.46 eV)/2 = 0.4 ± 0.1 eV) seems to be inevitable. With ∆G OOH-OH fixed, the tunable parameter which has been found to largely determine the activity of electrocatalyst surfaces is the adsorption free energy difference between O* and OH* ∆G O-OH . 75 The highest OER activity 94 Herein, for both OER and HER, non-noble metal compounds generally required higher overpotentials than noble metals in acidic media, whereas in basic media many non-noble metal compounds performed on par or better compared to noble metals. Hence, transition metal compounds could potentially replace noble metals in industrial electrocatalytic water splitting. The hydrogen production cost would then be slightly lower, on the condition that similar stabilities as for noble metal electrocatalysts can be achieved. The lower activation energies of certain transition metal compounds led also to their introduction as cocatalysts in photo(electro)chemical water splitting. [95][96][START_REF] Chen | Reduction-induced Fermi level pinning at the interfaces between Pb(Zr,Ti)O 3 and Pt, Cu and Ag metal electrodes[END_REF] For a more thorough description of electrocatalyst kinetics and an overview of the recent state of the art HER and OER systems, I refer the reader to the review written by Seh et al. 54 Figure 2.22: Specific activities per electrochemically active surface area, expressed as the overpotential needed to obtain a current density of 10 mA cm -2 , for a series of HER and OER electrocatalysts tested in acidic and basic media. 94

Photo(electro)chemical water splitting materials

The primary material in photo(electro)chemical water splitting is the photoabsorber. As discussed in Section 2.1.3 only few photoabsorbers are suitable for overall water splitting with regards to their band gap and band positions. Therefore, also photoabsorbers with reduced band gaps are considered, which can only be used to perform either water reduction or water oxidation, due to the photoabosrbers having a suitable conduction band minimum or valence band maximum, respectively. As such, more than 150 compounds have been recognised for (partial)

water splitting, which include metal oxides, [START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF][START_REF] Körber | Barrier heights at the SnO 2 /Pt interface: In situ photoemission and electrical properties[END_REF] sulphides, [START_REF] Giesecke | Modification of the Schottky barrier height at the RuO 2 cathode during resistance degradation of Fe-doped SrTiO 3[END_REF] [START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF]109,110 In the next Subsections, a short overview of the properties, design strategies and state of the art photo(electro)chemical water splitting efficiencies are given for the photoabsorbers selected for this work: BiVO 4 and CuFeO 2 .

BiVO 4

BiVO 4 is an intrinsic n-type semiconductor for which four common polymorphs have been distinguished: orthorhombic pucherite, tetragonal dreyerite, monoclinic clinobisvanite and the closely related tetragonal scheelite-like phase. 25,111,112 However, only the monoclinic clinobisvanite is active for photo(electro)chemical water oxidation. The monoclinic BiVO 4 crystal structure (Figure The electronic states in the BiVO 4 VBM have been found to originate from the hybridization of O2p and Bi6s orbitals, whereas the CBM consists of V3d x 2 -y 2 /z 2 states. 113,114 Depending on the determination method, CBMs between 0.0 and 0. m -2 V -1 s -1 for W:BiVO 4 and related this weak mobility to electron trapping at tungsten induced bulk defects. 122 Other authors, however, state that the low majority charge carrier mobility is intrinsic to the BiVO 4 electron conduction mechanism, for which the thermally activated small polaron hopping mechanism iii has been postulated. 114,119 For BiVO 4 the electrons would then reside at the vanadium lattice sites, implying a reduction of V 5+ to V 4+ . [123][124][125] To move from vanadium to vanadium lattice site, a polaron hopping barrier of 0.46 eV has been calculated by iii Polarons often form in transition metal oxides when the velocity of an electron is low enough, so that the position of the atoms surrounding the electron are altered due to the electron's presence. The atom displacement creates a potential well, which reduces the carrier's energy. As a consequence, the carrier can become self-trapped and may only move when thermal vibrations lead to favourable atom displacement. Self-trapping distinguishes itself from trapping, due to the self-trapped carrier still being able to move. A distinction is made between small and large polarons which extent over one or multiple structural units, respectively. Large polarons, generally, have drift mobilities > 1 x 10 -4 m -2 V -1 s -1 , whereas small polarons exhibit drift mobilities of << 1 x 10 -4 m -2 V -1 s -1 . The band gap of BiVO 4 has been a point of discussion, with respect to it being direct or indirect. However, now it has been generally accepted that the band gap of BiVO 4 is indirect with only a 70 to 200 meV difference between the first indirect allowed optical transition and the first direct allowed optical transition. 113,128 The direct allowed transition is mostly considered to be at 2.4 -2.5 eV. 114,121,129,130 High photocurrents have been demonstrated with BiVO 4 based thin film photoanodes under external biases or through coupling to a photovoltaic system. In particulate water splitting, BiVO 4 can be used as the water oxidation compound in a Z-scheme type system (Figure 2.12). 134 Into these BiVO 4 Z-scheme systems, CdS, 135,136 Zn x Cd 1-x Se 137 and rhodium-doped SrTiO 3 36,134,138 have been incorporated most frequently as the water reduction compounds. Liquid redox mediators like Fe 2+ /Fe 3+134 and IO 3 -/I -137 have been used as well as solid state redox mediators, such as carbon dots 135 and reduced graphene oxide. 34 The highest STH efficiencies with Z-scheme type systems have been reported when the particles were immobilised on a conductive sheet. By immobilising particulate Mo:BiVO 4 /RuO x and Rh,La-codoped SrTiO 3 /Ru on a 10 x 10 cm conductive carbon film (Figure 2.24), unassisted pure-water splitting with a state of the art 1.0% STH efficiency could be realised under ambient conditions (331 K and 91 kPa). Also the immobilization on a gold layer yielded a high 1.1% STH efficiency under reduced pressure (331 K and 10 kPa). 36 For particulate photocatalysts, contact materials were until recently solely deposited in a random manner. Random cocatalyst coverage, however, provide no distinct flow path for photogenerated holes and electrons, so that the problem of bulk and surface recombination could still persist due to insufficient charge carrier separation. Anisotropic photodeposition now allow cocatalysts to be deposited at specific sites on anisotropic crystalline particles. The idea is that after Based on the reported band gaps, photocurrents of up to 15 mA cm -2 should be feasible with CuFeO 2 . 148 However, the published water splitting efficiencies have been rather low so far. 95 These reports show that the addition of a "cocatalyst" benefits the CuFeO 2 water reduction efficiency. Furthermore, hydrogen evolution with CuFeO 2 only seems to be feasible under an applied bias. The need for a bias suggests that the necessary photovoltage for water reduction cannot be reached. Indeed, Prévot et al. found that, due to Fermi level pinning, only a photovoltage of 0.3 V could be generated, whereby the Fermi level pinning was associated with surface states at the CuFeO 2 /electrolyte interface. Other suggested reasons for the weak hydrogen evolution efficiencies include Fermi level pinning at inter-gap states in the bulk, 162 photocorrosion 152 and poor catalytic activity. 95 In this study (Chapter 6), particulate CuFeO 2 with a defined morphology was synthesized according to a hydrothermal procedure published by Jin et al. 28 However, the authors found that a thin film, prepared by dropcasting the hexagonal 2H-CuFeO 2 nanoplatelet powder, only led to a limited photocurrent of 8.8 µA at 0 V RHE . In this work, we wanted to study whether heterostructuring of the CuFeO 2 nanoplatelets, would allow unbiased particulate photochemical water reduction. Due to the specific morphology and the observed anisotropic charge carrier dynamics, we theorized that the creation of anisotropic CuFeO 2 heterostructures should be feasible through photodeposition. Hence, the photodeposition of several metal(oxide)s and the characteristics of the resulting CuFeO 2 heterostructures was studied. The performance of the CuFeO 2 heterostructured powders was assessed through photochemical sacrificial water splitting in a batch reactor equipped with an automatic burette. Furthermore, interface experiments were performed between pelletized CuFeO 2 powder as substrate, and ITO and water as contact layers to investigate bulk Fermi level pinning phenomena, which could limit the CuFeO 2 photovoltage.

Photoelectron spectroscopy

Photoelectron spectroscopy (PES) focuses on the energy analysis of electrons emitted from a substance through the photoelectric effect. The photoelectric effect takes place when an electron is excited by a photon, which has a high enough energy to overcome the binding energy of the electron to the substance, so that the electron is ejected from the substance as a result. Electrons liberated in this way are called photoelectrons. The energy needed to eject the electrons depends on how the electrons are bound in the substance. Photons with an energy of a few eV are sufficient to release delocalized electrons, whereby the lowest photon energy that allows the ejection of the electrons from a solid, is referred to as the solid's work function. Electrons residing in the valence orbitals of an atom, molecule or compound have a higher binding energy and generally require photons from the ultraviolet spectral region to be released. PES with a UV light source is commonly referred to as ultraviolet photoelectron spectroscopy (UPS). Electrons closer to the atomic core, have an even higher binding energy than the outer valence electrons, necessitating the absorption of higher energy photons. Generally, photons from the X-ray spectral region are needed to liberate these core electrons. X-ray photoelectron spectroscopy (XPS) encompasses the technique, whereby photoelectrons, emitted by an X-ray source, are analysed.

In a common photoelectron spectrometer (Figure 2.27), a monochromatized source of photons is focused onto a sample. The photoelectrons released from the sample pass through a set of electrostatic lenses, which decelerate the electrons and focus them into the entrance slit of a hemispherical analyser, which is used to only let electrons with a specific kinetic energy (pass energy) reach the detector channels. The path from sample to detector should be operated under UHV to avoid inelastic scattering of the photoelectrons. Photoelectron spectrograms are generally recorded in sweeping mode, whereby a certain energy region is scanned one or multiple times by fixing the pass energy and sweeping the voltage applied to the electrostatic lenses in such a manner that electrons with a selected kinetic energy are counted for a selected amount of time by the channels of the detector. A certain set of lens functions correspond to a certain electron kinetic energy (E kin ) and, thus, to a certain electron binding energy (E B ) according to: Φ sp is a parameter, specific to the photoelectron spectrometer, which can be corrected for by measuring the Fermi level of a cleaned metallic sample and appointing it to a binding energy of zero. This energy calibration is maintained for other conducting samples in electrical contact to the spectrometer, as their Fermi level align with that of the spectrometer. Therefore, the photo-electron binding energies of a conducting sample are measured in reference to the Fermi level of the sample.

E B = hν -E kin -Φ sp (2.
An important notice which has to be made is that the binding energy mentioned in Equation 2.42 not only depends on its initial bonding state but also on so called "final state" effects, which refer to the state of the photohole after the electron is ejected, and extrinsic effects through which the photoelectron loses/gains kinetic energy in between its creation and its escape from the material. An extensive overview of those effects can be found in the PhD dissertation of Mirko Weidner. 163 Some effects worth mentioning include spin-orbit coupling, electron shakeup, electron shake-off, plasmon excitation and phonon interactions. Spin-orbit coupling is a final state effect, whereby the spin of an unpaired core electron left behind after photo-emission couples with the angular momentum of the orbital. This effect is noticeable for photo-emission of electrons residing in p,d,f,… orbitals and is the reason why, for instance, the Bi4f core level spectrum exhibits two distinct signals, denoted as Bi4f 7/2 and Bi4f 5/2 , whereby 7/2 and 5/2 are the possible quantum number sums of angular (f-orbital: 3) and spin quantum number (+1/2 or -1/2).

Electron shake up and electron shake off phenomena are extrinsic effects, whereby an energy transfer takes place from the photoelectron to another electron that is excited to a higher state or ejected from its atomic core, respectively. These phenomena are in part responsible for the satellite features seen in many transition metal 2p spectra such as Co2p, Fe2p and Ni2p. During plasmon and phonon interactions energy transfer takes place between photoelectrons and the electron gas (delocalized electrons) or phonons, respectively, whereby seemingly higher binding energies are registered for the photoelectrons that were involved in those processes.

PES is, furthermore, a surface sensitive technique. The liberated photoelectrons have kinetic energies which correspond to mean free paths in the range of 0.5 -5 nm (Figure 2.28) and the PES information depth is generally considered to be ± 3 times the electron mean free path. Hence, the information depth is limited from a few atomic layers to ± 15 nm. For a more detailed description of photoelectron spectroscopy, I refer the reader to the work of Stefan Hüfner. 7,164 Next, it will be explained how the chemical state of an element can be determined from its most intense core level spectrum, specifically for the core level spectra displayed in this work.

Transition metal core level spectra

Deriving the chemical state of a transition metal from one of its photoelectron core level spectra may be challenging, due to the variety of final state effects that may take place, especially for oxidation states where the metal d-orbitals are partly filled. For some transition metals, it may even be impossible to quantify how much of each species is present in a certain composition, due to the spectra of these species not being well resolved from each other. For these transition metals, only qualitative statements can be made. The discussion below will focus on how certain oxidation states may be determined from the shape of the spectra as well as from the binding energies of certain features of the main core level photoelectron lines of transition metal (hydr)oxides.

V2p

The V2p spectrum of vanadium oxides with a single oxidation state generally contain two main signals which are split from each other due to spin-orbit coupling. 166 These signals are generally denoted as V2p 3/2 and V2p 1/2 . The V2p 1/2 signal is broadened with respect to V2p 3/2 due to Coster-Kronig effects, 167 which involve the rapid (in the order of 10 -15 s) filling of the photohole by an electron of a higher-energy shell with the same principal quantum number as the photohole. Coster-Kronig effects give rise to broad line shapes. 168 As can be seen in Figure 2.29 and Table 2.1, oxides with different vanadium oxidation states have distinct V2p 3/2 and V2p 1/2 binding energies for each oxidation state, whereby peak fitting can be applied to give an estimate of the surface composition of a mixed vanadium oxide material. The V2p 3/2 line of V 5+ is sharper with respect to the V2p 3/2 line of the other oxidation states, because V 5+ does not contain any unpaired electrons which may give rise to additional final state effects. Contrarily, V 4+ and V 3+ have unpaired electrons, causing a multiplet splitting and, thus, a broader line shape.

Fe2p

The Fe2p spectrum is considerably more complicated than the V2p spectrum. Most

Co2p

The Co2p

Ni2p

One of the most complicated transition metal spectra with respect to species determination is undoubtedly the Ni2p spectrum, as can be seen in

Cu2p

The Cu2p 3/2 and Cu2p 1/2 lines are separated by about 20 eV. The main Cu2p 3/2 emission line can be used to distinguish cupric oxide from cuprous oxide and metallic copper, as the binding energy of the Cu2p 3/2 line of copper(II) oxide generally lies at about 933.6 eV, whereas the Cu2p 3/2 binding energies of metallic copper and copper(I) oxide generally can be found at about 932.4 -932.6 eV, as can be seen in Figure 2.34 and Table 2.2. 166,177 Additionally, the Cu2p spectrum of Cu 2+ distinguishes itself from metallic copper and Cu + by an apparent shake-up satellite feature at about 941 eV. 178 The Cu2p spectra of metallic copper and cuprous oxide are nearly indistinguishable. Instead, the CuLMM Auger spectrum needs to be recorded to make the distinction between metallic copper and cuprous oxide, as both have a different CuLMM binding energy, as can be seen in Figure 2.35 and Table 2.2. The Ag3d spectrum is used most often to determine the silver species. Generally, the Ag3d spectrum consist of two main signals, Ag3d 5/2 and Ag3d 3/2 , which originate from spin-orbit coupling. Remarkably, the Ag3d 5/2 and Ag3d 3/2 binding energies of silver oxides have been found at lower binding energies (Table 2.3), which is apparently caused by initial state effects related to an ionic charge shift and the lattice potential. 179 Additionally, for metallic silver the line shapes are slightly asymmetric due to bulk and surface plasmon loss features.

Pt4f

Just like the Ag3d spectrum, the Pt4f spectrum consists of two clear signals, which are denoted as Pt4f 7/2 and Pt4f 5/2 . Due to bulk and surface plasmon loss features slightly asymmetric line shapes are found for metallic platinum. The commonly found binding energies for Pt, PtO and

Bi4f

The Bi4f spectrum is very similar to the Pt4f spectrum in that it only consists of two main signals, split from each other by spin-orbit coupling, Bi4f

O1s

The O1s spectrum is often used to determine the oxygen species and consists for a singular oxygen species of 1 main signal with a Pseudo-Voigt shape. Metal oxides can be easily distinguished in O1s spectra as they mostly exhibit O1s binding energies in the range of 529.0 -530.2 eV. 166,171 Metal hydroxides can be found in the range of 531 -532 eV. 177,180 However, to determine the metal hydroxide content the sample should not have been exposed to air as the range of 531 -533 eV is also typical for alcohol and carbonyl groups. 181,182 Water can also easily be distinguished, because it gives rise to a broad signal in the range of 533 -534 eV. 177,183,184

Conclusion

In this Chapter an overview was made of the most important fundamental principles and recently published insights regarding semiconductors, solid state and semiconductor-liquid interfaces, photo(electro)chemical water splitting, BiVO 4 , CuFeO 2 and photoelectron spectroscopy. These fundamentals and insights are essential to understand and appreciate the results that will be displayed in the following Chapters of this work, concerning the interface studies with BiVO 4

and CuFeO 2 substrates, the creation of particulate BiVO 4 and CuFeO 2 heterostructures, and the assessment of the photochemical water splitting performance of the heterostructured powders.

In the next Chapter, the experimental methods, techniques and procedures will be detailed which were used to prepare and analyse the BiVO 4 and CuFeO 2 based materials and interfaces.

(15) Fujii, K.; Karasawa, T.; Ohkawa, K.

Preparation of BiVO 4 microparticles

The BiVO 4 microparticles were synthesized according to a published procedure. 1 First, 1.25 mmol bismuth(III) nitrate pentahydrate and 1.25 mmol ammonium metavanadate were dissolved in 25 mL of a 2.0 mol/L nitric acid solution and 25 mL of a 2.0 mol/L ammonia solution, respectively.

To dissolve the ammonium metavanadate, 5 minutes of sonication was occasionally necessary.

To the bismuth(III) nitrate solution, 0.85 mmol sodium dodecyl sulphate was added as a templating agent. Under violent stirring (± 1000 RPM) the bismuth(III) nitrate solution was dropwise added to the ammonium metavanadate solution. After a few drops of the bismuth(III) nitrate solution the reaction mixture started to foam and a precipitate formed. After 30 min of stirring at 1000 RPM, the pH was altered from ± 0.5 to a more basic value (exact value depends on the preferred morphology (See Section 4.2 for exact values) using a 2.0 mol/L ammonia solution.

After additionally stirring for 2 h, the reaction mixture was added to a specially designed glass vessel. This glass vessel was put into a 200 mL pressure reactor and heated up to 180 ℃ using an electric clamp-on band heater for 24 h. Then, the heater was removed and the pressure reactor was allowed to cool down. After 1 h the reaction mixture containing a vivid yellow powder was washed three times with distilled water and three times with ethanol to remove unreacted precursor, waste products and templating agent. Finally, the BiVO 4 powder was dried for 8 h at 70

℃ in a muffle furnace.

Preparation of hexagonal CuFeO 2 nanoplatelets

The preparation of the hexagonal CuFeO 2 nanoplatelets was adapted from a previously published procedure. 2 In a typical synthesis, 14 mL ultra pure (18 mQ) water and a triangular stir bar were added to a 125 mL Teflon beaker. After placing a long needle in the water, nitrogen was bubbled through for at least 30 minutes while stirring at 750 RPM. 3 mmol (571 mg) copper iodide was then added at a stirring speed of 750 -1000 RPM, which formed a grey suspension. Immediately, 3 mmol iron chloride hexahydrate (811 mg) was added, making the suspension turn brown-green.

After 15 min stirring at 1000 RPM, 0.22 mol (8.8 g) NaOH was added, pellet by pellet, over a period of 5 min. The addition of NaOH went together with heat formation and made the suspension turn chocolate-brown. After 30 min of stirring at 1000 RPM the Teflon beaker was inserted in a fitting stainless steel pressure reactor (Figure 3.2) and placed in a preheated oven at 200 ℃ for 3.5

h. The pressure reactor was then removed from the oven and placed on the bench to cool down for 1 h. Subsequently, the black CuFeO 2 powder was separated from the reaction mixture and washed with distilled water by a sequence of centrifugation and redispersion that was repeated for at least 6 times until reaching neutral pH. Finally, the powder was dried at 70 ℃ during 8 h.

Photodeposition

In this work, anisotropic heterostructured photocatalysts were created through photodeposition.

To prepare a certain deposit, photocatalytic particles were first added to a honeypot-shaped glass

Thin film deposition

Magnetron sputtering

Magnetron sputtering was the main technique used in this work to prepare thin films. Magnetron sputtering is a physical thin film deposition method, wherein a microwave generator is used to transform a gas into a plasma of positively charged ionic species, which are accelerated by an electric field towards a cathode material (target), which will be sputtered as a consequence of the ion bombardment. The sputtered species will then move away from the cathode and may condense on a substrate which was placed underneath the sputter target. Figure 3.5a shows how a typical magnetron sputtering set up looks like.

A series of gases may be used as sputter gas, depending on the target material and the desired thin film composition. Argon is often used as the main magnetron sputtering gas, because of its inertness in non-ionic state and its appropriate mass which allows the sputtering of many compounds. Oxygen is regularly added to the sputter gas when oxide-containing targets are used.

Because the sputter process frequently causes the preferential sputtering of lighter elements, oxide containing targets may become slowly deprived of oxygen. The addition of oxygen to the sputter gas may then counter the oxygen deprivation of the target. Thus, preventing the deposition of non stoichiometric oxide thin films. This sputtering process is also called semireactive sputtering, because part of the oxide in the deposited thin films comes from the sputter gas. Also reactive sputtering may be performed, whereby all oxygen in an oxide thin film, for instance, originates from the sputter gas.

The magnetron sputtering set up consists of an inner part that functions as the cathode and an outer part that is regarded as the anode. Between those two parts the electric field is applied.

As the anode should have a larger area than the cathode, an electrical contact between the former is that the continuous bombardment of the target by cations cannot be sufficiently compensated by electrons, so that the negative potential of the cathode gradually decreases, until the sputter process eventually stops. For weak conducting compounds, an AC signal would be a better choice. Herein, an alternating electric field is generated at a frequency for which no impedance is created at the targets. This requirement is fulfilled for a frequency of 13.56 MHz and has, thus, been chosen as the standard frequency. Because this frequency lies in the radio frequency regime, the technique is also called radio frequency (RF) sputtering. At the beginning of an RF sputtering process, the AC signal is symmetric around zero. However, this changes over time. Due to the alternating electric field, both ions and electrons will be accelerated towards the sputter target.

Because the electrons have a lower mass compared to the cations, more electrons than cations will reach the target surface, so that an electron excess forms at the target surface over time.

As a consequence, the potential of the target diminishes, allowing the cations to reach the target surface at a steady pace, a mechanism called "self-biasing". In a common RF magnetron a circular magnet array is located above the target. The magnetic field functions as a sort of trap to raise the electron density around the north pole, which as a result causes an increased ionization rate in this area and, thus, a higher sputter rate. Due to the target being consumed at a higher rate in this circular zone, it is also commonly denoted as a "race track".

The DC and RF magnetrons used in this work were supplied by Thin Film Consulting and the RF signal was generated by a Hüttinger PFG 300RF RF generator.

Contact films for interface experiments

All interface experiments, for which the exact procedure will be explained further in this Chapter, were performed in the DArmstadt Integrated SYstem for MATerials science (DAISY-MAT) (Figure 3.6). The DAISY-MAT is a system which include chambers for surface cleaning, thin film deposition and surface analysis connected to each other via a distribution chamber operated under ultra-high vacuum (UHV). Because the samples are transferred in vacuum, surface contamination in between thin film deposition and surface analysis is limited. Substrates can be introduced into the DAISY-MAT through a load lock, which can be pumped from atmospheric pressure down to high vacuum.

The interface experiments required the deposition of contact thin films, whereby the thin film deposition was performed through magnetron sputtering. The deposition parameters of the investigated contact materials, RuO 2 , NiO, CoO x and Sn-doped In 2 O 3 (ITO), can be found in Table 3.1. RuO 2 and NiO were reactively sputtered from a metallic Ru target (Lesker) and metallic Ni target (Alineason), respectively. CoO x was semireactively sputtered from a ceramic Co 3 O 4 target (Alineason). ITO was sputtered from a ceramic ITO target with 10% Sn (Lesker). All targets had a diameter of 2 inches. For the ITO(400 ℃) deposition, heating of the substrates was performed in the process gas atmosphere through a halogen lamp, whereby the temperature was measured through a thermoelement. 10 ℃/min was the approximate heating and cooling rate. Adapted from [4].

BiVO 4 thin films

The polycrystalline gradient BiVO 4 thin films were prepared by Fuxian Wang from the group of our collaboration partner Dr. Klaus Ellmer of the Helmholtz-Zentrum Berlin through reactive magnetron sputtering using a home-built magnetron sputtering system equipped with a load lock. 5 A co-sputtering set up was used with two separate 2 inches diameter metallic Bi (99.9%) and V (99.99%) targets, inclined by 30°to each other and focused onto a soda lime glass substrate with a size of 22 x 22 mm 2 (Figure 3.7a). The glass substrates were coated with a highly conductive and transparent fluoride doped tin oxide F:SnO 2 film, which served as back contact during the XPS measurements. The sputtering atmosphere was an Ar/O 2 gas mixture (20% O 2 ) at a total sputtering pressure of 1 Pa. The DC power at the Bi target was about 20 -30 W, while the DC power at the V target was 450 W. While the V-target-to-substrate distance was 6 cm, the Bi-target-to-substrate distance was 12 cm. Since no rotation stage was used during sputtering, every sample contained a Bi/V compositional gradient along the sample width. After sputtering, the films were annealed in a muffle furnace at 500 ℃ in air for 2 h, using a temperature ramp of

CuFeO 2 pellet preparation

CuFeO 2 pellets were prepared from the CuFeO 2 powder (synthesis procedure in Section 3.1.3), which was not deemed fit for the photodeposition experiments (more thorough explanation in Section 6.5.1). 600 -700 mg of powder was first homogeneously ground with a pestle and mortar.

Then, the powder was uniaxially pressed at room temperature using an hydraulic powder press (CA NC II OSTERWALDER AG) at 300 bar for 15 min. Next, the pellet was further compacted at room temperature with a cold isostatic press (KIP 100E, Paul Otto Weber, Remshalden) at 5000 bar for 30 min. The result was a sturdy pellet with a thickness of 0.5 -1 cm.

Characterization

UV-Vis diffuse reflectance spectroscopy

To assess the optical properties of the particulate photoabsorbers and their heterostructures, UV-Vis-NIR diffuse reflectance spectroscopy was carried out. Herein, a light beam is focused onto a powder sample, which is commonly loaded into a cuvette, and the diffuse scattered light is collected, mostly using an integrating sphere. Normal UV-Vis-NIR spectroscopy, which is based on the transmission of light cannot be applied for powdered samples due their opaqueness.

For semiconductors the main electronic excitations induced by light concern the transition of electrons from valence band to conduction band and are governed by the band gap of the semiconductor. Most semiconductors have band gap energies between 1 eV and 4 eV. By applying the following equation:

λ = hc E (3.1)
with E as the excitation energy, λ as the wavelength, h as the Planck constant and c as the speed of light, one can calculate that band gap optical transitions are induced by photons with wavelengths between 300 -1240 nm, thus photons in the UV-Vis-NIR region of the light spectrum. For a semiconductor with a certain band gap, photons with energies below the band gap will pass through, whereas photons with an energy higher than the band gap may be absorbed if the electronic selection rules are fulfilled.

In this work, UV-Vis-NIR diffuse reflectance spectroscopy was performed with a Lambda 900 spectrometer (Perkin Elmer), equipped with an Ulbricht integrating sphere. Diffuse reflectance spectra were measured for a wavelength range of 300 -800 nm and a wavelength range of 300 -1200 nm for the BiVO 4 and CuFeO 2 powder samples, respectively. All samples were loaded into a glass cuvette with a pathlength of 1 mm. BaSO 4 powder was used as a reference.

The measured reflectance R was transformed into the Kubelka-Munk function F(R) according to:

F(R) = (1 -R) 2 2R (3.2)
whereby the Kubelka-Munk function is equivalent to the absorption coefficient according to the Kubelka-Munk theory: 7,8

F(R) = α s (3.3)
with α as the absorption coefficient and s as the scattering coefficient. respectively. Due to the equivalence of α and F(R) the optical band gap E g of a semiconductor can be estimated from:

F(R)hν = A(hν -E g ) n (3.4)
with A being a constant and ν being the light frequency. This equation was first introduced by Tauc and is therefore also referred to as Tauc's equation. 9 n relates to the type of optical transition taking place in the semiconductor. Specifically, n = 1/2 and n = 2 refer to the direct and indirect allowed optical transition in the semiconductor. Thus, creating plots of (F(R)hν ) 2 vs hν and (F(R)hν) 1 2 vs hν allows for an estimation of the direct and indirect optical band gap, respectively, from the intercepts with the x-axis of the extrapolated linear part of these plots.

Raman spectroscopy

Raman spectroscopy is a spectroscopic technique used to obtain chemical and structural information of a sample through the recording of inelastic scattered light of a monochromatic light source. Next to the prevalent elastic scattering (Rayleigh scattering), photons can inelastically scatter (Raman scattering) due to interactions with molecular vibrations or phonons. The energy shifts of the Raman scattered photons (Raman shifts) depend on the energy levels of the functional group vibration they interact with. Thus, Raman spectrograms of scattered light intensity versus Raman shift yield vibrational information that can be used to identify molecules and crystalline phases. Raman spectroscopy is complimentary to IR spectroscopy in the sense that Raman scattering only takes place after interaction with vibrations that change the (molecular) polarizability, while IR absorption bands arise from interactions with vibrations that induce a change in the overall dipole moment.

In this work, a LabRAM HR 800 Raman spectrometer was used to record the Raman spectra of all thin film and powder samples. An argon ion laser was used as an excitation source operating at a wavelength of 488 nm and an approximate laser power of 10 mW on the sample. A 10X objective was used to generate a spot size of approximately 2 µm. Before each set of measurements, the Raman shift wavenumber was calibrated by measuring a silicon substrate and setting the peak position of the main peak to the expected value of 521 cm -1 .

X-ray diffraction

Powder X-ray diffraction (PXRD)

Powder diffraction through X-rays is a popular method for detecting the crystalline phases in a powder sample. When X-ray photons hit a crystalline lattice they are diffracted into specific directions due to the periodical ordering of the atoms in the lattice. This phenomenon is called Bragg diffraction and is a consequence of the interference between waves reflecting from differ-ent crystal planes. The condition for constructive interference is summarized in Bragg's law:

2dsinθ = nλ (3.5)
where λ is the incident X-ray wavelength, d is the crystal plane separation and θ is the incidence angle. Typically, the Bragg-Brentano θ /2θ configuration is used to measure powder samples and consists of a fixed X-ray source, a sample stage that is tilted by an angle θ and a detector rotating around the same rotation axis by 2θ . In this configuration the lattice planes parallel to the sample stage surface are probed and for a homogeneously mixed powder sample all the main lattice planes should be registered. The set of recorded diffraction angles and relative intensities allows the determination of the crystalline phases, present in the sample, by comparison with diffractograms from databases.

The θ /2θ diffractograms of the powder samples were commonly recorded by a Bruker D2 phaser with a Cu Kα X-ray source (λ = 1.5418 Å), continuously scanning from 10°to 90°. The powder samples were loaded into small disc-like silicon containers. The diffractograms of the CuFeO 2 samples, however, exhibited a strong X-ray fluorescence background due to the iron atoms absorbing the Cu Kα X-ray photons. To avoid X-ray fluorescence and to obtain higher resolution θ /2θ diffractograms, a PANalitycal X'pert MPD-PRO diffractometer was used, which was equipped with a Cu Kα X-ray anode (operated at 45 kV and 40 mA), a secondary monochromator and an X'celerator detector. The secondary monochromator allowed to filter out the X-ray fluorescence signal generated by the CuFeO 2 samples. Herein, the powder samples were placed onto silicon wafer sample holders and flattened with a piece of glass.

Rietveld refinement

Rietveld refinement is a more thorough analysis of X-ray powder diffractograms, which is able to determine the amorphous fraction, give an estimate of the crystallite size, allow a more precise determination of the crystalline phases and provide the relative quantity of each crystalline phase. In short, a Rietveld refinement uses a least squares procedure, whereby a theoretical line profile is refined until it matches the measured profile. To begin the refinement, a set of starting parameters need to be given, which include the lattice parameters and crystalline structure of the estimated crystalline phases and the estimated size and morphology of the crystallites. Rietveld analysis was, herein, used to determine whether other crystalline phases were present besides monoclinic clinobisvanite in the BiVO 4 powder samples.

The θ /2θ diffractograms of the powder samples used for the Rietveld analysis were measured with a STOE STADI P diffractometer, equipped with a Ge 111 monochromator, a Mo Kα 1 Xray source and a position sensitive detector. Rietveld analysis was carried out with the Fullprof software suite. Figure 3.8: Schematic of the interaction of a primary electron beam with a solid state sample, including the escape depth of all measurable secondary signals. 10 Secondary electrons form the main signal in SEM. Secondary electrons are "free" electrons that were ejected from their atomic shells, after primary electrons, X-rays or other secondary electrons transferred enough energy to overcome the atomic binding energy of bound electrons.

Because of their low kinetic energy, most secondary electrons only escape the sample if they were generated at a sample depth of a 1 -10 nm (Figure 2.28). Additionally, the number of secondary electrons that escape the sample depends on the angle of incidence of the primary beam with the sample surface. Generally, more secondary electrons are released as the curvature and slope of the sample increases. Thus, by scanning the primary beam in a raster-like pattern along the surface and detecting how many secondary electrons are released, a composite topological image of the surface can be created. This technique is called secondary electron imaging (SEI). Image resolutions of less than 0.5 nm are feasible.

As a result of the ejection of core electrons by the primary beam, X-rays can be generated if outer shell electrons migrate to inner shell vacancies. These X-rays have an energy equivalent to the energy difference between the two involved atomic energy levels. Thus, each element correspond to a specific set of "characteristic" X-rays. Consequently, the energy analysis of X-rays emitted from a sample after electron beam bombardment, allows obtaining chemical information of that sample. This X-ray analysis technique is called energy-dispersive X-ray spectroscopy (EDS). The characteristic X-rays are typically detected up to a depth of a few micrometers, due to the strong penetration of the primary electron beam and the high escape depth of X-rays.

The lateral resolution varies from 50 to a few 100 nm, depending on the primary electron beam energy and the type of detector. Besides X-rays, Auger electrons can be released when electrons recombine with core holes. These Auger electrons are also element characteristic and have small escape depths of a few angstrom (Figure 3.8) due to their limited inelastic mean free path (Figure 2.28). Therefore, scanning Auger microscopy (SAM) is much more surface sensitive than EDS.

The primary beam electrons can also elastically scatter on the atoms in the sample and when they leave the sample again they are called backscattered electrons. Due to an approximate square root dependence of the scatter probability on the atomic number, elemental contrast can be achieved if the atoms in the sample have sufficiently different atomic numbers. Since the kinetic energy of backscattered electrons barely changes with respect to the primary electrons, the escape depth reaches up to a few 100 nm. Consequently, backscattered electrons yield a lower lateral resolution with respect to secondary electrons.

A Hitachi TM 1000 was used to make low resolution images (up to 10.000 magnification) through SEI of the BiVO 4 and CuFeO 2 powder samples. A JEOL 6700F, equipped with a field emission gun and detectors for SEI, backscattered electrons (BSE) and EDS, was used to make higher resolution images (up to ± 700.000 magnification) and to perform chemical microanalysis of the (heterostructured) powder samples. SAM was performed with a PHI 710 scanning Auger nanoprobe, having a maximal lateral resolution of 10 nm and a maximal SEI resolution of 3 nm, to elucidate how the contact materials of the heterostructured powder samples were distributed along the particulate photoabsorbers. Different types of measuring modes were used, including point measurements (Auger spectra acquisition on a certain point), zone measurements (Auger spectra acquisition on a certain area) and mapping (Auger intensity acquisition of a certain area by scanning the area in a raster like pattern). Also SAM depth profiling was occasionally tried, whereby an Argon ion gun (sputter rate 10 nm/min) was used to gradually sputter a sample and Auger spectra were measured after each sputter step.

Transmission electron spectroscopy (TEM)

In transmission electron microscopy (TEM) an electron beam is focused onto a sample, just as in SEM. However, not the secondary electrons are captured, but the electrons which pass through the sample are. Therefore, samples in TEM are mostly ultrathin specimens, below 100 nm, or nanocrystalline powders suspended onto a conductive grid which allow electron transmission.

By capturing the transmitted electrons (bright field) or the diffracted electrons (dark field) by a fluorescent screen or a charge-coupled device (CCD) camera, an image can be made from the sample. Due to the small de Broglie wavelength of electrons, resolutions up to the atom scale can be obtained. Besides imaging, TEM allows the registration of selected area electron diffraction (SAED) patterns for crystalline samples. SAED patterns are measured by placing a selected area aperture into the electron beam path after it passes through the sample. In this way, information on the crystalline structure, crystallinity and texture of a certain area of the sample can be collected. Also, EDS may be performed with TEM, and due to TEM samples being relatively thin, the electron interaction volume will be smaller than for SEM (See Figure 3.8). Additionally, to perform EDS mostly a scanning transmission electron microscope is used, which differs from a conventional TEM, in that it is able to focus the electron beam on a very fine spot (with a size between 0.05 -0.2 nm) on the sample. Therefore, EDS with a TEM apparatus allows for a higher lateral resolution compared to SEM-EDS. Additionally to the aforementioned uses of TEM, several other analyses, including electron energy loss spectroscopy (EELS) and high-angle annular dark-field (HAADF) imaging can be performed depending on the chemical and structural nature of the samples. For further reading on the technical principles behind TEM and its applicability in materials science, I refer the reader to the book of Williams and Carter. 11 TEM measurements, in this work, were performed with a JEOL 2100F, equipped with a field emission gun (FEG) and an Oxford Instruments EDS system. The TEM was operated at an acceleration voltage of 200 kV. The TEM powder samples were prepared by first making a suspension of the powder in ethanol and then putting a small droplet of the suspension onto a holey carboncoated copper grid. After letting the sample dry, a thin carbon layer was sputtered onto the sample to avoid charging under the electron beam. Due to the relatively weak stability of the BiVO 4 powder samples under a focused electron beam, the electron beam was slightly defocused to reduce the intensity of the electron beam hitting the samples.

Photoelectron spectroscopy (PES)

In Chapter 2 in Section 2.4 an overview is given of the fundamentals of photoelectron spectroscopy as well as an explanation on how the core level photoelectron spectra, which will be discussed in this work, may be interpreted. In this Section the way how the photoelectron spectra were measured as well as how data were extracted from the photoelectron spectra is summarised.

All photoelectron spectra in this work were measured with the Physical Electronic PHI 5700 multi-technique surface analysis system, which is part of the DAISY-MAT system (described in Section 3.2.2). The photoelectron spectroscopy chamber is operated at about 10 High resolution XP core level spectra contains plenty of information. First of all, the overall shape of the core level spectrum needs to be inspected, as the appearance of secondary features in the form of shoulders or additional peaks can point to an altered chemical composition or the presence of adsorbates. Next, the binding energy of the analysed core level can be determined from the maximum line intensity. Herein, a fitting function was used, which assumed a Gaussian line shape around the maximum of the core electron signal, which is a valid assumption for straightforward core level spectra, such as Bi4f, V2p 3/2 , O1s and Cu2p 3/2 core level spectra. For more complicated spectra, such as Fe2p 3/2 , Co2p 3/2 and Ni2p 3/2 , the maximum of the core level spectra had to be determined manually by drawing a straight line towards the maximum of intensity.

Additionally, the chemical surface composition can be determined from the XP core level spectra. The surface ratio x A/B of two elements A and B can be calculated according to the following formula: The intensities in Equation 3.6 can be determined by integrating the corresponding high resolution core level spectra, after removing the inelastic scattered electron background from the core level spectra. Polynomials could be used to remove the background, but that would not be correct from a physical point of view. Better is to use the more physically correct photoelectron background removal procedure introduced by Shirley in 1972 12 or the one described by Tougaard and Jansson in 1992. 13 Also, for difference spectra, whereby the difference is taken of spectra from the same core level region to inspect minor chemical differences, background subtraction had to be performed. In this work, the Bi/V surface ratio was determined frequently by taking the ratio of the integrated Bi4f intensity and V2p 3/2 ratio. However, as will be shown in Chapter 5, a correction factor has to be applied to obtain the real Bi/V surface ratio, due to the tabulated sensitivity factors not correcting adequately for the different photoelectron generation cross sections of Bi4f and V2p 3/2 . In addition, difference spectra for the Bi4f, V2p 3/2 , Cu2p 3/2 and Fe2p region were occasionally created. Before making the difference spectra, the background of the V2p 3/2 , O1s and Cu2p 3/2 spectra were removed with a Shirley function, whereas the background of the Bi4f and Fe2p spectra were removed with a Tougaard function. Figure 3.10 shows the effect of the Shirley and the Tougaard background removal functions on a high resolution Bi4f spectrum.

The differences between both functions are minor, however, the Tougaard function follows the curvature of the Bi4f spectrum in between the Bi4f 7/2 and Bi4f Binding energy (eV) Binding energy (eV)

φ = hν -E SEC -φ bias (3.7)
with hν corresponding to the photon energy, which for UPS is mostly 21.2 eV (He I) and φ bias being the applied bias (here -4 V). The onset of the valence band region in UPS (right most region) correspond to the valence band maximum E UPS VBM and can be determined (Figure 3.11b) by extrapolating the linear part of the valence band region to the lower part of the low binding energy side. In a similar way, the valence band maximum (E XPS VBM ) can be determined from the valence band region in XPS (Figure 3.11c). Unless stated otherwise, VBMs were collected in this work from the XPS valence band region. The main reason is that the UPS valence band regions of CuFeO 2 and BiVO 4 at the low binding energy side appeared to be too rounded, so that a linear extrapolation would not be reliable.

Interface experiments

Interface experiments in this work were performed to determine the junction properties between photoabsorbers (BiVO 4 and CuFeO 2 ) and selected contact materials (ITO, RuO 2 , NiO and CoO x ).

The preparation of the BiVO 4 thin film substrates and CuFeO 2 pellets is described in Section 3.2.3 and Section 3.3, respectively. The contact materials were sputtered onto the substrates through magnetron sputtering (See Section 3.2.1) using the parameters listed in Table 3.1.

In a particular interface experiment, the surface of a substrate was first cleaned with distilled water and, subsequently, dried using compressed air to remove any dust attached to the substrate surface. Then, the substrate was mounted onto a stainless steel substrate holder and kept in place using screws and washers. Next, the sample was introduced into the DAISY-MAT system (described in Section 3.2.2) through a load lock to move from ambient pressure to vacuum (10 -7 -10 -9 mbar). After that the substrate surface was subjected to a further cleaning step to remove water and carbonaceous species adsorbed onto the sample surface. For the CuFeO 2 pellets, this cleaning step represented a heating step whereby the samples were heated for at least 3 h in an oxygen atmosphere (flow: 5 sccm; 4 x 10 -4 mbar) at a temperature of 200 ℃. The BiVO 4 were cleaned with a more rapid oxygen plasma treatment. The working principle of an oxygen plasma source can be found in the paper of Anton et al. 14 In this work, the oxygen plasma treatment was carried out with a GenII Atom/Ion Hybrid Plasma source by Tectra. This plasma source generated an atom flux of 2 x 10 16 atoms cm -2 s -1 at the working distance of 10 cm. For the oxygen plasma treatments, an oxygen flow of 2 sccm was used, which resulted in a working pressure of approximately 1 x 10 -4 mbar. A 40 mA plasma current was maintained for 15 minutes.

The plasma was operated at a power of about 180 W. After cleaning the sample surface, a first set of PES measurements (description in Section 3.4.4) was performed to assess the chemical and electronic properties of the bare substrate. After that a contact thin film was deposited onto the substrate in a stepwise fashion, whereby PES measurements were performed after each deposition step to determine how the substrate core level spectra change as a response to the growing film. In the last deposition step, a relatively thick layer of the contact material was deposited to measure the electronic and chemical properties of the contact film by PES. The whole procedure of an interface experiment is also schematically represented in Figure 3.12a. To represent the interface experiments in this work, all relevant core level spectra and valence band spectra of substrate and film will be stacked above each other in the same figure, whereby the stacking order from bottom to top represents the subsequent deposition steps.

All the PES data measured during an interface experiment could then be used to interpret the interfacial band alignment between the two materials that form the contact. For the bare Figure 3.12: (a) Representation of an interface experiment, (b) Evolution of the binding energy with film thickness for the core levels, VBM and CBM of the substrate and the deposited contact film, (c) Evolution of the VBM of the substrate and the film with respect to the Fermi level position as a function of the film thickness. Adapted from [15].

substrate, first, the VBM (E VB,0 (sub)) and the energetic distance between core levels and VBM (E VB CL,0 (sub)) would be extracted. Then, the UPS measurement performed on the bare substrate was used to determine the work function of the bare substrate. Next, all substrate and film core level binding energies after each deposition step were determined. Due to the relatively low film thickness after the first few deposition steps, the substrate core photoelectrons could still be detected, thus, allowing the determination of the substrate related core level binding energies.

After a relatively thick film had been deposited, the substrate core level photoelectrons could not penetrate through the film any more. After this observation, a final deposition was made to obtain a sufficiently thick film so that the VBM (E VB,final (film)) and the energetic distance between core levels and VBM (E VB CL,final (film) of the thick film could be extracted from the PES measurements on this thick contact layer. Additionally, the work function of the contact layer would be determined from a UPS measurement taken of the contact film. When the chemical state of the elements in substrate and film did not change during the interface experiment, the distance between VBM and core levels was considered to be constant for both the substrate and the film. It should be noted that in most cases the VBM of the substrate cannot be determined any more after the first contact film deposition step, because of the overlap of the contact film valence band spectrum with that of the substrate. In Figure 3.12b it is shown that the core levels and VBMs of substrate and film change in parallel with increasing contact film thickness. These data can then be used to determine how the VBM and core level binding energies shift with respect to the Fermi level (Figure 3.12c), which allows to construct a band alignment diagram between substrate and film to visualize phenomena like band bending and Schottky type barriers. Additionally, the shape of the core level spectra should be checked to analyse whether any chemical changes take place during the interface experiment.

Additionally, for the BiVO 4 substrates, the thickness of the deposited contact films could be estimated from the difference in attenuation of the Bi4f and V2p 3/2 core level intensity. Since the kinetic energy of the photoelectrons from the Bi4f emission is higher than for the V2p 3/2 emission, the Bi4f photoelectrons are attenuated less because of the higher inelastic mean free path of these photoelectrons. The thickness of a thin film growing on top of the BiVO 4 substrate can then be evaluated, under the assumption that the surface V/Bi ratio does not change during film growth, according to following formula:

t = ln ( I V x I Bi,0 I B i x I V,0 ) x λ V x λ Bi λ V -λ Bi (3.8)
where t: film thickness; I V : integrated V2p 3/2 core level intensity at film thickness t; I Bi : integrated Bi4f core level intensity at film thickness t; I V,0 : integrated V2p 3/2 core level intensity at zero film thickness; I Bi,0 : integrated Bi4f core level intensity at zero film thickness; λ V : electron effective attenuation length of V2p 3/2 photoelectrons; λ Bi : electron effective attenuation length of Bi4f photoelectrons. The electron effective attenuation lengths were calculated using the "NIST Electron Effective-Attenuation-Length Database", based on the density, number of valence electrons and band gap of the deposited film.

Water exposure experiments

In addition to normal interface experiments, where a contact thin film was gradually sputtered onto a substrate, a special type of interface experiments were performed, whereby a water layer was adsorbed onto the substrate. These water exposure experiments were just like the other interface experiments carried out to determine how the substrate binding energy would shift after a water layer was adsorbed onto the substrates. In addition, also the chemical changes related to the water exposure could be interpreted through PES measurements.

In a typical water exposure experiment, a cleaned substrate was first placed into the atomic layer deposition (ALD) chamber in the DAISY-MAT. The ALD chamber is one of the chambers in the DAISY-MAT, which is mainly used to deposit Al 2 O 3 thin films via ALD, but was used in this work solely for the water exposure experiments. Onto the ALD chamber, a closed glass vessel containing distilled water was attached via an ALD valve. The ALD valve forms the connection between the glass vessel, at ambient pressure, and the chamber, under high vacuum. By opening the ALD valve for a short period of time, water vapour flows into the chamber and may condense onto the substrate surface. The complete water exposure experiment followed several cycles, whereby the ALD valve of the water reservoir was opened for a certain duration, followed by a period of pumping down the chamber. In our experiments, we performed 15 cycles with an exposure duration of 100 ms and a pumping down duration of 300000 ms. Due to the pumping down to vacuum no physisorbed water remains on the substrate surface. Before and after the water exposure, PES measurements of the substrate were taken.

The closed gas vessel contains two arms which are needed when the water in the reservoir is replenished and needs to be degassed. For the degassing, the water placed in one arm of the reservoir is frozen with liquid nitrogen, after which the ALD valve is opened for a few seconds to create a vacuum in the glass reservoir. Then, the other arm is cooled with liquid nitrogen while the arm with the frozen water is heated. In this way, the water evaporates in the one arm and condenses in the other. Thereafter, the ALD valve is opened again for a few seconds. This procedure was repeated 4 times in total.

The ALD valve used, was a Swagelok 6LVV-ALD3FR4-P-C valve, which was pneumatically controlled with a MAC 34C-ABA-GDFC-1KT controller. For a complete overview of the ALD chamber I refer the reader to the work of Mirko Weidner. 16

Water desorption experiments

Water desorption experiments were carried out in the Darmstadt Integrated System for Fundamental research (DAISY-FUN), 17 which consists of several chambers operated under UHV and dedicated to sample preparation and/or surface analysis (Figure 3.13). The characteristics of DAISY-FUN are described in detail elsewhere. 18 For a water desorption experiment, the BiVO 4 thin film substrate was first of all thoroughly rinsed with distilled water and dried with compressed air. Then, the substrate was mounted onto an Omicron sample holder and inserted via a load lock into the DAISY-FUN. Subsequently, a calcination step was performed to remove the atmospheric surface contaminations. The calcination was performed for 5 h at 400 ℃ under an O 2 atmosphere (10 -3 mbar).

Afterwards, the BiVO 4 sample was placed into the XPS-chamber, operated at a base pressure of 1 x 10 -9 mbar. The XPS-chamber is equipped with a SPECS Phoibos 150 analyser and a SPECS Focus 500 X-ray source to perform XPS measurements using the monochromatized AlKα line Additionally, a glass water vessel is mounted onto a small pipe connected to the XPS-chamber.

This water vessel contains ultra-pure water, which was degassed through the freeze-pump cycling procedure explained above. The dosage of water was precisely controlled by a leak valve and the exposure time was varied stepwise to match the desired dosage using uncorrected pressure readings.

For the water desorption experiment, the BiVO 4 sample was first cooled to 77 K using liquid nitrogen (LN). Then, frozen water was stepwise deposited onto the BiVO 4 sample by opening the leak valve connected to the water vessel for specific amounts of time. After the deposition of 10 -15 langmuir of water, the sample was slowly heated up to room temperature. XPS and UPS measurements were performed before cooling down, after cooling down, in between each deposition and after heating up the sample.

Peak fitting

For the V2p The fitting parameters for V 5+ and V 4+ correspond quite well with what was found earlier for vanadium oxides. 19,20 Namely, that for V 4+ the L/G decreases and the full width at half maximum (FWHM) increases with respect to V 5+ . The 0.76 eV binding energy difference between V 5+ and V 4+ is, however, lower than the 1.2 eV difference found for vanadium oxides. The higher FWHM of the second signal in the O1s spectra measured during the water desoprtion experiment in DAISY-FUN is due to the presence of carbonaceous species which have a similar binding energy as hydroxyl groups. 21

Photocatalytic efficiency assessment

Rhodamine B degradation

As a first test of the photocatalytic efficiency, the particulate BiVO 4 samples were tested for their dye degradation performance, where rhodamine B (Rhd B) was used as the dye. Rhd B was not chosen for any particular reason besides that the wavelength of maximum absorbance λ max (554 nm), is higher than the band gap of BiVO 4 (516 nm). The dye degradation experiments were performed with the same set up as the one used for the photodeposition experiments (Figure 3.3). All experiments were performed under ambient pressure and temperature. In a typical dye degradation experiment, 50 mg of the photocatalyst powder was loaded into the honeypot shaped glass beaker together with 100 mL of a Rhd B solution (10 mg/L in distilled water). The beaker was covered with parafilm and placed for 10 min in an ultrasonic bath to disperse the powdered photocatalyst. After that, the top of the glass vessel was covered with a loose transparent Petri dish and the sides of the glass vessel were covered with aluminium foil. Then, a 1.2 mL aliquot of the dispersion was taken. Subsequently, the dispersion was stirred at 750 RPM in the dark for 30 min, after which another 1.2 mL aliquot was taken to determine whether the dye already degraded in the dark. At that point the shutter of the solar simulator, which was allowed to warm up for 30 min, was opened and the dispersion was illuminated. Then at time intervals of 30 min, 1.2 mL aliquots were take from the dispersion during a period of 2 h. Next, all recuperated aliquots (normally 6 in total) were centrifuged to separate the photocatalyst powder from the Rhd B solution. For each aliquot, 1 mL was taken, diluted to 3 mL with distilled water, and loaded into a quartz cuvette with a 1 cm path length. Thereafter, the cuvette was placed into a UV-Vis spectrophotometer (Shimadzu, UV-1650PC) and a UV-Vis spectrum in the range from 200 nm to 800 nm was recorded, as well as the absorbance at λ max . Then, absorbance plots as a function of degradation time were made to visualize how the absorbance and, thus, the concentration of the undegraded dye decreases. Also, blank experiments in the absence of photocatalyst were performed. All experiments were at least repeated once to assess the reproducibility of the experiments.

Sacrificial photochemical water splitting

To test the photochemical water splitting performance of BiVO After stirring for 10 min, the light source was uncovered and the sacrificial water oxidation/reduction was initiated. The amount of evolved gases were continuously determined by the automatic gas burette, connected to a pressure sensor. After stopping the reaction, a 1 mL gas sample was taken and analysed with a GC (Agilent GC HP 6890N) equipped with a carboxen 1000 column and a thermal conductivity detector (TCD), having argon as carrier gas. the nickel contact layer was deposited through photodeposition or electrodeposition. [4][5][6] In this Chapter, first, the experimental parameter optimization for the BiVO 4 synthesis and its twofold upscaling will be detailed, whereby the effect of each parameter on the morphology and crystalline phase composition will be shown. Next, the structural and chemical properties of the BiVO 4 /Pt, BiVO 4 /Ag, BiVO 4 /CoO x OH y and the novel BiVO 4 /NiO x OH y heterostructured powders will be described, which were manufactured through photodeposition. Herein, the emphasis is laid on analysis techniques, which helped to demonstrate whether the metal(oxide) deposits that adhered onto the BiVO 4 microcrystalline particles exhibit a certain regioselectivity.

Finally, the newly synthesized BiVO 4 /NiO x OH y heterostructure will be compared to the other heterostructured BiVO 4 powders with regards to its ability for dye degradation and sacrificial water oxidation.

Synthesis of monoclinic BiVO 4 microcrystals

To be able to explore new anisotropic BiVO 4 heterostructures, truncated bipyramidal BiVO 4 with exposed {010} and {110} crystal facets had to be synthesised in large quantities first. According to literature several synthesis procedures are available to obtain the truncated bipyramidal morphology (see Table 4.1). However, none of the published synthesis procedures, we tried, resulted straightforwardly in monodisperse monoclinic BiVO 4 powders with the truncated bipyramidal morphology. This specific morphology could only be obtained after tuning several reaction parameters. In Section 3.1.2 the optimised synthesis procedure is explained in detail and in what follows the way towards the optimised synthesis conditions is described.

Most published synthesis procedures make use of Bi(NO 3 ) 3 and NH 4 VO 3 as the bismuth and vanadium precursor, respectively, whereby stoichiometric amounts of both precursors are first dissolved and then thoroughly mixed to form an aqueous reaction mixture that is subsequently subjected to a hydrothermal treatment, after which monoclinic BiVO 4 microcrystals with a truncated bipyramidal shape are obtained. As can be seen in Table 4.1, there seems to be some flexibility in the reaction parameters since the presence of a surfactant is not necessarily required, the reaction mixture which is put in the autoclave can be acidic or neutral, the autoclave temperature can be varied from 160 ℃ -200 ℃ and even an autoclave treatment does not appear to be compulsory.

For a start, the synthesis procedure mentioned in the first reference in Table 4.1 was tried, thus, with sodium dodecyl sulphate (SDS) as surfactant, with a reaction mixture brought to pH 7, at an autoclave temperature of 180 ℃ and with an autoclave time of 24 h. 7 These parameters will be referred to as the baseline parameters throughout this Section. Through this synthesis procedure, monoclinic BiVO 4 was obtained, as evidenced by the X-ray powder diffractogram in Figure 4.1a. Indeed, the diffraction pattern corresponds exclusively to monoclinic BiVO 4 . 16 However, the morphology of the obtained powder seems quite irregular with no indication of truncated bipyramidal shaped crystals (Figure 4.1b). Therefore, a series of syntheses were performed, whereby the parameters of the baseline experiment were varied one by one to figure out which ones affected the BiVO 4 morphology. Specifically, the pH, autoclave temperature, autoclave time, autoclave stirring and presence of surfactant was changed, since previous research pointed out that these parameters had an influence on the appearance of the BiVO 4 particles. 17,18 In Figure 4.2 the recorded secondary electron images show the effect of the pH on the morphology. For a pH between 3 and 7, the resulting BiVO 4 particles have mostly irregular shapes

V 2 O 5 - 0 - - [14] Bi 2 O 3 V 2 O 5 - 0 - - [15]
EDTA: ethylenediaminetetraacetic acid All in all, bringing the reaction mixture pH to pH 1.5 and increasing the autoclave reaction time to 64 h seem to be adequate measures to obtain BiVO 4 with a truncated bipyramidal morphology. However, to ensure that those measures lead systematically to truncated bipyramidal BiVO 4 particles, several reproductions of those synthesis conditions had to be performed. Therefore, the procedures with a pH of 1.5 and an 64 h autoclave time reaction conditions were repeated several times. However, only the adjustment of the pH to pH 1.5 led to reproducible results over multiple syntheses. As can be seen in Figure 4.8, all powders synthesized at pH 1.5, are purely monoclinic BiVO 4 and consist of particles with clearly exposed crystalline facets. Next, the synthesis procedure had to be upscaled, because only ± 300 mg of BiVO 4 powder was obtained after each synthesis, while we estimated that a batch of 10 g was needed to perform all the planned photodeposition experiments. Multiplying the amount of every reactant by two, however, did not result in the preferred particle characteristics any more. Indeed, Figure 4.9 shows that the powder looks to be more aggregated and irregularly shaped. Non-aggregated homogeneously shaped BiVO 4 particles could be obtained through two modifications to the synthesis procedure; namely by multiplying the amount of SDS by three instead of two and by stirring the reaction mixture during the hydrothermal treatment in the autoclave.

The fully optimised BiVO 4 synthesis procedure yields ± 600 mg and was repeated 17 times to establish a 10 g batch of truncated bipyramidal BiVO 4 for the photodeposition experiments.

Photocatalysis

The observation that different unexpected morphologies were obtained during the truncated bipyramidal BiVO 4 experimental parameter optimization was welcomed to explore the characteristics and oxidation performance of several morphologies further. In particular, the irregular, dendritic and truncated bipyramidal shaped BiVO 4 powders were studied further, and compared with each other and to a commercial BiVO 4 powder purchased from Alfa Aesar.

First, Rietveld refinement was carried out to determine whether other crystalline phases besides monoclinic BiVO 4 are present in any of the samples. However, due to the similarity of the diffractograms of tetragonal and monoclinic clinobisvanite BiVO 4 phases 20 new diffractograms were recorded with a Mo K-alpha 1 X-ray source, which has only one X-ray line compared to the two X-ray lines in a standard Cu K-alpha X-ray source. As a result, diffractograms with an overall lower peak broadening and, thus, a higher resolution could be obtained. On these newly acquired diffractograms, Rietveld refinement was performed. As is shown in Figure 4.10, the presence of tetragonal BiVO 4 for the truncated bipyramidal shaped and irregular shaped powders seems unlikely, due to the absence of any reflections characteristic for tetragonal BiVO 4

and the Rietveld refinement suggesting a presence below 5% for tetragonal BiVO 4 . In contrast, for the commercial BiVO 4 sample, reflections of tetragonal and orthorhombic bismuth vanadate phases can be easily seen in Figure 4.10. The Rietveld refinement suggested that the commercial BiVO 4 powder seemingly contains 5% of the orthorhombic bismuth vanadate phase and 15% of the tetragonal bismuth vanadate phase, however the real share of the tetragonal phase will very likely be lower due to the mismatch of the Rietveld fit with the measured diffractogram.

In Figure 4.11 the morphology of the commercial BiVO 4 powder is shown. A very high size dispersity can be noticed with particles having a size between < 1 µm and 50 µm. Additionally, the powder morphology seems to be completely irregular, as no distinct shapes can be recognized.

Diffuse reflectance spectroscopy was performed to determine the direct optical band gap of 2theta (°) Intensity (arb. units)

Truncated bipyramidal

Irregular Commercial the different powders (Figure 4.12). As can be seen in the Tauc plot [F(R)hν ] 2 vs hν , the absorption onset lies around 2.5 eV for the truncated bipyramidal and irregular shaped BiVO 4 , which agrees well to the 2.4 -2.5 eV direct band gap that others have found for undoped monoclinic BiVO 4 . 16,21,22 For the commercial BiVO 4 powder, the absorption onset lies lower, at 2.35 eV, which can be due to the presence of the Bi 2 VO 5.5 phase, as detected during the Rietveld refinement analysis. Indeed, an optical band gap of about 1.9 eV was previously determined for Bi 2 VO 5.5 . of BiVO 4 , the core level spectra of the BiVO 4 powder samples show small signs of inhomogeneous surface charging, like the small bump at a binding energy of 162 eV in the Bi4f spectrum. 24 The inhomogeneous sample charging during the XPS measurements, however, does not prevent the determination of the binding energies in this case. The Bi4f 7/2 core level binding energy is situated between 159.0 eV and 159.3 eV, which matches to Bi 3+ in BiVO 4 and the V2p 3/2 core level position between 516.7 and 517.0 eV is typical for V 5+ in BiVO 4 . 25 The 0.3 eV difference is possibly related to electronic shifts due to different doping levels. Additional oxidation states cannot be observed in the V2p 3/2 and the Bi4f 7/2 spectra, because inhomogeneous surface charging induces line broadening and extra features at higher binding energies, which interfere with potential minority signals belonging to other elemental oxidation states or different local chemical environments. In the O1s spectrum two signals can be seen at about 530 eV and at about 532 eV, which correspond to metal oxides and alcohol+carbonyl groups, respectively. 26 The hydroxyl+carbonyl signal is due to the sample being prepared under ambient conditions. The photodegradation of Rhd B through BiVO 4 is considered to be a first order chemical reaction, 27 which should follow the first order rate equation:

ln ( A t A 0 ) = kt (4.1)
with t being the degradation time and k corresponding to the first order rate constant. The first order rate constant can then be determined from the slope of a ln (

A t A 0
) vs t plot. These plots and all first order rate constants are shown in Figure 4. 16. The first order rate model seems to fit more or less, as the linear regression yield R 2 values in the range of 0.96 -0.99. Especially, the irregular shaped BiVO 4 powder seems to have a high first order rate constant. Thus, one could argue whether the truncated bipyramidal morphology is the best BiVO 4 morphology to use for oxygen evolution. However, the dye degradation efficiency does not always correspond to the water splitting performance, 2,28 most likely due to a difference in reaction mechanisms.

Anisotropic photodeposition on BiVO 4

After synthesizing a batch of 10 g of coexposed {010}/{110} truncated bipyramidal BiVO 4 powder, photodeposition experiments were carried out to test whether new BiVO 4 based anisotropic heterostructures could be synthesised. The general photodeposition procedure is disclosed in Section 3.1.4 and the experimental details for the photodeposition experiments can be found in Table 4.2.

BiVO 4 /Ag

The photoreduction of silver nitrate to metallic silver is a standard test to assess whether a certain light absorber is qualified for the photodeposition process, because the light induced silver reduction is a straightforward reaction that only involves the transfer of one electron: AgNO 3 (aq) + e - Ag(s) + NO - 3 (aq) Also, due to the relatively large size and cubic morphology, silver deposits are easy to distinguish in FESEM. 29,30 The facet-selectivity of the silver photodeposition on truncated bipyramidal BiVO 4 has been tested before and was found to result in the selective deposition of silver onto the {010} facets, 1 which are the facets on which, after light absorption, electrons reportedly accumulate. 31 Herein, the photodeposition of silver was tried out to verify whether our synthesized truncated bipyramidal BiVO 4 particles were susceptible for photodeposition and whether silver would again be selectively deposited on the {010} facets. As can be seen in the FESEM images in Figure 4.17 Some deposits that can be seen in Figure 4.18a do not show up in the silver EDS map, which is due to the shadowing of the characteristic silver X-rays by the BiVO 4 particles. no extra signals besides those belonging to Bi 3+ and V 5+ can be observed. In the Ag3d spectrum two signals can be seen with binding energy maxima situated at 368.3 eV and 374.3 eV, indicative of Ag3d 5/2 and Ag3d 3/2 , respectively. Those binding energies are characteristic for metallic silver, 26 confirming that the photodeposition of silver follows a reduction process. Please also take a look at Section 2.4 where all the main core level spectra, which can be found in this work, are described.

ANISOTROPIC PHOTODEPOSITION ON BiVO

1 μm Bi V Ag (a) (b) (c) (d) (e)

BiVO 4 /Pt

The next material which was used to coat the BiVO 4 particles through photodeposition was platinum. Overall, platinum is considered to be one of the best water reduction electrocatalysts 3 and just like silver, platinum has been found to be selectively deposited onto the {010} facets of BiVO 4 through the photodeposition method. 1 The Pt deposits were, however, considerably smaller than the Ag deposits, more in the range of 5 -10 nm, which make them slightly more complicated to spot in FESEM images. An additional complication is the use of the acidic H 2 PtCl 6 precursor for Pt photodeposition, 1,32,33 weight percentage that is aimed for, influences the pH of the reaction mixture. These pH differences can have an impact on the morphology and oxidation state of the platinum deposit, as was shown previously for the pH controlled photodeposition of Pt onto TiO 2 single crystals. 34 In our experiments, a reaction mixture pH of approximately 2.5 and 3.6 were measured for the photodeposition experiments where a 10 wt% Pt and 1 wt% Pt, respectively, were intendedly deposited.

In Figure 4.20 the effect of different pHs and Pt wt% can be seen. At a pH of 2.5 the deposited Pt particles seem to be scarce and have sizes of about 100 nm. Additionally, the BiVO 4 particles have lost part of their shape, which is most likely due to the chemical instability of BiVO 4 in an acidic environment. 35 Therefore, an additional experiment was performed at a higher pH, namely at pH 5.8 through the adjustment of the pH with NaOH. The deposit seems to be considerably smaller and more spread out over the BiVO 4 particles. Facet-selectivity could not be clearly observed, but the deposited particles on the {010} facets seem to be somewhat larger. For the 1 wt% Pt deposit, the platinum particles cannot be easily distinguished in the FESEM images due to the relatively small size, just a few nm, and the low contrast of the images. However, the Pt particles seem to have selectively accumulated onto the {010} exposed surfaces.

EDS could not be used any more to study the spatial distribution of the deposited Pt particles, because of the platinum particles being smaller than the 50 nm lateral resolution of EDS.

Instead, scanning Auger microscopy (SAM) was tried, because of the higher surface sensitivity and the slightly better lateral resolution, as explained in Section 3.4.3.1. There were, however, a few difficulties to collect elemental maps with SAM. The possibility of obtaining elemental maps with a higher lateral resolution compared to EDS is an advantage, which in consequence require more points to be scanned in order to acquire a complete elemental map of a certain area.

Combined with the long acquisition times per measurement point, the scan time to obtain a complete elemental map of a certain area with SAM is substantially longer than with EDS. Long scan times are, however, problematic for the poor electron conducting micron-sized BiVO 4 particles, because they tend to charge during the exposure to the primary electron beam. This charging phenomenon caused image drift during map acquisition, and for scan times exceeding 30 minutes, often led to the (partial) degradation of the BiVO 4 particles. Therefore, the total acquisition time had to be drastically reduced, which was done here by selecting just a few measurement zones, which were not scanned but measured as a whole. Specifically, zones across the exposed {010} and {110} facets were chosen. On these zones a SAM sputter depth profile was acquired.

However, due to the argon ions approaching from a certain direction not all regions were sputtered in the same proportion. The results of the SAM acquisition of the BiVO 4 /Pt(10 wt%) sample prepared at a pH of 5.8 can be seen in Figure 4.21. Approximately, the same Pt N 7 VV Auger signal intensity is obtained for measurement zones on both the {010} and {110} facets, meaning that the Pt photodeposition in this case was not selective to a certain crystalline facet. This observation agrees well with the FESEM images shown above. The Pt N 7 VV Auger signal intensity decreases in different proportion for different zones which confirms the orientation dependence of the sputtering process. The BiVO 4 /Pt(1 wt%) sample prepared at a pH of 3.6 was not selected for the SAM analysis, because we perceived the deposited platinum particles as too small with respect to the lateral resolution of the SAM technique. Therefore, we reasoned that a platinum elemental map would not have been able to show the exact spatial distribution of platinum. However, zone measurements would have been able to obtain information on the facet-selectivity of the Pt deposits. Unfortunately, the sample selection process was made before the problems with BiVO 4 SAM mapping were observed and a switch to zone measurements was made.

High resolution XP core level spectra for several heterostructured BiVO correspond to the Pt4f 7/2 binding energies of metallic platinum, platinum hydroxides and platinum oxides, respectively. 26 Additionally, the O1s signal at a binding energy of 531.

BiVO 4 /CoO x OH y

For {010}/{110} dual exposed BiVO 4 particles, electrons have been found to accumulate on the {010} facets, while holes tend to accumulate on the {110} facets, which has allowed the selective deposition of transition metal oxides such as MnO x and CoO x onto the {110} facets through a photooxidation process. 1,2 We again tested the facet-selectivity of the photo-oxidation of Co(NO 3 ) 2 on our truncated bipyramidal BiVO 4 particles. Indeed, the increased roughness on the {110} facets on the BiVO 4 particles, which is seen in Figure 4.23, corresponds most likely to cobalt based particles, as Li et al. found a similar morphology for heterostructured truncated bipyramidal BiVO 4 /CoO x OH y . 2 The {010} facets remain largely uncovered.

Again, SAM was tried to figure out how cobalt is exactly distributed over the BiVO 4 particles.

Luckily, the BiVO 4 /CoO x OH y sample did not degrade during the acquisition of the elemental maps. The elemental maps in Figure 4.24 confirm more clearly that cobalt is more or less restricted to the {110} facets. Another remarkable finding in these maps is that the bismuth and vanadium signal intensity vary for the different facets. The surface of the {010} facets seems to be more vanadium rich while the {110} facets seem to be more bismuth rich. Previously, anisotropic charge carrier mobilities 39 or higher upwards band bending for the {011} facets 31,40 have been cited as potential reasons for the stronger hole accumulation on these facets. However, the observation that the V/Bi surface stoichiometry is facet specific could also influence charge carrier separation. XPS was performed to determine the oxidation state of the cobalt species of the heterostructured BiVO 4 /CoO x OH y (Figure 4.25). Inhomogeneous surface charging seems to be largely suppressed, as no substantial line broadening or additional unexpected signals appear in the V2p 3/2 and Bi4f core level spectra. Thus, the Co2p 3/2 spectrum could be carefully analysed to find indications of the chemical speciation of cobalt. The Co2p 3/2 spectrum is quite complicated, consisting of an asymmetric main line and several satellites due to plasmon loss features, electron correlations and final-state processes. 41,42 Through peak fitting the exact cobalt species can be obtained. 41 However, it has been stressed that high resolution spectra should be used and that preferably the Co2p 1/2 is included in the fitting procedure. 43 Fitting of the Co2p spectrum was in our case not feasible due to the relatively high amount of noise and the possibility of inhomogeneous surface charging. However, the Co2p 3/2 spectrum in Figure 4.25 can be compared to the spectra of well prepared, well defined cobalt species. The Co2p 3/2 spectrum of our sample is similar to CoO, and Co(OH) 2 with a broad main line at a binding energy around 780 eV and a shoulder at 786 eV. 41,44 The Co2p 3/2 satellite structure of CoOOH and Co 3 O 4 is considerably different, as these species show a much smaller shoulder at 786 eV and an additional bump at a binding energy of 790 eV, 41,45 which some authors explain by the (partial) cobalt oxidation state change from Co 2+ to Co 3+ for CoOOH and Co 3 O 4 . 44,46 Thus, our sample contains seemingly more Co 2+ , indicating that cobalt is most likely only partly oxidized during the photodeposition process. Furthermore, Li et al. found a similar Co2p 3/2 spectrum for their photodeposited cobalt species. 2 Additionally, the cobalt species has been found to change during the photo-oxidation of water. 44,47

Co

BiVO 4 /NiO x OH y

So far we know, the regioselectivity of the photodeposition of nickel based oxides onto particulate truncated bipyramidal BiVO 4 and its effect on the photocatalytic efficiency has not yet been extensively studied, although nickel based compounds are considered to be performant oxidation electrocatalysts. 48,49 Nevertheless, the photodeposition of nickel based compounds has already been realised before; having been successfully deposited onto several semiconductors, such as SiO 2 , 50 CdS 51 and even on BiVO 4 thin films. 52 Moreover, Ni based compounds have proven to enhance the photochemical water oxidation efficiency of BiVO 4 thin film photoanodes. 4,52,53 We have used the experimental nickel photodeposition procedure published by Choi et al, 52 which involves Ni(NO 3 ) 2 as nickel precursor and a potassium borate (KBi) buffer solution to fix the reaction mixture pH to 9.2. However, the nickel concentration in the reaction mixture should not be too high, since nickel then reacts to form Ni(OH) 2 , which directly precipitates. For instance, a blue Ni(OH) 2 precipitate was detected when we wanted to prepare a 10 wt% nickel solution. In this case, the precipitate was discarded and the solution which then yielded less nickel was used, so that the deposit that resulted from this reaction mixture is referred to as the ± 10 wt% nickel deposit. For the 1 wt% and 0.1 wt% nickel solution no Ni(OH) 2 was perceived.

The FESEM images in Figure 4.26 seemingly show that the regioselectivity of the nickel-based deposit depends on the weight percentage of nickel we wanted to obtain. For the samples containing ± 10 wt% and 1 wt% Ni, nickel seems to be unselectively deposited over the entirety of the BiVO 4 particles, which is in contrast to the observed facet-selective photo-oxidation of cobalt.

Possibly, the pH plays a role in the regioselectivity of the nickel species, as was also previously noticed for the platinum photoreduction. So far, only the silver and cobalt photodeposition processes followed the expected facet-selectivity and these were deposited at pH 7. Elemental EDS mapping of nickel (Figure 4.26c) also shows that the surface of the BiVO 4 particles seems to be completely covered by nickel deposits for the BiVO 4 /NiO x OH y (± 10 wt%). Contrarily, for the low 0.1 wt% nickel deposit, a certain regioselectivity can be observed in the FESEM images, as the nickel deposit seems to be preferably deposited on the edge between two {110} facets.

The regioselectivity of the nickel deposit was further studied with SAM. Unfortunately, no elemental maps could be made due to gradual sample degradation under the electron beam. Hence, zone measurements, which require a much shorter acquisition time, were again performed to probe the different facets. The edges between the {110} facets could not be selectively measured due to the limited lateral resolution. For the BiVO 4 /NiO x OH y (1 wt%) sample (Figure 4.27), an intense Ni L 3 VV Auger signal was detected for every zone. Hence, the SAM zone measurements confirm the observations made earlier based off the FESEM images. Namely, that nickel seems to be omnipresent. On the other hand, for the BiVO 4 /NiO x OH y (0.1 wt%) sample (Figure 4.28), the Ni L 3 VV Auger signal is only visible when zones on the {010} surfaces were measured. Therefore, the 0.1 wt% NiO x OH y deposit appears to be limited to the {010} facets and to the edges between the {110} facets. to the Co2p core level spectrum, the Ni2p spectrum consists of many satellites which require proper peak fitting to be certain of the analysed nickel species. However, inhomogeneous surface charging and a low intensity-to-noise ratio did not allow proper peak fitting, so that the nickel species can only be roughly estimated from the Ni2p line shape. Pure NiO can be excluded, since that would give rise to a noticeable peak splitting of the main Ni2p 3/2 line at a binding energy around 856.5 eV, 57 which is not seen in our Ni2p spectra. Instead, the Ni2p line shape resembles more that of a mixed nickel oxyhydroxide. 57,58 Furthermore, in the O1s core level XP spectrum of the BiVO 4 /NiO x OH y (1 wt%) sample, a strong signal at a binding energy of 532.1 eV can be seen, which is an indication of the presence of nickel hydroxides. 26,58 The hydroxide signal is commonly present for samples prepared under ambient conditions, as these conditions promote the adsorption of carbohydrates and water onto the exposed samples. However, the signal here is considerably stronger than usual and even higher than the O 2-signal at a binding energy of 530.3 eV. Therefore, the deposited nickel species is denoted as NiO x OH y . Above, the photoreduction of Ni 2+ to metallic nickel was perceived as a possible reaction for solutions with a high enough Ni 2+ concentration. However, in the Ni2p spectrum no indication of metallic nickel, which has a characteristic Ni 2p 3/2 binding energy of 852.6 eV, 59 could be found. This could be due to the high reactivity of metallic nickel, having been found to oxidise at a relatively high rate at room temperature. In a photochemical process the deposited metallic nickel is most likely involved in a follow-up redox reaction, wherein metallic nickel reacts with hydroxides according to:

E NiOOH/Ni 2+ (RHE) = E 0 + 0.059 V 1 x log( [H + ] 3 [Ni 2+ ] 2 ) + 0.059 V x pH (4.9) E NiOOH/Ni 2+ (RHE) = E 0 -0.059 V x log([Ni 2+ ] 2 ) -0.118 V x

BiVO 4 /RuO x

RuO x is another good water oxidation electrocatalyst 48 for which, to our knowledge, the regioselectivity on truncated bipyramidal BiVO 4 has not been studied before. Photodeposition of RuO x has already been demonstrated through a photo-oxidation process, using RuCl 3 .3H 2 O as ruthenium precursor, 60 and a photoreduction process, using KRuO 4 as ruthenium precursor. 61 We tried both ruthenium precursors but their limited water solubility was a serious problem. Only ruthenium solutions for the deposition of 0.1 wt% Ru or lower could be prepared with RuCl 3 .xH 2 O and KRuO 4 . At the time, we thought that 0.1 wt% Ru would be too low to study the ruthenium regioselectivity with FESEM and SAM, so that the BiVO 4 /RuO x (0.1 wt%) samples were not further investigated. However, as is shown above for the BiVO 4 /NiO x OH y samples, which were prepared later, the regioselectivity of deposits with a coverage as low as 0.1 wt% may be studied through FESEM and SAM. Thus, in principle one could look again at the regioselectivity of the photodeposition of RuO x on truncated bipyramidal BiVO 4 with only 0.1 wt% Ru water based solutions.

Another way to deposit more ruthenium is to use another solvent system in which the solubility of ruthenium is higher. Ruthenium acetylacetonate, Ru(acac) 3 , has a high solubility in acetic acid and THF, making it possible to establish a 10 wt% Ru solution. Because sodium iodate did not dissolve in either THF and acetic acid, FeCl 3 was added instead as an electron scavenger.

Figure 4.31a shows, however, that the immersion in acetic acid had an adverse effect on the BiVO 4 morphology, as the particles appear to be more aggregated and seem to have lost their truncated bipyramidal shape in part. These adverse effects are most probably due to the instability of BiVO 4 in acidic media. For the photodeposition in THF, the BiVO 4 particles maintained their shape and a deposit can be seen, which is seemingly spread out over the entirety of the BiVO 4 particles (Figure 4.31b). However, when measuring the Ru3d core level spectrum, only a signal at a binding energy of 285 eV could be detected, which is characteristic for C1s photoelectrons of adventitious carbon. 62 Thus, RuO 2 seems not to be present, as it should have generated a signal at a binding energy of about 281 eV, the characteristic Ru3d 5/2 photoelectron binding energy for RuO 2 . 26 Instead, iron characteristic photoelectrons were detected, suggesting that the deposit seen in Figure 4.31b was the result of iron photoreduction. We were, however, not interested in the deposition of iron based particles, so that we did not further investigate the supposed photoreduction of iron.

Binding

Rhd B photodegradation

With all the heterostructures fully characterized, the Rhd B photodegradation efficiency was tested. The procedure of Rhd B photodegradation is explained in Section 3.4.5.1. In Figure 4.15 the evolution of a Rhd B photodegradation experiment is depicted by means of UV-Vis spectra, which were measured for aliquots taken from the Rhd B solution at specific points in time. The mean first order rate constants k mean and mean R 2 linear fitting parameter are denoted in the graph. For each heterostructure there was a 0.1 wt% loading of contact material.

Sacrificial water oxidation

Photochemical sacrificial water oxidation experiments were performed to assess how the performance of the newly synthesized BiVO 4 /NiO x OH y relates to previously studied heterostructures.

The sacrificial water oxidation experiments were performed within a closed reactor, coupled to an automatic burette to register the gas volume increase during the reaction. The exact experimental procedure is described in Section 3.4.5.2. In Figure 4.35 the gas evolution of all the BiVO 4 heterostructures is depicted. The gas evolution curves show that the sacrificial water oxidation with BiVO 4 follows a photocatalytic mechanism, as in the absence of light or BiVO 4 (Figure 4.36) no steady gas evolution occurs. Nevertheless, a volume increase of about 1 mL can be seen for the blank experiments without BiVO 4 (Figure 4.36). This increase is largely due to the set up searching for a new equilibrium after the light source is focused onto the reactor. In addition, a small amount of air present in the reaction medium likely contributes to the small amount of detected gas as well, since GC measurements, ran after the blank experiment, revealed an increased nitrogen signal with respect to the GC measurements ran before the blank experiment.

A first remarkable observation is that for all BiVO 4 powders, gas evolution takes place until a gas volume of approximately 12 mL (or 11 mL after subtracting the blank volume) is reached.

This volume corresponds to the concentration of sacrificial agent (1 mmol), as 11. In other words, to produce 0.5 mmol (11.2 mL) of O 2 , which consumes 4 holes (See Section 2.1.3), 1 mmol of NaIO 4 , which consumes 2 electrons, is needed. The IO - 4 /IO - 3 reaction has a standard reduction potential of 1.589 V, 55 which means that NaIO 4 is a strong oxidative agent and should be easily reduced by BiVO 4 .

Thus, the oxygen evolution stops when no IO - 4 is present any more and, as can be seen in Figure 4.35, the gas evolution then stops quite suddenly, reaching a volume of about 12 mL. This is a remarkable observation, first, because it demonstrates that the reduction of IO - 4 is not the rate-limiting step and, secondly, because it means that IO - 3 , here, does not act as a sacrificial agent, despite the reactions: having high standard reduction potentials of 1.210 V, 1.154 V and 0.269 V, respectively. 55 Due to the high standard reduction potentials, these reactions should not be thermodynamically limited for BiVO 4 . Indeed, NaIO 3 has been used sometimes in sacrificial water oxidation 1,2 or as a redox shuttle (IO - 3 /I -) in Z-scheme water splitting systems. 63,64 However, the tests we ran with NaIO 3 as sacrificial agent did not show any gas evolution.

IO - 3 + 6 H + + 5 e -
To determine the oxygen evolution rate of all heterostructures, a linear regression was performed on the linear part of the gas evolution curves, i.e. at a reaction time between 120 min and 360 min (Figure 4.35). The raw data of the linear regression can be found in Appendix A.

Seemingly, the implementation of CoO x OH y and Pt as contact materials boost the water oxidation performance of BiVO 4 by about 10 times and 5.9 times, respectively, with respect to bare BiVO 4 . In comparison, Li et al. found that bare BiVO 4 , BiVO 4 /Co 3 O 4 and BiVO 4 /Pt yielded oxygen evolution rates of 1.0, 2.5 and 16.8 µmol h -1 , respectively. 2 Thus, qualitatively, their results compare well to ours, in the sense that we both found boosted water oxidation performance for cobalt and platinum based contact materials. However, we found much higher absolute gas evolution rates for all BiVO 4 systems, while the relative boost to the water oxidation performance by Pt and Co is lower than what Li et al. found. Under the assumption that we can compare the absolute results, our bare BiVO 4 powder seems to be of a better quality than the one synthesized by Li et al, especially because the authors used a threefold higher amount of photocatalyst and sacrificial agent in their procedure. Additionally, it is remarkable that in the publication of Li et al. Pt, perceived as a reduction cocatalyst, yielded a higher boost than Co 3 O 4 , a perceived oxidation cocatalyst. This could be an indication that in their testing procedure, the reduction of the sacrificial agent was the rate limiting step, which is not a good way to assess the water oxidation performance of BiVO 4 heterostructures. The problem could be that the NaIO 3 sacrificial agent, as observed in this work, has a much lower reactivity than NaIO 4 . The use of NaIO 3 instead of NaIO 4 could also explain why Li et al. obtained lower absolute oxygen evolution rates.

Unfortunately, our newly synthesized BiVO 4 /NiO x OH y heterostructure exhibits a rather poor water oxidation performance compared to the other heterostructures, as the oxygen evolution rate is only two times higher with respect to bare BiVO 4 .

Conclusion

The {010}/{110} dual facet exposed BiVO 4 microcrystalline powder could not be straightforwardly synthesized according to any of the previously published synthesis procedures we attempted. However, by optimizing the parameters of one of the most promising synthesis procedures we could obtain BiVO 4 powder with the appropriate morphology and phase purity. The optimised parameters included a pH of 1.5, having an autoclave temperature of 200 ℃, an autoclave time of 20-24 h, and having SDS as surfactant. The synthesis could then easily be upscaled by a factor of two by multiplying the amount of all chemicals used by a factor of two, except for the amount of SDS, which had to be multiplied by three. Additionally, the reaction mixture was stirred at 300 RPM while performing the hydrothermal treatment. A remarkable finding was that syntheses performed at different pH yielded BiVO 4 powders with different morphologies, giving rise to dendritic and irregular shapes besides the desired truncated bipyramidal shape. It was shown that the irregular morphology exhibited a slightly better dye decomposition performance than the other morphologies. BiVO 4 /Pt, BiVO 4 /Ag, BiVO 4 /CoO x and the new BiVO 4 /NiO x OH y heterostructures could be successfully synthesized through a photodeposition procedure. Thorough electron microscopy analysis clearly showed that silver was selectively deposited onto the {010} facets and CoO x OH y could be selectively deposited onto the {110} facets. This facet-selectivity was also observed previously in literature. Platinum seems to be deposited selectively onto the {010} facets for weight percentages below 10%. However, this facet-selectivity is somewhat more uncertain, because the platinum particles could not be easily visualised through FESEM. The regioselectivity of NiO x OH y was remarkable, because there seems to be an increased amount on the edges of the {110} facets as well as on top of the {010} facets, as was shown by FESEM and SAM.

For dye decomposition, BiVO 4 /Ag and BiVO 4 /NiO x OH y were almost twice as efficient compared to bare BiVO 4 , whereas BiVO 4 /CoO x OH y did not improve upon bare BiVO 4 . In contrast, for sacrificial water oxidation, BiVO 4 /CoO x OH y yielded a tenfold increase in oxygen evolution rate with respect to bare BiVO 4 . This proves again that heterostructured photocatalyts which are efficient at dye degradation are not necessarily good at water oxidation and vice versa. Further, coating BiVO 4 with platinum effectively increased the water oxidation performance by a factor of 5.9. The water oxidation performance boost with cobalt and platinum contact materials was observed before by Li et al., 2 however we surprisingly obtained a seven times higher absolute oxygen evolution rate for bare BiVO 4 with three times as less BiVO 4 powder. An unfortunate finding was that our newly synthesized BiVO 4 /NiO x OH y heterostructure only boosted the oxygen evolution rate by a factor of two, thus much lower than what we obtained with CoO x OH y as contact material.

Without question, our results show that heterostructuring a photoabsorber with noble and transition metal compounds have a noticeable boost on its photocatalytic performance. But why those contact materials exactly benefit the photoabsorber is still an open question. In the next Chapter we try to solve part of this puzzling question. (10) Successful structuring strategies that have previously been identified are nanostructuring, cocatalyst deposition and creation of p-n junctions. [1][2][3][4][5] Combining these strategies with the incorporation of hexavalent ions, like Mo 6+ or W 6+ , have allowed BiVO 4 based photoanodes to achieve photocurrents which are among the highest that have been measured for photoanodes, admittedly using an external bias of 1.23 V RHE or through coupling to a photovoltaic system. 3,[6][7][8] To reach high efficiencies heterostructuring seems essential, however, the exact influence of a specific contact material on the water oxidation efficiency is not yet fully understood. Therefore, current research efforts focus on determining and quantifying the effects of heterostructuring to figure out whether surface passivation, improved reaction kinetics or enhanced charge carrier collection has the strongest influence on the efficiency of BiVO 4 .

Especially cobalt and nickel containing compounds have been found to enhance the BiVO 4 water oxidation efficiency, whereas traditionally high performing noble metal based electrocatalysts have only been sparsely reported as BiVO 4 cocatalysts. [1][2][3]6,[8][9][10][11] Diverse reasons have been stated for the good effects of the transition metal oxides ranging from improved kinetics, 9,10,12,13 suppression of surface recombination 10,14,15 and more efficient charge carrier separation. 14,16,17 However, many of those conclusions were drawn based on theoretical models or through techniques that cannot fully differentiate between surface passivation, oxidation kinetics and charge carrier recombination. Techniques which are able to elucidate specific effects of heterostructures include intensity modulated photocurrent spectroscopy (IMPS), 15 which can be used to distinguish between surface recombination and charge carrier injection, transient absorption spec-troscopy (TAS), 17 which gives information on charge carrier dynamics, surface photovoltage spectroscopy (SPS), 18 which can determine the photovoltage and detect charge carrier trapping, and dual-working-electrode (DWE) photoelectrochemistry, 19 which is useful to study junction behaviour in situ.

Although, IMPS, TAS, SPS and DWE photoelectrochemistry may help to understand heterostructures, they are still "indirect" techniques in the sense that they give no direct information on the interfacial electronic band structure and how that affects the contact properties in the photoanode. Due to the apparent lack of straightforward experimental techniques, the interfacial contact properties are often estimated from the band positions before contact. [20][21][22][23] One way of determining the band positions before contact is through (electrochemical) flat band potential measurements at the point of zero zeta potential, whereby the Mott-Schottky analysis is used most often. [24][25][26][27][28] These band positions should, however, be carefully interpreted since the Mott-Schottky approach is based on several assumptions which are not always met experimentally. 29 Another way of estimating the band positions was popularised by Butler and Ginley, who derived empiric equations that allowed to estimate the valence band edge from the electronegativity of the semiconductor, which was calculated from the electronegativity of the constituent atoms of the semiconductor. 30 Xu et al. demonstrated that for many semiconductors the band edge positions obtained through flat band potential measurements or electronegativity calculations were the same or within a few 100 meV from each other. 31 The semiconductor band positions are often referenced with respect to the hydrogen reduction potential. Now, to obtain the interfacial alignment of two materials from the individual band positions one can subtract the energy level alignment with respect to the hydrogen evolution potential. However, this is not a direct measurement but assumes that band alignments are transitive. Up to now, this hypothesis has not been demonstrated yet for electrochemical alignments and has been proven to not be valid for classical semiconductors. 32 Moreover, this kind of interfacial band alignment neglects the actual surface and interface structures and the possibility of Fermi level pinning.

In this Chapter an alternative approach, based on the combination of thin film sputtering and photoelectron spectroscopy, for studying the interfacial alignment will be described with respect to BiVO 4 based heterostructures. This technique, which was first used by Waldrop and Kraut, gives direct information on the interfacial electronic structure. 33 A more detailed description of this technique can be found in Section 3.4.4.1. These so-called interface experiments were carried out with thin film BiVO 4 substrates, which were fabricated through magnetron sputtering and pulsed laser deposition. In the first Section of this Chapter the surface properties of these thin films will be described. Then, the results of the Waldrop and Kraut method will be revealed for the BiVO 4 /ITO, BiVO 4 /RuO 2 , BiVO 4 /CoO x and BiVO 4 /NiO heterostructures. The contact materials were chosen because of two criteria. On the one hand, contact materials with different work functions were chosen, with ITO having a low work function and NiO, RuO 2 and CoO x having high work functions. [34][35][36] On the other hand, the aforementioned contact materials were considered because of their role in BiVO 4 based devices, with NiO, CoO x and RuO 2 being identified as performant oxidation electrocatalysts, 37 and ITO as a transparent conductive electrode that can serve as a back contact. 6,38 5.2 BiVO 4 thin film substrates

Gradient BiVO 4

Two different BiVO 4 thin film substrates were used, which are denoted throughout the manuscript as gradient BiVO 4 thin films and as PLD BiVO 4 thin films. The gradient BiVO 4 thin film substrates were fabricated through reactive magnetron cosputtering. The complete experimental details can be found in Section 3.2. In this Subsection the characteristics of the gradient BiVO 4 thin films will be exhaustively described.

Due to the absence of sample rotation during the BiVO 4 thin film co-sputtering process, a lateral variation in the V/Bi composition is observed. Raman spectroscopy was used to identify the phases along the Bi/V gradient. Raman spectra at specific positions along the lateral Bi/V gradient are shown in Figure 5.1. Along the entire gradient, monoclinic BiVO 4 is observed as the main phase with characteristic bands at 127, 213, 328, 370, 712, 828 cm -1 . 39 Features at 828 and 712 cm -1 can be attributed to antisymmetric and symmetric stretching vibration modes of the VO 4 tetrahedra. 40 The bending modes of the VO 4 tetrahedra are the cause of the bands at 370 and 328 cm -1 whereas those at 213 and 147 cm -1 correspond to external modes 41 and crystal lattice vibration of BiVO 4 . 40 In the V rich part of the film additional signals appear at 147, 287, 531, 997 cm -1 . These signals most likely correspond to the frequencies of the Raman active modes of orthorhombic V 2 O 5 . 39,42 Towards the middle of the film, at 10 mm distance from the Bi edge, no extra bands can be observed any more, which indicates that no other phases besides monoclinic BiVO 4 are present. The amount of noise increases while moving further to the Bi rich edge, making it difficult to distinguish any other phases besides monoclinic BiVO 4 .

Photoelectron spectroscopy was used to study the surface composition and electronic properties of the gradient BiVO 4 thin films. First, the samples were cleaned by an oxygen plasma treatment to remove the upper contamination layer containing carbonaceous species and hydroxides.

As can be seen in the XP survey spectra (Figure 5.2a) no core level emissions corresponding to carbon could be detected after the oxygen plasma treatment. Besides the removal of carbon from the surface, the oxygen plasma caused slight binding energy shifts but did not change the chemical composition of the film as can be noticed from the more detailed core level spectra of Bi4f, More detailed core level spectra were recorded for Bi4f, V2p 3/2 and O1s (Figure 5.4). The Bi4f 7/2 core level binding energy is found around 159.0 eV, which matches to Bi 3+ in BiVO 4 and the V2p 3/2 core level position around 516.7 eV is typical for V 5+ in BiVO 4 . 39 All core level binding energies shift towards higher energies with increasing V content in the film. Likely, part of the energy shift is due to a change in Fermi level position, as the Bi4f and O1s spectra both shift similarly by 0.18 eV from Bi rich edge to V rich edge. However, the V2p 3/2 spectra shift 0.30 eV up in energy, which is slightly different from the other core level shifts, and is probably due to a change in surface chemistry. The change in surface chemistry can be due to the presence of V 2 O 5 , which was detected with Raman spectroscopy. The change in surface chemistry can be noticed in the valence band spectra as well, as the shoulder at a binding energy of 3 eV, which corresponds to the hybridised Bi6s/O2p state, decreases in intensity moving from Bi rich edge to V rich edge of the gradient BiVO 4 thin film, whereas the signal at 5.5 eV, which belongs to the V3d state, is more pronounced around the V rich edge. 43 Binding energy (eV) Intensity (arb. units) The core level spectra were also used to determine the surface V/Bi ratio by integrating the areas of the Bi4f and V2p 3/2 core levels, and multiplying the integrated areas by the tabulated sensitivity factors, 7.632 and 1.275, respectively. 44 In Figure 5.5 the V/Bi ratios are shown as a function of the distance to the Bi rich edge of the film. Preferably, interface experiments are performed on stoichiometric BiVO 4 for which a V/Bi ratio of 1:1 would be expected. However, no position on the gradient BiVO 4 thin film exhibit a 1:1 V/Bi ratio: from 4 mm to 13 mm the V/Bi ratio increases almost linearly from 0.5 to 0.76, indicating a Bi rich surface, and at 15 mm from the Bi rich edge, a sudden V/Bi ratio increase to 1.4 can be noticed, indicating a V rich surface.

The higher V/Bi ratio near the V rich edge is probably due to the presence of a vanadium oxide phase, as was detected by Raman spectroscopy.

However, during Raman spectroscopy no additional vibrations besides those of BiVO 4 were detected, so that the presence of a crystalline Bi rich phase seems unlikely. Possibly, the calculated V/Bi surface ratios differ from the actual V/Bi surface ratios, which can happen when the sensitivity factors used to calculate the V/Bi ratios, do not properly correct for the difference in detection probability of Bi4f and V2p 3/2 core electrons. Tabulated instrument-specific sensitivity factors were used, but it could be that they are not applicable to BiVO 4 . As a matter of fact, whenever compounds with large cations were being studied with our surface analysis system improper surface stoichiometry was observed, as was the case for example for CdS. 45 In order to determine whether the calculated V/Bi ratios are incorrect, a monoclinic 1% Mo doped BiVO 4 single crystal with a (010) exposed crystal facet was tested as a reference in our surface analysis system. The Raman spectrum of the Mo:BiVO 4 single crystal in Figure 5.1 confirms the monoclinic BiVO 4 phase, because the observed bands (at 127, 213, 328, 370, 712, 828 cm -1 ) correspond to the Raman active modes expected for monoclinic BiVO 4 . 39 The single crystal was polished along the (010) surface to expose the internal structure, then introduced in the DAISY-MAT via a load lock and, subsequently, sputter cleaned with an oxygen plasma for 15 minutes. In and secondary electron cut-off of the single crystal are shifted by about 2.5 eV towards higher binding energies, a shift too large to be caused exclusively by impurity doping. The observed surface charging has, however, no effect on the XP core level areas, so that V/Bi surface ratios could still be reliably determined from the integrated Bi4f 7/2 and V2p 3/2 areas (inset Figure 5.5).

After polishing and O 2 plasma exposure, a V/Bi ratio of 0.79 was obtained using the standard instrument sensitivity factors. This V/Bi ratio differs largely from the expected ratio of 0.99, considering that 1% of the vanadium lattice sites are substituted by molybdenum. Thus, the V/Bi surface ratios calculated from the tabulated sensitivity factors are most likely incorrect and should be adjusted by a correction factor. A correction factor of 1.25 can be calculated by dividing the expected surface V/Bi ratio of 0.99 through the calculated ratio of 0.79. Using this correction factor the surface V/Bi ratios for the gradient BiVO 4 thin films were calculated again (Figure 5.5).

A V/Bi surface ratio of 0.95 is obtained at 13 mm from the Bi edge, which is close to stoichiometric BiVO 4 . To obtain the bulk V/Bi ratio EDS was used. At 11 mm from the Bi rich edge a bulk V/Bi ratio of 1.02 was determined, thus, very similar to the surface V/Bi ratio (Figure 5.7). Closer to the V rich edge no increase in the bulk V/Bi ratio is observed with EDS, indicating that the V 2 O 5 phase is probably limited to the BiVO 4 surface. From Raman, EDS and XPS measurements it can be concluded that at 13 mm distance from the Bi rich edge the gradient film consists of single phase monoclinic polycrystalline nearstoichiometric BiVO 4 . Therefore, at 13 mm distance from the Bi rich edge all photoelectron spectra were measured that were used to evaluate the interface properties of the interface experiments carried out with the gradient BiVO 4 thin films.

PLD BiVO 4

Part of the interface experiments were carried out using BiVO 4 thin films prepared through PLD.

PLD BiVO 4 thin films were chosen after the complicated surface stoichiometry of the gradient BiVO 4 thin films was discovered. However, some time was needed to prepare those thin films, so that a first set of interface experiments was performed with the gradient BiVO 4 thin films.

The exact synthesis conditions of the PLD BiVO 4 thin films can be found in Section 3.2. The PLD BiVO 4 thin films were prepared by our collaboration partners from the Institut de Recerca en Energia de Catalunya (IREC) who have published the structural and compositional characteristics of those thin films. 46 The reader is, thus, referred to this publication for additional details on the PLD BiVO 4 thin films.

In this Subsection, the surface analysis of the PLD BiVO 4 thin films is described as well as the differences with regards to the gradient BiVO 4 thin films. Murcia et al. characterised the PLD BiVO 4 thin films before, whereby PXRD revealed that the thin films were composed of phase pure BiVO 4 and EDS showed that the thin films had a bulk V/Bi ratio which was approximately 1:1. 46 Hence, the PLD BiVO 4 thin films seem to be more stoichiometric and homogeneous than the gradient BiVO 4 thin films. 46 To determine the surface V/Bi ratio, the BiVO 4 thin film substrates were mounted on a stainless steel substrate holder, introduced into the DAISY-MAT, cleaned with an oxygen plasma and analysed with our surface analysis set up. As can be seen in Figure 5.8a the C1s signal disappears after the oxygen plasma treatment, indicating that the oxygen plasma treatment effectively removed the upper contamination layer. Moreover, the BiVO 4 core level spectra (Figure 5.8b) show that the oxygen plasma did not induce a chemical change, only a small binding energy shift of about 0.15 eV for all core level spectra. The binding energies of 159.3

eV for Bi4f 7/2 and 517.0 eV for V2p 3/2 differ by 0.3 eV with respect to the Bi4f 7/2 and V2p 3/2 binding energies of the gradient BiVO 4 thin films, which is, however, still well within range of what can be expected for BiVO 4 . 39 A series of PLD BiVO 4 thin film samples was analysed by photoelectron spectroscopy and the V/Bi surface ratios of these samples were calculated (Table 5.1). Again, the tabulated surface sensitivity factors were not appropriate for calculating the surface V/Bi ratio.

Thus, the correction factor determined with the Mo:doped BiVO 4 single crystal (Section 5.2.1) was used to obtain more accurate surface V/Bi ratios. Some variation in the V/Bi ratio was observed with some films being apparently vanadium deficient. However, difference spectra of Bi4f and V2p 3/2 between films with different V/Bi ratios did not show any dissimilarities between the spectra. In addition, only a variation of about 0.15 eV is observed for the ∆BE V2p 3/2 ,Bi4f 7/2 binding energy difference. Therefore, the measured off-stoichiometry is not likely due to the presence of additional crystalline phases.

PES on bare BiVO 4 thin films

Figure 5.9 summarizes all photoelectron spectra which were measured on bare (no contact layers) BiVO 4 thin film substrates for which reasonable V/Bi surface ratios were determined. Reasonable V/Bi ratios here means ratios between 0.64 and 0.96 or between 0.8 and 1.2 after the 1.25 correction factor was applied. The Bi4f 7/2 , V2p 3/2 and O1s binding energies all linearly correlate with one another with slopes close to 1, as is expected for samples whose chemical structure does not change. Thus, interface experiments using BiVO 4 thin film substrates with a corrected V/Bi surface ratio between 0.8 -1.2 are all considered to be representative for BiVO 4 . The linear correlation of E F -E VBM with the core level binding energies is less clear, since R 2 is quite low for the linear fitting of the data points in the E F -E VBM vs BE Bi4f 7/2 plot. The low R 2 is partly due to the lower precision of the E F -E VBM determination, through linear extrapolation, compared to the core level binding energy determination, through peak maximum fitting. Another reason could be that there is also a small chemical shift of the binding energies, which is not unreasonable since the data were collected on samples before and after cleaning the surface. Anyhow, the overall difference in the shift between core level binding energies and E F -E VBM is smaller than the shift itself. The shift of the O1s, Bi4f 7/2 and V2p 3/2 binding energies, seen for all samples, amounts to 0.35 eV, whereas the E F -E VBM amounts to 0.29 eV. Hence, the shift in the core level binding energies is mostly due to band bending or due to a shift in the Fermi level position.

Both possibilities are reasonable considering that the different surface treatments can alter the number of surface states and the number of oxygen vacancies. Oxygen vacancies can alter the BiVO 4 Fermi level position as DFT calculations have shown that they act as shallow donors in BiVO 4 . 47,48 The small discrepancy between the E F -E VBM shift and the core level binding energy shift can be due to the aforementioned relative imprecision of the E F -E VBM determination or due to small changes in the valence band structure around the VBM. To verify whether the valence band structure stays the same, the O1s, V2p 3/2 , Bi4f 7/2 and valence band difference spectra of two samples, one with high core level binding energies and one with low core level binding energies, were determined (Figure 5.10). No significant change in the valence band, Bi4f 7/2 or V2p 3/2 XP spectra can be seen. Only in the O1s region a signal at 531.2 eV can be observed for the sample with high core level binding energies. Some authors suggest that the signal at 531.2 eV is representative for oxygen vacancies. 49,50 Another possibility is that peroxo-species, also found around 531.2 eV, 51 were formed due to the sample being treated by an oxygen plasma. For other non-cleaned samples, an O1s signal around 531 eV, is often related to the presence of hydroxyl and carbonyl groups. O1s binding energy (gradient BiVO 4 after 15 min oxygen plasma (yellow) and a sample with a low O1s binding energy (PLD BiVO 4 after 2 h of heating at 400 ℃ in oxygen atmosphere (blue). The spectra were normalized and shifted to the same binding energy.

3/2 BE V2p (eV) 3/2 BE Bi4f (eV) 7/2 BE Bi4f (eV) 7/2 E -E (eV)

Interface alignment

Interface experiments

With the surface stoichiometry of the different polycrystalline BiVO 4 thin film substrates fully characterised, interface experiments were carried out to elucidate the interface structure between the BiVO 4 thin films and a few selected contact materials. All the procedural steps of an interface experiment are described in Section 3.4.4.1. In short, in a typical interface experiment a BiVO 4 thin film substrate was first introduced into the DAISY-MAT system, after which the surface was cleaned through an oxygen plasma treatment, followed by a stepwise coverage of another thin film oxide through magnetron sputtering. Photoelectron spectra were recorded in between each deposition step. By analysing the change in the high resolution photoelectron spectra of the BiVO 4 core levels, information could be gathered about the in situ formed junction. Due to the samples being transferred in UHV between deposition and analysis no surface contamination could take place. The deposition parameters of the selected contact films can be found in Section First, the interface between BiVO 4 and RuO 2 was studied. In Figure 5.11 the X-ray photoelectron Bi4f, V2p 3/2 , O1s, Ru3d core level and valence band spectra after each RuO 2 deposition step are shown. The growing RuO 2 film thickness is noticeable from the increasing Ru3d core level line intensity, the decreasing intensity of the Bi4f and V2p 3/2 core level lines and the O1s intensity not changing much with each RuO 2 deposition step. The spectral features as measured for the growing RuO 2 films are in good agreement to spectra measured before for RuO 2 samples. 35,52 Normally, to investigate how the energy bands between semiconducting BiVO 4 and metallic RuO 2 align, the change in valence band maximum (VBM) of BiVO 4 should be followed, which is, however, not feasible here, due to the superposition of the BiVO 4 and RuO 2 valence band spectra.

Therefore, the changing VBM has to be evaluated from the binding energy shift of the BiVO 4 core levels, a method commonly referred to as the Kraut method. 33 Both Bi4f and V2p 3/2 core levels exhibit a similar downwards binding energy shift of 0.71 eV, which was confidently attributed to a BiVO 4 Fermi level change, caused by RuO 2 induced band bending.

Binding energy (eV) Intensity (arb. units) Next, the interface between ITO and BiVO 4 was investigated with ITO deposited at room temperature. The XP Bi4f, V2p 3/2 , O1s, In3d 3/2 core level and valence band spectra for the stepwise sputter deposition of ITO on BiVO 4 are displayed in Figure 5.12. As the In3d core level intensity increases, the Bi4f, V2p 3/2 core level intensity drops, while the O1s core level intensity remains approximately constant. In contrast to the RuO 2 deposition, the Bi4f 7/2 and V2p 3/2 core levels shift to higher binding energies after ITO deposition, which is expected regarding the substantially lower work function of ITO compared to that of RuO 2 . After the first deposition step the V2p 3/2 and Bi4f 7/2 spectra both shift by approximately 0.16 eV. Thereafter, both spectra stay around the same value with increasing film thickness. Due to both V2p 3/2 and Bi4f spectra shifting in the same way, this shift is likely to be caused by downward band bending. Also, in the V2p 3/2 spectrum a signal at a lower binding energy seemingly appears after the first ITO deposition step, which may indicate a reduction of V 5+ to V 4+ . Further in this Chapter a closer look is taken at this feature. The final 452.3 eV In3d 3/2 binding energy is typical for ITO, 53 as well as the valence band spectral features observed here for the growing ITO film. 35,54 Binding energy (eV) Intensity (arb. units)

E VBM Reduction ?

Figure 5.12: O1s, V2p 3/2 , In3d 3/2 , Bi4f core level and valence band XP spectra for BiVO 4 /ITO(RT) interface. ITO deposition times are denoted in the O1s spectra.

Since the ITO deposited at room temperature was not yet degenerately doped, a second BiVO 4 /ITO interface experiment was performed, but this time the BiVO 4 substrate was heated to 400 ℃ before each ITO deposition step, while maintaining a pure argon atmosphere during heating and cooling down. In this way, the amount of oxygen vacancies in ITO and, thus, the Fermi level position could be increased. Indeed, for ITO now the E F -E VBM amounts to a value of 2.8 eV (see Figure 5.13), indicating that the Fermi level coincides with the CBM, considering a 2.8 eV band gap. 55 This ITO interface experiment again demonstrates that the Bi4f 7/2 and V2p 3/2 spectra shift towards higher binding energies. More precisely, a shift of 0.30 eV can be observed after the first deposition step after which the Bi4f and V2p 3/2 core level spectra do not shift any more with further ITO deposition steps. As was observed for the interface experiment with ITO deposited at room temperature, an additional signal seems to appear at a lower binding energy in the V2p 3/2 spectrum after the first ITO deposition step. Moreover, in the Bi4f spectrum two additional signals appear as well, whereby the signal at 157.1 eV is characteristic for metallic bismuth. 44 Further in this Chapter the apparent reduction of the bismuth vanadate substrate due to the contact with ITO will be discussed. photoelectrons, the contact layer thickness can be calculated i and, indeed, a relatively thick 1.3 nm NiO film seems to be deposited after the first 5 s deposition step (Figure 5.15). After the second deposition step, the Bi4f and V2p 3/2 core level spectra shift an additional 0.1 eV. Afterwards, no more shifts are observed for further deposition steps. After a total deposition time of 8 min a i Due to the Bi4f photoelectrons having a higher kinetic energy than the V2p 3/2 photoelectrons, the inelastic mean free path of the Bi4f photoelectrons is larger, so that the attenuation of Bi4f photoelectrons is lower. The thickness of a thin film deposited on top of a BiVO 4 substrate can then be evaluated, under the assumption that the surface V/Bi ratio does not change during film growth, according to following formula: t = ln(

I V * I Bi,0 I Bi * I V,0 ) * λ V * λ Bi λ V -λ Bi
with t: film thickness; I V : integrated V2p 3/2 core level intensity at film thickness t; I Bi : integrated Bi4f core level intensity at film thickness t; I V,0 : integrated V2p 3/2 core level intensity at zero film thickness; I Bi,0 : integrated Bi4f core level intensity at zero film thickness; λ V : electron effective attenuation length of V2p 3/2 photoelectrons; λ Bi : electron effective attenuation length of Bi4f photoelectrons. The electron effective attenuation lengths were calculated using the "NIST Electron Effective-Attenuation-Length Database", based on the density, number of valence electrons and band gap of the deposited film. rather thick NiO layer has been deposited. As a consequence no Bi4f and V2p 3/2 photoelectrons could be detected any more. The Ni2p spectrum measured after the last deposition step shows a main line and satellite structure typical for NiO. 56 Therefore, the nickel based contact layer can be assumed to be NiO. In the O1s spectrum after the last deposition step a main emission line at 529.5 eV and a smaller one at 531.5 eV can be observed. Dupin In the final interface experiment, a CoO x contact layer was stepwise sputtered on top of BiVO 4 (Figure 5.16). The Bi4f 7/2 and V2p 3/2 binding energies, 158.7 eV and 516.4 eV, respectively, are 0.3 eV lower compared to the ones measured for the other BiVO 4 substrates. These slightly lower binding energies are most likely due to a shift of the Fermi level, as the BiVO 4 valence band maximum (E F -E VBM ) of this substrate is situated at an 0.3 eV lower position as well, at 1.5 eV.

Similar to the BiVO 4 interfaces with nickel oxide and ruthenium oxide, the Bi4f and V2p 3/2 spectra shift towards lower binding energies as a reaction to the contact with the cobalt oxide contact layer. The binding energy shift is less drastic compared to the BiVO 4 /NiO experiment due to CoO x being deposited at a lower rate, which is also noticeable from the Bi4f 7/2 and V2p 3/2 line intensities not dropping strongly after the first deposition step. The Bi4f 7/2 and V2p 3/2 binding energies both shift in a similar fashion by 0.45 eV, seen over all deposition steps. After 1 h of sputtering a sufficiently thick CoO x layer was obtained, so that the Bi4f 7/2 and V2p 3/2 core level photoelectrons could not be detected any more. Co 2+ and Co 3+ can be distinguished from each other in the satellite structure of the Co2p spectrum with a Co 2+ feature being apparent at a binding energy of 786 eV and Co 3+ at 789 eV. 57 Here, both features are clearly visible, which indicates that the deposited cobalt oxide is a mixed oxide. Also, the valence band spectrum resembles the valence band spectrum of a mixed cobalt oxide. 58 Therefore, the deposited cobalt oxide film is denoted as CoO x . Off course, Co 3 O 4 also yields both Co 2+ and Co 3+ . However, due to the CoO x thin film deposition being performed at room temperature, the thin film was assumed to be amorphous. Thus, it cannot be associated to Co 3 O 4 , which is a crystalline structure with a particular stoichiometry. Additionally, ultraviolet photoelectron spectra were recorded for all BiVO 4 thin films and contact materials (Figure 5.17). The secondary electron cut-offs in the UP spectra were used to determine the work functions. The work function of the BiVO 4 substrates, used for the RuO 2 , ITO and NiO interface experiments, can be found at 5.6 -5.9 eV, whereas a 6.6 eV work function is determined for the substrate used in the BiVO 4 /CoO x interface experiment. This higher work function finds again its origin in the lower Fermi level position of this substrate, which is here also noticeable in the valence band region of the BiVO 4 UP spectrum. The work function of RuO 2 was found to be 6.5 eV, similar to earlier published values. 59,60 Work functions of 5.4 eV and 5.7 eV were determined for NiO and CoO x , respectively. These work functions agree with earlier determined values within our group. 61,62 However, Greiner et al. measured work functions between 6 eV and 7 eV for NiO and Co 3 O 4 after an extensive photoelectron spectroscopic analysis of the first row transition metal oxides. 34 Prior, Greiner et al. found that NiO thin films prepared through in situ oxidation yielded work functions up to 6.7 eV, which, however, dropped below 6.0 eV over time in vacuum, which the authors explained by the adsorption of residual gases in the vacuum. 63 A similar effect could take place on sputtered CoO x thin films. Additionally, as was explained above, an extra line besides the main O 2-line was observed in the O1s spectrum of the CoO x and NiO thin films, which could be due to O -species in the subsurface. 51 Reasonably, these species could create a positive dipole at the surface and, thus, be the cause of a lower than expected work function. Since the NiO VBM could not be unambiguously determined from the XP valence band spectrum, the UP spectrum was used instead. A NiO VBM of 0.4 eV was obtained, confirming the p-type character of the NiO film. A 4.4 eV and 4.5 eV work function were found for ITO(RT) and ITO(400 ℃), respectively, which should actually be higher, considering the relatively low Fermi level positions of 2.5 eV and 2.8 eV and the correlation between Fermi level position and work function suggested by Klein et al. 64 Possibly, this dissimilarity in measured and expected work function indicates a surface dipole at the ITO surface.

Band diagrams

The Bi4f 7/2 and V2p 3/2 core level binding energy shifts for all interface experiments can be found in Figure 5.15 and are represented more visually appealing in Figure 5.18. The Bi4f 7/2 and V2p 3/2 binding energies for the BiVO 4 /NiO, BiVO 4 /RuO 2 , BiVO 4 /ITO(RT) and BiVO 4 /ITO(400 ℃) interface experiment shift in a similar way, which indicates an electronic shift, i.e. Fermi level position shift, rather than a chemical shift. The shift is less coherent for the BiVO 4 /CoO x interface experiment. First of all, the initial binding energy positions differ from the other substrates which, as explained above, is most likely due to a different doping level in this BiVO 4 substrate. Additionally, the Bi4f 7/2 binding energy shift is 40 meV lower than the V2p 3/2 shift, so that chemical changes during the CoO x deposition are not unlikely. Since CoO x was reactively sputtered in an oxygen atmosphere part of the cobalt ions reaching the BiVO 4 could extract oxygen atoms from the surface instead from the sputtering atmosphere. Nevertheless, the 40 meV difference is relatively small compared to the overall 0.45 eV shift, meaning that the observed binding energy shift originates mainly from an electronic shift. Also remarkable is that the binding energy shift for all contact materials, besides RuO 2 , is quite abrupt, indicating that a very thin contact layer (see Figure 5.15 for thickness estimations) already fixes the electronic properties of the junction.

The most important energetic values of the interface experiments are summarised in Table 5.2 and based on this accumulated photoelectron spectroscopy data complete energy band diagrams can be constructed for all the studied interfaces. To construct these band diagrams, the band gaps 2.4 eV for BiVO 4 , 2.8 eV for ITO and 3.6 eV for NiO were taken from literature. 55,[START_REF] Tokunaga | Selective preparation of monoclinic and tetragonal BiVO 4 with scheelite structure and their photocatalytic properties[END_REF][START_REF] Irwin | p-Type semiconducting nickel oxide as an efficiency-enhancing anode interfacial layer in polymer bulk-heterojunction solar cells[END_REF] The band gap of CoO x is somewhat chosen arbitrarily since the exact stoichiometry is not known and because widely varying band gaps have been published for mixed cobalt oxides, for instance the optical band gap of Co 3 O 4 has been found to range between 0.8 -2.5 eV. 58,62,[START_REF] Barreca | Composition and Microstructure of Cobalt Oxide Thin Films Obtained from a Novel Cobalt(II) Precursor by Chemical Vapor Deposition[END_REF][START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF] The energy band diagrams displayed in Figure 5. V2p binding energy (eV) Binding energies (BE) of the Bi4f 7/2 and V2p 3/2 core levels at 0 s coverage and for the final spectrum that could still be observed, valence band maxima (VBM) and work functions (WF) for the interface experiments. CL: contact layer. All values are denoted in eV. *Determined from UPS measurements.

3/2 / C o O x B i V O 4 / N i O B i V O 4 / R u O 2 B i V O 4 / I T O R T B i V O 4 / I T O R T B i V O 4
BiVO 4 /ITO interfaces might be caused, as explained above, by a surface dipole at the ITO surface.

It is also known, however, that semiconducting oxides often exhibit a huge dipole potential step upon forming a Schottky barrier with RuO 2 . [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF] These Schottky barrier heights would imply a work function of RuO 2 of about 5.6 eV; if this work function would be considered instead of the measured 6.5 eV work function, a similar dipole potential step would be present at both the BiVO 4 /RuO 2 and BiVO 4 /ITO interfaces. The dipole potential step is possibly partly due to the oxygen plasma treatment of the BiVO 4 surface, as the oxygen plasma may create an O 2 -surface layer, which increases the surface dipole. [START_REF] Lee | Mechanism for the increase of indium-tin-oxide work function by O 2 inductively coupled plasma treatment[END_REF]71 During the interface experiments this surface dipole would then disappear due to the formation of the oxide contact layer which would incorporate the excess oxygen in its lattice. Additionally, the experimentally determined work function of BiVO 4 is higher than expected since the BiVO 4 electron affinity obtained in this work is situated between 5.3 -5.9 eV, which is in contrast to an electron affinity of 4.5 -4.8 eV commonly published in literature. 3,42,43 As explained above the experimental NiO and CoO The second signal in the Bi4f 7/2 spectra has a binding energy of around 157.0 eV, which corresponds to the binding energy of metallic bismuth. 44 The exact binding energy of the second signal in V2p 3/2 is difficult to obtain since V 5+ can be reduced to V 4+ or V 3+ which have similar binding energies. The reduction of V 5+ and Bi 3+ could be due to Fermi level pinning, whereby electrons are not situated in the conduction band states, but form small polaron states with bismuth and vanadium atoms as the centres of these polaron states. The creation of a small polaron state on a vanadium lattice site after the injection of an electron into the BiVO 4 lattice has been theoretically predicted through DFT calculations with the polaron state having V 3dz 2 character and shifted from the conduction band edge into the band gap. 72,[START_REF] Liechtenstein | Density-functional theory and strong interactions: Orbital ordering in Mott-Hubbard insulators[END_REF] Furthermore, the localized polaron state has been experimentally verified for molybdenum and tungsten doped BiVO 4 single crystals through resonant XPS. 74 Indeed, upon closer inspection of the valence band spectrum after 4 s of ITO deposition (Figure 5.21) an additional signal appears in the band gap about 0.9 eV from the CBM, the position where the localized polaron state due to V 4+ would be expected. [72][START_REF] Liechtenstein | Density-functional theory and strong interactions: Orbital ordering in Mott-Hubbard insulators[END_REF][74] With further ITO deposition the mixing of the ITO and BiVO 4 valence band states does not allow further observation of the supposed polaron state.

For NiO, RuO NiO and CoO x no significant change in the Bi4f 7/2 core level spectrum can be observed (Figure

5.22).

There is a small bump appearing at a higher binding energy around 160.5 -161.0 eV. This could be due to the oxidation of Bi 3+ to Bi 5+ , however since this level only shows up in a low intensity we could not reliably deduce its existence. In addition, the distinction between Bi 3+ and Bi 5+ in Bi4f spectra can be vague, as was observed for mixed bismuth oxides such as BaBiO 3 . 75-77 Fermi level position can be changed relative to the band positions. The BiVO 4 band positions are here given with respect to the SHE scale based on previously performed Mott-Schottky measurements and work function measurements, which were carried out to determine the energetic distance of the BiVO 4 band positions with respect to the H + /H 2 reduction level. 20,43,78 However, to my knowledge none of the reliable Mott-Schottky measurements were performed at the pH ZPC .

Therefore, the presence of a Helmholtz layer cannot be excluded, so that the band positions at the absolute vacuum scale are actually unknown. Please take a look at Figure 2.18 for the relation between SHE, RHE and vacuum level. Therefore, in the discussion below the relative energetic positions are relevant, but the absolute positions should be taken with a grain of salt. The SHE and not the RHE scale is taken here, because the equivalence of the SHE scale with the absolute vacuum scale (E Vac -E SHE = 4.44 eV) allows to position the polaron levels as well as the standard reduction potentials in a straightforward manner.

Based on the interface experiments, the range over which the Fermi level position can be shifted seems to be limited to 1.1 eV with the lower limit at -1.5 eV vs SHE and the upper limit at -0.4 eV vs SHE. Thus, defects intrinsic to BiVO 4 do not prevent water oxidation since the H 2 O/O 2 reduction potential is situated well within the range over which the Fermi level can be changed.

Nevertheless, in many occasions Fermi level pinning has been identified as a main reason for poor BiVO 4 photoanode efficiency. 15,[79][80][81] However, as has been recently pointed out, various synthesis methods could introduce different native defects, which may be removed by post synthetic treatments. 81 Thus, by removing the native carbon and hydroxide layer, as was done in For the core level spectra the background was subtracted using a Tougaard function for the Bi4f spectra and a Shirley function for the V2p 3/2 spectra. Then, the core level spectra were normalized and shifted along the x-axis to align the line maxima. The contact layer deposition times are denoted in the spectra. Then, the core level spectra were normalized and shifted along the x-axis to align the line maxima. The contact layer deposition times are denoted in the spectra.

this work through an oxygen plasma treatment, the BiVO 4 surface states, which previously have been identified as recombination centres and the cause of Fermi level pinning, may be removed.

Nevertheless, because of the differing work functions of CoO x , NiO and RuO 2 , the BiVO 4

Fermi level limit at -1.7 eV vs SHE may indicate intrinsic defect compensation, involving gap state (polaron) formation, whereby such states could still act as charge carrier recombination centres. The observed lower Fermi level limit at -1.7 eV vs SHE differs, however, strongly from the recently theoretically calculated BiVO 4 hole polaron level, situated at -2.5 eV vs SHE. [START_REF] Liechtenstein | Density-functional theory and strong interactions: Orbital ordering in Mott-Hubbard insulators[END_REF] Alternatively, the hole defect compensation level can be estimated from the standard reduction potentials involving bismuth and vanadium. As presented in Figure 5.23, the Bi 5+ /Bi 3+ , redox couple at -1.59 eV vs SHE, 82 is close to the observed lower Fermi level limit. Hence, defect compensation by oxidation of Bi 3+ to Bi 5+ is still a possibility. Similarly, the Fermi level limit at -0.6

eV vs SHE can be due to the intrinsic defect compensation involving the reduction of Bi 3+ and/or V 5+ . Indeed, indications of Bi 3+ and V 5+ reduction were observed in the XP core level spectra of the BiVO 4 /ITO interface experiment (Figure 5.21). The reduction of Bi 3+ could be linked to the standard reduction potential of Bi 3+ /Bi(s), at -0.31 eV vs SHE. Pinpointing the standard redox levels involving V 5+ is more complicated since the V 5+ /V 4+ and V 4+ /V 3+ standard redox potentials do not formally exist, possibly due to V 5+ and V 4+ being unstable in aqueous electrolytes, wherein these ions form (hydr)oxide based compounds. Nevertheless, the V 5+ /V 4+ reduction can be represented by the V 2 O 5 /VO 2+ standard redox potential at -1.0 eV vs SHE. The Bi 3+ /Bi(s) and V 2 O 5 /VO 2+ standard reduction potentials differ, however, from the Fermi level limit at -0.6 eV vs SHE, which may be due to the solid state reduction potentials being different from the standard reduction potentials measured in an electrolyte, because of the different coordination of metal ions in an electrolyte compared to a crystal lattice. Additionally, the calculated V 5+ /V 4+ electron polaron level position 72,[START_REF] Liechtenstein | Density-functional theory and strong interactions: Orbital ordering in Mott-Hubbard insulators[END_REF] is about 0.3 eV lower than the observed upper Fermi level limit. Further reduction of V 4+ to V 3+ may be possible, but the reduction of V 3+ to V 2+ is unlikely due to its unsuitable standard reduction potential at 0.26 eV vs SHE. 20,43,78 Right: a series of standard reduction potentials related to bismuth and vanadium species are shown on the standard hydrogen electrode scale. 82,83

Solid state band alignment

Mostly, band alignment between two solid state compounds is not evaluated through interface experiments, but simply by observing how the bands of the individual compounds align before contact. [84][85][86] The band positions are in this approach most often experimentally derived through flat-band measurements. Indeed, capacitance-derived flat-band potential measurements from Mott-Schottky plots give, next to information on the kinetics and energetics of the semiconductor-liquid junctions, information on the location of the band positions with respect to the reference electrode. Such measurements were often performed in the early days of semiconductor electrochemistry (1970s -1980s), mostly on defined surfaces of single crystals. These were used to align the commonly studied semiconductor electrodes relative to the H + /H 2 redox potential. [26][27][28]87 By assuming transitivity of band alignment this provides a comparison of band edge energies of the different photoelectrodes. The energy band alignment of different oxides inferred from such studies is shown in Figure 5.24. From the electrochemical alignment of the energy bands, which positions the valence and conduction band energies relative to the H + /H 2 redox potential and which is established in literature, 26,28,88 the relative alignment between two semiconducting oxides can be directly extracted by assuming transitivity. In principle, the same can be achieved using the Fermi level position at the interfaces of the oxides with the same metal, also assuming transitivity. 89 The interface experiments with RuO 2 and ITO described in this Chapter can, in principle, both be used to compare the band edge energies. 90,91 Transitivity of band alignment is, however, not necessarily valid when Fermi level pinning is involved in interface formation, as this affects the measured Fermi level positions. [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF][92][93][94] In the extreme case of pinning, known as Bardeen limit, the Fermi level at an interface becomes independent on metal work function. In such a case, the Fermi level at the interfaces with RuO 2 and ITO would be identical. This is definitely not the case, as evidenced by Figure 5.24. Additionally, the band edge positions or dipole layer at the interface may vary due to adsorbates as well as ionic surface terminations. Such offsets may be expected, especially for semiconductor/electrolyte interfaces, but they have hardly been studied yet in a systematic manner. 32 For classical semiconductors, Fermi level pinning might be induced by virtual gap states 89,95 or by deposition-induced defects. 96 The latter are almost always observed during deposition of elementary metals on clean oxide surfaces, [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF][START_REF] Chen | Reduction-induced Fermi level pinning at the interfaces between Pb(Zr,Ti)O 3 and Pt, Cu and Ag metal electrodes[END_REF][START_REF] Körber | Barrier heights at the SnO 2 /Pt interface: In situ photoemission and electrical properties[END_REF] but can be avoided by contact formation with conducting metal oxides, depending on the applied deposition parameters. [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF] The Fermi levels observed at contacts of BiVO 4 (this work) and other oxides with RuO 2 (high work function)

F E ITO F E RuO E (eV) vs F E RuO 1 2 Φ B,RuO 2 ΔΦ B Φ B,RuO 2 (a) (b) 1.1 1.0 ΔΦ B -Bi 2 O 3 β -Bi 2 O 3 β 2
and ITO (low work function) are included in Figure 5.24. In the present case, the variation of Fermi level position at the interfaces of BiVO 4 with RuO 2 and ITO amounts to about 1.0 eV. This variation is higher than observed for rutile and anatase TiO 2 , 52 lower than the variation observed for SrTiO 3 [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF][START_REF] Giesecke | Modification of the Schottky barrier height at the RuO 2 cathode during resistance degradation of Fe-doped SrTiO 3[END_REF]100 and about the same as the variation determined for β-Bi 2 O 3 . 35 In the case of (Ba,Sr)TiO 3 and Pb(Zr,Ti)O 3 , the alignment determined using transitivity with ITO or RuO 2 and the alignment extracted from direct interface formation all agree within 0.2 eV, giving substantial credit to this approach. 90 However, there is substantial deviation between the electrochemical alignment and those obtained by aligning the Fermi energies at interfaces with RuO 2 and ITO for the materials included in Figure 5.24. This difference is valid even if the different magnitude of Fermi level splitting for the different oxides is taken into account. An instructive example is given by anatase and rutile TiO 2 , where it was shown recently using complementary approaches that the valence band of rutile is about 0.7 eV higher than that of anatase, which substantially differs from the alignment obtained using electrochemical studies. 52,101,102 The valence band maximum of BiVO 4 and β-Bi 2 O 3 is about 0.4 eV above that of SrTiO 3 according to the electrochemical alignment from flatband potential measurements. 78,103,104 Using the RuO 2 and ITO interfaces, the valence band of BiVO 4 would be 0.75/1.25 eV higher than that of SrTiO 3 . The latter alignment would correspond well with the alignment of SrTiO 3 with other perovskites having Bi 3+ or Pb 2+ as A-site cations. 90,105,106 The higher valence band maximum of the latter perovskites is explained by the contribution of the occupied 6s orbitals of Bi or Pb. 107 As higher valence band maxima are quite generally observed for such compounds, it seems reasonable to expect that the valence band maximum of BiVO 4 is also about 1 eV higher than that of SrTiO 3 , which would better agree with the alignment obtained from transitivity of RuO 2 and ITO interfaces.

The deviation of the energy band alignment obtained using transitivity from electrochemical or solid-state interfaces may be explained either by the presence of Fermi level pinning or by a fundamental difference between solid/electrolyte and solid/solid interfaces. For solid/solid interfaces, the presence or the degree of Fermi level pinning can be obtained by comparing the variation of the Fermi energy using contact materials with different work functions. Such experiments should, however, take into account that Fermi level pinning might also be caused by lattice defects, which are frequently introduced on oxide surfaces by metal deposition. [START_REF] Klein | Interface Properties of Dielectric Oxides[END_REF]108 However, even in the case of unreactive interfaces, such as those between two oxides used in the present study, bulk or surface defects may also result in a substantial variation of band alignment. 94,[109][110][111] As a consequence, one may expect a substantial dependence of the band alignment on the preparation conditions. For instance, different concentrations of oxygen vacancies in ZnO and metallic precipitates in Cu 2 O result in a variation of the band alignment by more than 1 eV at the interface between these materials. 94 In contrast, the effect of defects on the band alignment at oxide/liquid interfaces has been less extensively studied. In addition, specific surface interaction mechanisms must be considered for solid/electrolyte interfaces, which will be different in solid/solid interfaces. First of all, the substrate surface may be terminated in a different way due to specific adsorption of H + , OH -, or molecular H 2 O species from the solution phase.

Specific adsorption of supporting electrolyte species and solution contaminants may also play a role. The adsorption of ionic species results in the formation of a charged inner Helmholtz layer, which does not exist in similar form on solid surfaces. Additionally, an electrochemical double layer, the outer Helmholtz layer is formed due to electrostatic attraction of charged species or dipoles in the electrolyte to the (charged) solid surface. Due to thermal diffusion and dipolar interactions this layer can extend up to several hundreds of nm into the electrolyte solution (Gouy-Chapman or Stern layer). Finally, the different procedures in contact formation may lead to additional differing defect formation or passivation reactions at the interface. For solid/solid interfaces, defect passivation or surface reactions typically take place during contact formation, after which the contact properties show little or no further changes. In contrast, the properties of the solid/electrolyte interface can be easily and often reversibly modified afterwards by e.g. applying a potential, changing the pH of the solution, or by illumination in case of a semiconducting solid.

BiVO 4 /H 2 O

In addition to the solid state BiVO 4 interfaces studied in the previous Section through interface experiments, we studied the interaction of BiVO 4 with water through water exposure and water desorption experiments. In a water exposure experiment, a freshly calcined PLD BiVO 4 thin film was exposed to several bursts of water in a chamber operated under high vacuum, in the absence of oxygen and light (see Section 3.4.4.2 for more experimental details). In a water desorption experiment a freshly calcined PLD BiVO 4 thin film substrate was first exposed to defined amounts of water at 77 K and high vacuum, after which the water was desorbed by allowing the sample to heat up again to room temperature (see Section 3.4.4.3 for more experimental details). Photoelectron spectroscopy was used to study the electronic and chemical changes after the water exposure/desorption.

Before the water exposure experiment was carried out, the PLD BiVO 4 thin film was calcined for 2 h at 400 ℃ in a pure oxygen atmosphere to remove the upper surface contamination layer.

As can be seen in the survey spectra in Figure 5.25 all the carbonaceous surface species were evidently removed after the calcination, because no C1s photoelectrons could be detected any more. On this seemingly clean surface, the water exposure experiment was carried out. In the survey spectrum after the water exposure no additional photoelectrons besides those expected for BiVO 4 were detected. In Figure 5.26 the effect of water exposure on the Bi4f, V2p eV, respectively. These high core level binding energies are probably a consequence of a high Fermi level position, as an E F -E VBM of 2.0 eV could be determined from the valence band spectrum of the thermally cleaned BiVO 4 sample. Hence, the thermal cleaning step, used before the water exposure, probably maintains a higher amount of oxygen vacancies than a typical oxygen plasma treatment, which was used as cleaning method before the interface experiments.

Moreover, additional features in the Bi4f 7/2 , V2p 3/2 and O1s spectra can be seen after the exposure of the BiVO 4 substrate to water. The features at lower binding energies in the Bi4f and V2p 3/2 spectra indicate the reduction of Bi 3+ and V 5+ , respectively. 44 The second feature in the V2p 3/2 spectrum after water exposure has also been seen before by Favaro et al. during near ambient pressure XPS measurements, which the authors also related to the reduction of V 5+ to V 4+ . 112 The feature at 157.4 eV in the Bi4f 7/2 spectrum most likely corresponds to metallic bismuth. 113 The feature at higher binding energy in the O1s spectrum is most likely due to the formation of hydroxides. 44,114 The BiVO 4 reduction could be partially reversed by exposing the sample to an oxygen atmosphere at room temperature, as the additional features in the O1s, V2p 3/2 and Bi4f seemingly decrease after the oxygen exposure. Furthermore, an oxygen plasma The intensities of the core level spectra in Figure 5.26, determined through integration of the core level signals after background subtraction, were used to determine surface intensity ratios, which are denoted in Table 5.3. Surface treatments may change these intensity ratios through changes of the surface chemistry or through the formation/removal of an overlayer. The change Also, the extra features that appear in the Bi4f, O1s and V2p 3/2 spectra after water exposure indicate the reduction of the BiVO 4 surface. Peak fitting was performed to better assess the relative intensity of the different components in the O1s and V2p 3/2 spectra peak fitting. The fitting parameters are denoted in Table 3.2 and the results of the fitting can be seen in Figure 5.27.

The second component in the V2p 3/2 core level spectrum after water exposure is most likely due to the presence of V 4+ , although the binding energy difference only amounts to 0.76 eV, where for vanadium oxides normally a binding energy difference of 1.2 eV is measured between V 5+ and V 4+ . 115 In the fitting, the strong reduction of the vanadium sample after water exposure is again apparent as the I V 5+ /I V 4+ drops to 0.39. Simultaneously, a second feature appears in the O1s spectrum at a binding energy difference of 1.05 eV to the metal oxide O 2-species, which is close to the 1.2 eV binding energy difference normally noticed between metal oxide and metal hydroxide species. 44,114 Thus, also the fitted spectra of the water exposure experiment hint at a chemical reaction that leads to the formation of hydroxides and the reduction of V 5+ and V 4+ .

The valence band difference spectra in Figure 5.28 seemingly confirm the conclusions drawn from the core level binding energy shifts and the appearance of the additional features in the O1s, V2p 3/2 and Bi4f spectra. The features denoted in the valence band spectra correspond to that of pristine BiVO 4 . In the difference trace a decrease in intensity after the water exposure at a binding energy of about 2.7 eV, 3.7 eV and 5.5 eV can be seen, which correspond to unhybridized O2p π mixed with Bi6s states, unhybridized O2p π states and hybridized Osp 2 /V3d states. 43 The decrease in intensity of these three features is most likely correlated to the perceived decrease in the amount of surface oxygen, as determined from the reduced I O1s /I V2p 3/2 surface intensity ratio.

The increase in signal intensity in the binding energy region from 8 -10 eV could possibly be due to the formation of hydroxides at the surface. 116 The appearance of the feature at a binding energy of 1.5 eV is most likely due to the filling of the V3d states and Bi6p states, as those feature have been linked before to V 4+ for tungsten and molybdenum doped BiVO 4 samples 74 and to metallic bismuth, 117 respectively. Also in the valence band region of the UP spectra (Figure 5.29) the effects of the water exposure on the BiVO 4 sample can be seen. Namely, the distinct signals at 3.6 eV, 5.4 eV and 7.2 eV characteristic to BiVO 4 cannot be seen any more after exposing BiVO 4 to water, but are restored after the oxygen plasma treatment. The SEC also moves to a higher binding energy after the water exposure, meaning that the work function is reduced.

The results of the water exposure experiment were cross-checked through the similar water desorption experiment, where a water layer was gradually deposited on top of a liquid nitrogen cooled BiVO 4 thin film substrate and, subsequently, desorbed by allowing the sample to heat up again to room temperature. First of all, the PLD BiVO 4 thin film was calcined under the same conditions as the thin film substrates used for the water exposure experiment. However, this surface cleaning step did not fully remove the carbon contamination layer here. As can be seen in Binding energy (eV)

V 4+ V 5+ V 5+ O 2- OH O 2-

O1s

V2p 3/2 = 0.39 = 1.96 = 0.61 The background was subtracted through the Shirley method, and the spectra were normalised on the intensity and the binding energy scale according to the Bi6s orbital.

V 4+ V 5+
and valence band stay somewhat the same during the whole experiment. There is only a small 0.1 eV shift towards a higher binding energy after cooling the calcined BiVO 4 sample to 77 K. Because part of the extrinsic contamination layer was still present on the BiVO 4 surface after calcination, a second feature in the O1s spectrum at higher binding energy can be seen. This feature can be more clearly observed after peak fitting (Figure 5.32). This additional O1s feature has a relatively high FWHM and originates likely from oxygen containing carbonaceous species and hydroxy groups at the surface. 44 During the water adsorption a third feature appears in the O1s spectrum, which shifts to higher binding energies with increasing amounts of adsorbed water. The binding energy of this feature after the deposition of 12.8 L of H 2 O is situated at 534.1 eV, which is typical for water molecules in a multilayer film. 114,[118][119][120] The binding energy shift of this feature has generally been explained by final state screening effects, whereby a core hole in water is screened more effectively by electrons at the surface than by electrons from adjacent water molecules in a multilayer. 120 The fitted O1s components in Figure 5.32 seemingly show that the intensity of the OH+CO component with respect to the intensity of the O Binding energy (eV) Intensity (arb. units) position obtained after calcination indicates that the BiVO 4 surface contains a reasonable amount of oxygen vacancies at which water dissociation could take place. To obtain more information on the water dissociation mechanism one may try to perform the water exposure experiment on a specific nearly defect free BiVO 4 surface.

After the water dissociation and the formation of hydroxy groups, these hydroxy groups somehow seemingly react with the lattice oxygen. This is a somewhat puzzling phenomenon because reductive water dissociation, to my knowledge, has only been described before for CeO 2-x by Berner et al. 121 The decrease in lattice oxygen should be due to the formation of a volatile oxygen containing compound, which is most likely oxygen or hydrogen peroxide. In Section 2.2.2 in Chapter 2 the electrochemical formation of oxygen and hydrogen peroxide from water is described. In both reaction pathways, hydroxy group intermediates are involved. However, the formation of oxygen and hydrogen peroxide from water under standard conditions is highly endothermic, being associated with standard reduction potentials of 1.23 V vs NHE and 1.78 V vs NHE, respectively. This means that at room temperature these reactions do not take place spontaneously. The conditions of the water exposure can, however, not be considered standard, as the pressure is situated at about 10 -6 -10 -7 mbar and not the standard 1 bar. However, the de- viation from the standard conditions will most likely not alter the spontaneity of either reaction.

For example, for the oxygen evolution reaction which corresponds to a standard Gibbs energy ∆G 0 of 237.13 kJ/mol, an equilibrium constant K eq of 10 -46 can be calculated according to:

K eq = 10 G 0 /2.303RT (5.1)
K eq = 10 237130 J.K / 2.303 * 8.31 J.K -1 .mol -1 * 273.15 K (5.2)

with R the gas constant and T the temperature. This means that the vacuum conditions are likely not sufficient to drive the reaction towards the formation of oxygen. Another important notice which has to be made is that hydrogen peroxide or oxygen are not formed with two surface hydroxy groups originating from water but that at least one oxygen in the reaction originates from the lattice, which may or may not be hydroxylated. This may alter the reaction mechanisms denoted in Section 2.2.2 and, thus, the standard Gibbs energy.

Due to the water oxidation reactions not being spontaneous, another source of energy may aid the transformation of surface hydroxy groups to oxygen or hydrogen peroxide. The only source of energy (besides thermal energy) are the X-rays photons used to measure the XP spectra.

V 4+ V 5+ V 5+ I V 4+ I = 3.19 V 5+ V 5+
Binding energy (eV) Intensity (arb. units) Herein, the photoelectric effect does not only create a core hole but also leads to a cascade of secondary electrons (and holes) in follow up processes. The holes created during these follow up processes should have enough energy to oxidise the surface hydroxy groups. To confirm this pathway one may try to perform the PES measurements at 77 K after having carried out the water exposure procedure at RT. The hydroxy groups may not be oxidised at these low temperatures.

This experiment may give more information since the water desorption experiment did not lead to a reduction of vanadium or bismuth while being cooled to 77 K, meaning that neither water nor any formed hydroxy groups were oxidised at this low temperature.

Another unlikely possibility is that the upward Fermi level shift precedes the chemical reactions that involve the lattice oxygen. The Fermi level shift would lead to a higher amount of reduced bismuth and vanadium due to the reasons specified in Section 5.3.3. Namely, the Fermi level would cross the Bi 3+ /Bi 0 and V 5+ /V 4+ reduction potentials leading to the reduction of bismuth and vanadium. As a consequence of this reduction, the Bi-O-V bond may become unstable after which the lattice oxygen reacts with hydroxy groups at the surface. This mechanism is unlikely, because water normally does not donate electrons which would increase the Fermi level of the substrate. Hence, the source of injected electrons and, thus, the origin of the upward Fermi level shift, which forms the base of this hypothesis, is rather unclear.

All in all, the water exposure experiment on BiVO 4 is surprising, because it revealed that a cleaned BiVO 4 substrate is reduced after it makes contact with water. The exact reaction mechanisms which cause the BiVO 4 surface reduction upon contact with water are, however, still unclear.

Conclusion

In conclusion, we have demonstrated that the Fermi level position within BiVO 4 can be effectively tuned over a range of 1. These findings clearly show that the alignment based on solid/electrolyte interfaces cannot be used to interpret the junctions of heterostructured materials. Thus, interface experiments, which allow the direct analysis of the contact properties of newly formed junctions, are needed to understand such heterojunctions and may help to design more efficient photoelectrodes.

One of the most noteworthy observations we made was during the BiVO 4 water exposure experiment, which clearly demonstrated that an oxidised BiVO 4 thin film was reduced upon contact with water. Namely, a reduction of Bi 3+ to Bi 0 , a reduction of V 5+ to V 4+ and a reduced amount of oxygen could be shown through photoelectron spectroscopy. Furthermore, this reduction proved to be reversible as the exposure of the reduced BiVO 4 substrate to oxygen at room temperature could partly reoxidise the surface and an oxygen plasma treatment led to a complete reoxidation of the BiVO 4 surface. The reduction of BiVO 4 upon contact with water is so remarkable, because it could provide insight in the BiVO 4 based photochemical water splitting mechanism, whereby the first step would not be the absorption of light but the reduction of the substrate and the oxidation of water.

Chapter 6

Heterostructured CuFeO 2

Introduction

All the information that was gathered on how to prepare anisotropic BiVO 4 heterostructures (Chapter 4) and how to analyse the interfaces of BiVO 4 heterostructures through photoelectron spectroscopy (Chapter 5) was used to develop and analyse heterostructures of a promising photocathode material, delafossite copper iron oxide or CuFeO 2 .

As described in the main Introduction (Chapter 2) there is an immediate need to study and develop new efficient visible light active photocathode materials. Few oxides have proven to be suitable for hydrogen production through photochemical water splitting. 1 Photochemical water splitting has been demonstrated with TiO 2 and SrTiO 3 , which, however, have an intrinsically low water splitting efficiency due to their relatively wide band gap (E g > 3 eV). 2 The main problem is that oxides generally have a low VBM, owing to the O2p character of their valence band. Strategies to enhance visible light absorption, such as impurity doping or mixing in nitrogen to form oxynitrides, only led to a small improvement of the (sacrificial) hydrogen evolution efficiency. 3,4 Some oxides, however, exhibit a higher VBM due to the hybridization of metal valence states with the O2p states, as is the case for copper based oxides. [5][6][7] Indeed, due to its narrow band gap of 2.0 -2.2 eV, Cu 2 O based photocathodes have been able to reach state-of-the-art cathodic photocurrents, up to -7.6 mA cm -2 at 0 V RHE , 8 admittedly, using TiO 2 protection layers to protect Cu 2 O against photocorrosion. 8,9 A promising alternative to cuprous oxide could be the delafossite copper iron oxide due to its even lower band gap of about 1.5 eV, 10,11 good carrier mobility of about 0.2 cm 2 V -1 s -1 , 12 and its suitable CBM, at -0.1 --0.5 eV vs RHE. 1,10,12,13 However, only poor to mediocre hydrogen evolution efficiencies have been measured so far with CuFeO 2 . Jin et al. showed that a thin film of dropcasted 2H-CuFeO 2 nanoplatelets only exhibited a photocurrent of 8.8 µA cm -2 at 0 V RHE , 11 which is far below the efficiency which can be expected from the absorption spectrum, which should allow to reach photocurrents up to 15 mA cm -2 . 12 Prévot et al. demonstrated photocurrents of up to 2.4 mA cm -2 for a host-guest CuAlO 2 /CuFeO 2 system, whereby however oxygen had to be used as an electron scavenger, so that water splitting could not be assumed any more. 14 Jang et al. found a fairly reasonable photocurrent of 2.4 mA cm -2 at 0.4 V RHE , here in the absence of oxygen, for a CuFeO 2 photocathode, which was modified through hybrid microwave annealing and a NiFe-layered double hydroxide/reduced graphene oxide electrocatalyst. 15 To the best of my knowledge, no higher photocurrents or hydrogen evolution efficiencies have been reported up to now. Another remarkable conclusion is that unbiased hydrogen evolution with CuFeO 2 has also not yet been achieved. Several reasons have been suggested to explain why CuFeO 2 have showed low hydrogen evolution efficiencies so far, which include Fermi level pinning at inter band gap states in the bulk, 16 Fermi level pinning at surface states, 12 photocorrosion 10 and poor catalytic activity. 15 In this Chapter the development of anisotropic CuFeO 2 heterostructures will be described as a strategy to improve the hydrogen evolution efficiency through optimization of the charge carrier separation and by including more performant hydrogen evolution electrocatalysts in contact with the electrolyte. Specifically, the 2H-CuFeO 2 polytype was chosen, because its hexagonal lattice system allows the preparation of hexagonal CuFeO 2 nanoplatelets with different exposed facets. 11 The upscaling of the CuFeO 2 nanoplatelet synthesis procedure will be detailed in this Chapter, as well as the exploration of regioselectivity of the photodeposition of platinum, silver and nickel(hydroxide). In the second part of the Chapter, the tunability of the Fermi level position in CuFeO 2 is explored by contacting water and ITO in individual interface experiments with 2H-CuFeO 2 pellets.

Synthesis hexagonal CuFeO 2 nanoplatelets

2H-CuFeO 2 nanoplatelets were synthesized according to a recently published synthesis procedure (See Section 3.1.3 for a thorough description). 11 Considering that the original synthesis procedure only yielded about 75 mg per synthesis, we tested whether the reaction parameters could be upscaled by a factor of 4 and a factor of 8, so that a batch of 6 g could be produced in a more timely manner. First, the hydrothermal treatment of the reaction mixture was tried in a one-way sealed glass vessel, which was however not a good option, because the glass vessel exploded, in the protective container where it was put in, due to pressure built-up and the very basic conditions of the reaction mixture, that probably led to glass corrosion. In the next try, the glass vessel was replaced by a Teflon lined stainless steel pressure reactor, capable of handling elevated pressures and possessing a high chemical inertness due to the Teflon lining. As can be seen in the secondary electron images in Figure 6.1a, the 4x an 8x upscaling of the reaction parameters led to a powder of which the particles still look hexagonally shaped. Furthermore, the powder X-ray diffractograms (Figure 6.1b) show that the powder almost completely consists of 2H-CuFeO 2 with 3R-CuFeO 2 being detected in only a minor amount. Another noticeable feature in the diffractograms is the higher than normal background, which was caused by X-ray fluorescence. The energy of the X-rays emitted by the copper anode (Cu Kα: 8046 eV) of our lab diffractometer can be absorbed by the iron atoms (Fe K-edge: 7112 eV) of the CuFeO 2 sample, which leads to fluorescence after radiative de-excitation. X-ray fluorescence can be easily overcome by using another X-ray anode with lower characteristic emission energies (such as Fe or Cr) or by installing a second monochromator before the detector. As such, the 4x upscaled powder was measured again with another PXRD apparatus that had a monochromator before the detector (Figure 6.2). In this diffractogram the background is greatly reduced and the presence of 3R-CuFeO 2 is more pronounced. The FESEM images of the 4x upscaled CuFeO 2 powder in Figure 6.3 show the hexagonal nanoplatelet morphology more clearly. The short diagonal length is ± 600 nm and the thickness lies in the range of 100 -200 nm. However, the exposed faces appear to be much rougher than those of the truncated bipyramidal BiVO 4 microcrystals, which can complicate the introduction and detection of fine deposits.

At first glance, the overall characteristics of the upscaled CuFeO 2 powder were quite satisfactory, so that the CuFeO 2 production with the 8x upscaled reaction parameters was continued.

Further upscaling was not considered, because the 125 mL stainless steel pressure reactor had Intensity (arb. units) reached its filling limit.

With the 8x upscaled parameters, each synthesis yielded about 600 mg CuFeO 2 , so that in principle the reaction had to be carried out 10 times to create a batch of 6 g. However, due to problems with reproducibility the synthesis had to be performed 25 times. In all cases monodisperse hexagonal CuFeO 2 particles were obtained. However, some samples seemingly exhibited stronger particle aggregation/agglomeration, which could hinder studying the regioselectivity of the photodeposition of additional materials, because of increased concealing of certain particle surfaces. Furthermore, particle aggregation may increase the chance of charge carrier recombination, which would have a negative impact on water splitting efficiencies. Figure After all the suitable CuFeO 2 powders were homogeneously mixed into one large batch of ± 800 mg, the whole batch was characterized again. The high resolution powder X-ray diffractogram in Figure 6.6 surprisingly shows that only a minimal amount of 3R-CuFeO 2 seems to be present in the batch. Furthermore, the FESEM images in Figure 6.7, taken from a sample of the CuFeO 2 batch, again show the hexagonal nanoplatelet morphology and, more remarkably, demonstrate a much lower surface roughness of the exposed CuFeO 2 nanoplatelet faces compared to the sample that was upscaled 8x times (Figure 6.3). However, there seems to be still a deposit present with a particular morphology on some of the nanoplatelets. Auger electron spectroscopy (AES) was performed on these samples to figure out if the additional deposits had another composition, but no differences could be found. To study the regioselectivity of the photodeposition process it would be good to know in the satellite structure, Fe 2+ gives rise to a signal at about 716 eV, whereas Fe 3+ give rise to a signal at about 719 eV. 23 In the Fe2p spectrum in Figure 6.9 no satellite signal can be seen around 716 eV only a shoulder at about 720 eV is present, thus indicating that Fe 3+ is the main oxidation state in the analysed CuFeO 2 sample. Peak fitting of the Fe2p spectrum would be the ideal way to assign the iron oxidation state, however the noise and the possibility of surface charging does not allow proper peak fitting. In the Fe2p spectrum a signal can be seen at a binding energy of 705 eV. This signal corresponds to In3p 1/2 24 and originates from the indium foil used to glue the powder sample to the substrate holder. In the O1s spectrum signals belonging to oxide (530 eV) and hydroxide/carbonyl (532 eV) can be seen. 24 The hydroxide/carbonyl signal is most likely due to the sample having been prepared ex situ.

Anisotropic photodeposition on CuFeO 2

Once the CuFeO 2 batch was fully analysed, samples were taken of this batch to perform photodeposition experiments to test whether certain metal(oxide)s could be selectively deposited on certain regions (regioselectivity) of the anisotropic hexagonal CuFeO 2 nanoplatelets. The photodeposition procedure is extensively explained in Section 3.1.4 and specific experimental parameters of the CuFeO 2 photodeposition experiments can be found in Table 6.1.

CuFeO 2 /Ag

Silver has been effectively deposited onto the CuFeO 2 nanoplatelets following the photodeposition process, as can be perceived in Figure 6.10. Evidently, more silver is present on the {0001} 6.11). On positions, where in the FESEM images a Ag deposit seems to be present (more brilliant spots), silver Auger electrons were detected, confirming that silver was effectively deposited onto the CuFeO 2 nanoplatelets. The number of silver deposits is, however, so sparse that the regioselectivity of the deposit could not be verified with high certainty from these AES point measurements. The silver oxidation state of the CuFeO 2 /Ag (10 wt%) sample was determined through XPS (Figure 6.12). First, the Fe2p, Cu2p and CuLMM core level spectra were checked for inhomogeneous sample charging, as this phenomenon made the oxidation state analysis for some BiVO 4 based powder samples (see Chapter 4) more complicated. However, the charge carrier mobility of CuFeO 2 , 10 -1 cm 2 V -1 s -1 , is one order of magnitude higher than for BiVO 4 , 12,25

CuFeO 2 /Pt

The photodeposition of platinum was next in line. For BiVO 4 , the acidic H 2 PtCl 6 precursor was found to change the pH of the photodeposition reaction mixture (Section 4.3.2). Moreover, the pH and platinum precursor concentration had a strong effect on the Pt distribution and oxidation state. In particular, a reaction mixture pH of 2.5 -3 resulted predominantly in metallic platinum deposits, whereas at higher pHs oxidized platinum species were preferentially deposited. Hence, for the photodeposition of platinum on CuFeO 2 the pH was not altered with sodium hydroxide and was preferentially held between 2.5 -3. After making a 10 wt% platinum solution a pH of 2.9 was obtained. As can be seen in Figure 6.13 the size of the platinum deposits after the photodeposition process are in the range of 10 -50 nm. Furthermore, the platinum coverage is also much higher than what was obtained for the silver photodeposition on CuFeO 2 .

Even more remarkably, the Pt deposit seems to be highly selectively deposited onto the {0001} facets, which constitutes a strong proof that the anisotropic shape of the CuFeO 2 nanoplatelets and holes accumulate on the {010} and {110} crystalline facets, respectively, making that reduction and oxidation reactions are spatially separated and take place on different crystalline facets. [26][27][28] Since then, the light-induced charge migration has also been demonstrated for anisotropically shaped SrTiO 3 and TiO 2 . 29,30 The particular crystal structure of CuFeO 2 and the proposed small polaron hopping charge carrier conduction mechanism 25,31 could be the reasons for regioselective charge accumulation. The CuFeO 2 crystal structure consists of alternating layers of [FeO 2 -] and [Cu + ] which are stacked perpendicular to the c-axis. When assuming that charge carriers follow the small polaron hopping mechanism, hole hopping occurs presumingly between copper sites (forming Cu 2+ ions) and electron hopping between iron sites (forming Fe 2+ ions). 25 In a simplistic electrostatic model, electrons must hop from iron to iron site over a saddle type configuration of oxygen anions, due to the iron ions being octahedrally coordinated by oxygen ions.

In contrast, the copper cations are, in line with the c-axis, coordinated by two oxygen anions, whereas no oxygen anions, and thus no barrier, are present in the [Cu + ] layers perpendicular to the c-axis, so that a higher hole mobility perpendicular to the c-axis would be expected. Indeed, electrical conductivity measurements on p-type CuFeO 2 single crystals demonstrated that the resistivity along the c-axis, 3 x 10 1 Ωm, was several orders of magnitude higher than perpendicular to the c-axis, 5 x 10 -3 Ωm. 32 Besides anisotropic charge accumulation, the reason for regioselective deposition of platinum on the CuFeO 2 nanoplatelets can possibly be due to a difference in surface polarity between the {0001} facets and the sides of the nanoplatelets. The {0001} facets should then be more positively charged than the sides, so that the [PtCl 6 ] 2-precursor species preferentially adsorb onto the {0001} facets, followed by reduction of the precursor to Pt(0). This alternative pathway of anisotropic deposition has already been documented by Wenderich et al., who showed that [PtCl 6 ] 2-preferentially adsorbed onto the small subordinate facets of platelike WO 3 through a comparison of dark impregnation and photodeposition experiments, and aqueous-phase atomic force microscopy (AFM) measurements. 33 Thus, in principle, dark impregnation experiments could be performed to determine whether platinum was selectively deposited onto the {0001} facets through lightinduced charge carrier migration or through facet-preferential [PtCl 6 ] 2-adsorption. However, the impregnation of platinum on the 2H-CuFeO 2 nanoplatelets could not be performed, because the impregnation process requires a heating step in a controlled atmosphere to convert the precursor to the desired deposit with the desired oxidation state and, unfortunately, the 2H-CuFeO 2 phase and hexagonal nanoplatelet morphology was not stable at the elevated temperature, i.e. 400 ℃, that is required for metal(oxide) impregnation (which will be shown in Section 6.5.1).

Nonetheless, the observation that the positively charged Ag + precursor is also selectively reduced and deposited onto the {0001} facets of 2H-CuFeO 2 hexagonal nanoplatelets speaks against the sorption-determined platinum deposition hypothesis.

The deposit, found on the CuFeO 2 particles, is indeed platinum, because Pt M 5 VV Auger electrons were detected when an AES point measurement was performed on a point which was assumed to be a deposit in the recorded secondary electron image (Figure 6.14).

XP core level spectra were measured to verify the platinum species (Figure 6.15). No signs of inhomogeneous sample charging were observed, so that the Pt4f spectrum could be used to analyse the oxidation state of the platinum deposit. Two intense signals at 71.0 eV and 74.4 eV can be seen which are characteristic for the binding energies of Pt4f 7/2 and Pt4f 5/2 , respectively, of metallic platinum. 24 Additionally, a shoulder can be seen at a binding energy of 77.4 eV, which is indicative of Pt4f 5/2 of platinum oxide 24 As was discussed for the photodeposition of platinum onto BiVO 4 (Section 4.3.2), platinum oxide can be formed when [PtCl 6 ] 2-is only partially reduced.

CuFeO 2 /NiO x OH y

Next, the photodeposition of nickel was investigated. The FESEM images (Figure 6.16) show that the deposited nickel particles have a distinct morphology, yielding sizes of just a few nanometer and with part of the particles being seemingly elongated. Nonetheless, the nickel deposit is slightly more difficult to identify due to the strong similarity with the deposit that was found on AES was performed to determine whether all of the supposed deposits were containing nickel.

However, the most intense nickel Auger lines overlap with either the Auger signals of iron or copper, and the less intense nickel Auger lines do not have a high enough signal-to-noise ratio.

Hence, nickel could not be detected through AES. EDS was not considered due to the small size of the deposits.

The reason that BiVO 4 and CuFeO 2 are evidently completely covered by photodeposited nickel could be that, at a pH of 9.2, regioselective deposition of nickel is inhibited, due to the surface polarity being the same on every exposed facet at that pH, suggesting that the deposition of nickel is sorption-determined. Another possibility is that Ni 2+ reduction and oxidation takes place simultaneously. In Section 4. using a 0.1 wt% nickel solution would be able to be determined. Due to the similarity of the nickel deposit for both CuFeO 2 and BiVO 4 , an alternative reaction mechanism which does not involve light can not be excluded, despite no alternative mechanisms having been mentioned in literature, as far as we know.

XPS measurements were performed to investigate the oxidation state of the deposited nickel species (Figure 6.17). Again, no signs of inhomogeneous surface charging can be found in the core level spectra, so that the deposited nickel species can be evaluated from the Ni2p spectrum.

UV-Vis-NIR diffuse reflectance spectroscopy

For all heterostructured CuFeO 2 powders, UV-Vis-NIR diffuse reflectance spectra were recorded (Figure 6.18a). Subsequently, the diffuse reflectance data were transformed into Tauc plots to determine the indirect and direct optical band gaps of the different heterostructures (Figure 6.18b).

The indirect and direct optical band gaps of the heterostructured CuFeO 2 powders are similar to bare CuFeO 2 and apparently lie within a range of 1.22 -1.26 eV and 1.43 -1.45 eV, respectively.

The direct band gap values are close to the previously determined direct optical band gap of 2H-CuFeO 2 . 11 Thus, eventual deviations in sacrificial water splitting between the different CuFeO 2 powders will not be caused by differences in light absorption. Furthermore, due to the 0.2 eV difference between the direct and the indirect band gap, radiative charge carrier recombination will be less likely.

Water reduction

All the CuFeO 2 heterostructures were tested for their sacrificial photochemical water splitting performance. The complete testing procedure is explained in detail in Section 3.4.5.2. The heterostructured CuFeO 2 nanoplatelets were expected to act as good sacrificial water splitting systems, because of the straightforward photodeposition of silver and platinum, the apparent spatial separation of photogenerated charge carriers, and the good electrocatalytic properties of the deposited cocatalysts; platinum for hydrogen evolution 34 and nickel based species for oxygen evolution. 35 However, none of the heterostructured CuFeO 2 powders showed any significant water splitting performance, even though two different sacrificial agents, methanol and triethanolamine, and two different weight percentages for every heterostructure were tested. Even when the reaction times were extended to 72 hours or when UV lamps were used instead of Xenon lamps, no gas evolution could be noticed by the automatic burette of the closed batch reactor. Nonetheless, gas samples were taken from the reactor for GC analysis after the hydrogen evolution experiments were terminated. Hydrogen could be detected for the CuFeO 2 /Ag and CuFeO 2 /Pt powders.

However, the hydrogen GC signal was so low that it did not reach the calibration range of the GC apparatus, so that the amount of evolved hydrogen was considered to be below 1 µmol/h. This unsatisfactory water splitting efficiency is in accordance with the previously determined poor photocurrent, 8.8 µA cm -2 at 0 V RHE , for a thin film of dropcasted 2H-CuFeO 2 nanoplatelets. 11 According to Prévot et al. Fermi level pinning at surface states, due to formation of metal hydroxides at the semiconductor electrolyte junction, impairs the PEC water splitting performance of CuFeO 2 . 12 In the following Section additional reasons for the bad hydrogen evolution efficiency of CuFeO 2 are explored through interface experiments of CuFeO 2 with the low work function material ITO and through water exposure experiments.

6.5 Interfacial band alignment of heterostructured CuFeO 2

CuFeO 2 pellet fabrication

Before any interface experiments could be performed, suitable CuFeO 2 substrates had to be prepared onto which ITO could be sputtered and for which X-ray photoelectron spectra could be measured. To our knowledge, no experimental procedures have been published for the fabrication of 2H-CuFeO 2 thin films so far, only some examples of 3R-CuFeO 2 thin films could be found in literature. 16,[36][37][38] Exploring the preparation of 2H-CuFeO 2 thin films ourselves seemed too complex due to the stringent conditions, high temperature and specific atmosphere, required to produce phase pure delafossite CuFeO 2 thin films 36,37 and because of the privileged growth of the 3R-CuFeO 2 phase when fusing copper and iron precursors. Thus, rather than preparing thin films, 2H-CuFeO 2 substrates were prepared by pressing a certain amount of previously prepared 2H-CuFeO 2 powder into pellets. The pellets were pressed with the 2H-CuFeO 2 powder that was discarded earlier because of its unsuitable properties for photodeposition and water splitting experiments (See Section 6.2). Obviously, pellets are a few order of magnitudes thicker than thin films, but inhomogeneous surface charging during XPS measurements was not expected due to the good electrical conductivity of CuFeO 2 and because the top surface of the pellets is grounded through a mask, which keeps the pellet in place. 2H-CuFeO 2 pellets were prepared through a combination of uniaxial and isostatic pressing. The complete synthesis procedure can be found in Section 3.3. Normally, to obtain compact CuFeO 2 pellets, sintering at temperatures above 1000 ℃ is required. 39,40 However, fully compacted pellets did not seem imperative for the execution of interface experiments and, more importantly, the 2H-CuFeO 2 crystalline phase was found to be not stable at temperatures > 400 ℃. As can be seen in Figure 6.19, the 2H-CuFeO 2 crystalline phase completely changed into Fe 3 O 4 and CuO after heating for 2 h in air at 400 ℃. Moreover, the hexagonal nanoplatelets seem to fragment into smaller particles during the calcination (Figure 6.20), which is most likely due to the phase change that takes place.

Since the pellets were made ex situ, the surface had to be cleaned before interface experiments could be performed. The survey XP spectra in Figure 6. Additionally, Cu2p, Fe2p, O1s and VB XP core level spectra and CuLMM Auger spectra were recorded for the different CuFeO 2 cleaning procedures (Figure 6.22). These measurements indeed confirm that the oxygen plasma was a too oxidative treatment, because in the Cu2p spectrum a broadening of the main Cu2p 3/2 line and an intensity increase in the satellite area around 940 -945 eV can be observed, which are both typical features for copper in the Cu 2+ oxidation state. 22 Furthermore, in the CuLMM spectrum a broadening of the main line can be noticed, which is most likely related to the presence of Cu 2+ , because the Cu 2+ CuLMM spectrum is usually shifted towards a lower binding energy compared to the CuLMM spectrum of Cu + . 22 In the Fe2p spectrum no noticeable change can be seen, mainly because the oxidation of Fe 3+ to Fe 4+ is unlikely, due to Fe 3+ being a more stable oxidation state for iron containing solid state compounds. For the heat treatment in oxygen atmosphere, the formation of Cu 2+ could not be observed and also the Fe2p region remains largely the same. However, the CuFeO 2 pellet, before treatment, exhibited a bump at a binding energy of 935 eV in the Cu2p spectrum, which disappears after the thermally cleaning step at 200 ℃ Furthermore, in the O1s region the shoulder at 531.6 eV, characteristic for hydroxyl and carbonyl groups, disappears after the thermally cleaning step, indicating that the oxygen heat treatment completely oxidises the surface, removing all hydroxides and carbonaceous species. Another observation is that, after the heat treatment, the valence band maximum shift from 0.41 eV to 0.19 eV, indicating that CuFeO 2 becomes more p-type.

As mentioned above, the 2H-CuFeO 2 pellets were made from the CuFeO 2 powder which was not appropriate for the photodeposition experiments, i.e. the powder which was more aggregated and which contained more 3R-CuFeO 2 . In Figure 6.23 the X-ray diffractograms of the powder, used to prepare the pellets, and the pellet after the 200 ℃ thermally cleaning step are shown.

Surprisingly, not much 3R-CuFeO 2 seems to be present in the powder. Moreover, 3R-CuFeO 2 could not be detected any more after pressing the powder into a pellet and cleaning it through the oxygen heat treatment, which may indicate that the 3R-CuFeO 2 phase converts into the 2H-CuFeO 2 at the used cleaning conditions. Another remarkable feature is the relatively higher intensity of the {001} reflections for the pellet, which most likely originate from the stacking of the nanoplatelets according to the [0001] direction during uniaxial pressing. This preferential stacking could also be the reason that 3R-CuFeO 2 cannot be observed any more in the θ -2θ scan of the pellet. Additionally, the [0001] orientation of the grains in the pellets means that the interface experiments on these pellets mostly yields information on how the {0001} surfaces of 2H-CuFeO 2 align with the contact material. can be raised is also more relevant for CuFeO 2 based hydrogen evolution compounds, because then information can be gained on the extent of feasible downwards band bending and an idea can be conceived on whether photovoltage will be limited due to bulk electronic processes. The sputtering parameters of ITO can be found in Section 3.2.2. The details of an interface experiment are described in Section 3.4.4.1. Figure 6.24 shows how the Cu2p, Fe2p, CuLMM, O1s

CuFeO 2 /ITO

and VB spectra evolve with increasing ITO coverage. The developing CuFeO 2 /ITO interface can be clearly followed, as the Cu2p, Fe2p and CuLMM spectra decrease and the In3d 5/2 intensity increases with increased ITO sputtering time. Meanwhile, the O1s intensity remains constant.

The shift of the Cu2p, CuLMM, Fe2p and valence band spectra towards higher binding energies can be noticed with increased ITO coverage, which indicates an electronic, rather than a chemical, shift. As E F -E VBM of bare CuFeO 2 is quite low, the ITO valence band spectrum does not overlap with that of CuFeO 2 , so that the shift in the CuFeO 2 VBM can be used to quantify the electronic shift. Quantifying the shift by following the shift in the CuLMM and Fe2p 3/2 is more complicated, because of the asymmetric line shape in both spectra and, additionally, the large number of satellites and mixing in of In3p 1/2 and Sn3p 3/2 in the Fe2p 3/2 spectra. The Cu2p 3/2 binding energy shift can be followed with ease due to its relatively strong intensity and well-defined Intensity (arb. units) 2theta (°) marginally from that of metallic copper, and the Cu2p 3/2 line of Cu 2+ yields only a slight shift and broadening compared to Cu + . 22 However, the broadening of the principal Cu2p 3/2 line, seen here, most likely does not correspond to an oxidation of Cu + to Cu 2+ , because then an increase in intensity in the satellite region between 940 eV and 945 eV, characteristic for Cu 2+ , should be seen as well. 22 Furthermore, the CuLMM Auger spectrum after 4 s of ITO deposition does not show any sign of Cu 2+ either, as then a signal should appear at a 0.9 eV lower binding energy than the main Cu + CuLMM line. 44 In addition, a change from Cu 2+ to Cu + is not consistent with the observed increase in binding energies. Hence, the broadening of the main Cu2p 3/2 line is not likely due to a change in copper oxidation state. Another possible reason for the increased broadening of the main Cu2p 3/2 line is inhomogeneous ITO coverage due to the large substrate surface roughness, which is plausible for a powder pressed pellet. In this case, the Cu2p 3/2 line at 4 s ITO coverage would then be a composite signal consisting of Cu2p 3/2 signals with different binding energy shifts, originating from positions with differing ITO thickness. Indeed, surface profiles show that the 2H-CuFeO 2 pellet surface roughness appears to be an order of magnitude higher compared to that of a BiVO 4 thin film (Figure 6.27). Thus, surface roughness seems to be a possible explanation for the observed Cu2p 3/2 broadening. However, the CuLMM and Fe2p 3/2 spectrum apparently do not suffer from line broadening, which would be expected if surface roughness was the cause of Cu2p 3/2 line broadening. Therefore, the exact cause of Cu2p 3/2 line broadening could not be confidently determined.

Furthermore, an extra signal at lower binding energy, i.e. around 709 eV, appears in the Fe2p 3/2 difference spectra after 4 s of ITO deposition (Figure 6.26). Since this binding energy position is characteristic for Fe 2+ , 45 Fe 3+ is likely partly reduced to Fe 2+ , as effect of the ITO coverage.

Moreover, in our group the appearance of Fe 2+ in the Fe2p 3/2 spectrum has been observed before for hematite, 46 lanthanum ferrite (unpublished results) and bismuth ferrite (unpublished results), while increasing the Fermi level position in these materials through ITO deposition.

The results of the CuFeO 2 /ITO interface experiment are summarized in Figure 6 CBM at 0.1 -0.5 eV. 1,10,12,13 Additionally, the apparent 0.32 eV Fermi level shift, as determined from the ITO interface experiment is shown. In Figure 6.28c a series of standard reduction po- ± -0.7 eV vs RHE (pH 7), which the authors assumed to be caused by surface states. 12 However, as our results suggest Fermi level pinning does not take place because of surface states, but due to the bulk Fe 3+ /Fe 2+ polaron level at -0.7 eV vs RHE. Bulk Fermi level pinning can, thus, strongly impair the potential water reduction capacity of CuFeO 2 , because then not enough photovoltage can be created to achieve water reduction and no solution could be engineered because the photovoltage limit would be intrinsic to CuFeO 2 . The -0.7 eV vs RHE Fermi level limit could also explain why PtCl 6 2-, with a reduction potential of -0.75 eV vs RHE, and Ag + , with a reduction potential of -1.02 eV vs RHE, could be reduced, upon light absorption, whereas water, with a reduction potential of 0 eV vs RHE, could not. Only the nickel photodeposition could disprove the hypothesis of bulk Fermi level pinning at the Fe 3+ /Fe 2+ polaron level, as the Ni 2+ /Ni reduction potential is situated at about -0.13 eV vs RHE. However, due to the absence of regioselectivity in the nickel deposition (Figure 6.16), a dark nickel impregnation mechanism could not be excluded.

Fermi level tuning was further investigated through water exposure experiments, which will be discussed in the next Section.

i The RHE scale is chosen because the CuFeO 2 band positions and, the H + /H 2 and O 2 /H 2 O reduction potentials are then independent of pH (See also Figure 2.18). However, the reduction potentials of the other listed reactions, although they do not involve any hydroxide anions or hydrogen cations, shift with pH on the RHE scale, due to the pH dependence of the reversible hydrogen electrode with respect to the standard hydrogen electrode. These reduction potentials are then calculated according to the modified Nernst equation: E RHE = E 0 SHE +0.059pH + 0.059 n loд(Q ini ); with the E 0 SHE taken from literature, 47 n the number of electrons involved in the reduction reaction and Q ini the initial reaction quotient. The standard reduction potentials of PtCl 6 2-/PtCl 4 2-and Ag + /Ag(s) are calculated for pH 3 and pH 7, respectively, because the photodeposition of platinum and silver were performed at pH 3 and pH 7, respectively. The Fe 3+ /Fe 2+ polaron level, situated at -0.29 eV vs SHE as was recently proven by Lohaus et al., 46 and reduction potentials of Cu + /Cu(s) and Fe 3+ /Fe 2+ are calculated at pH 7, because the bare CuFeO 2 pellet, in the absence of a surface dipole (see above) and under flat band conditions, is regarded as being at the pH of zero point charge, which is around pH 7. The experimental procedure of a water exposure experiment can be found in Section 3.4.4.2.

Before exposing the CuFeO 2 samples to water, they were calcined in a pure oxygen atmosphere to remove the extrinsic contamination layer. However, the calcination was not effective at removing the upper contamination layer, since the intensity of the C1s signal in the survey spectrum (Figure 6.29) does not decrease after the calcination procedure. The incomplete removal of the extrinsic carbonaceous species could be due to the low temperature that had to be used in the calcination procedure or due to the porosity of the pellets. In Figure 6.30 the effect of water exposure on the Cu2p 3/2 , Fe2p, O1s and valence band XP and CuLMM Auger spectra of CuFeO 2 is shown.

All spectra shift towards higher binding energies, as can be expected from the contact with a low work function material. 49 The determined valence band maxima show that water exposure had a stronger effect than ITO as here the VBM is pushed towards 0.7 -0.8 eV. Admittedly, the VBM of pellet 1 before water exposure was already positioned at a quite high position. In the O1s spectra an additional feature at a binding energy of 532 eV can be seen, which is characteristic for carbonyl bonds and hydroxides. 24 After water exposure this signal at 532 eV increases, indicating that hydroxides were formed during water exposure. This hypothesis is also seemingly confirmed by the valence band region in the UPS spectra. After water exposure, features at 5.3 eV, 7.4 eV and 11.2 eV can be seen, which are characteristic to the 1b 1 , 3a 1 and 1b 2 orbitals, respectively, of water, only shifted by about 2 eV. 50 The shift is likely due to the water being physisorbed to the CuFeO 2 substrate, as a 1 eV shift with respect to molecular water was seen earlier for physisorbed water on NiO. 51 Binding energy (eV) Intensity (arb. units) Difference spectra of Cu2p 3/2 , Fe2p 3/2 , CuLMM and Fe3p (Figure 6.32) were generated to verify whether the exposure of CuFeO 2 to water led to any chemical changes. In the difference trace of the main Cu2p 3/2 signal an apparent line narrowing can be seen. Possibly, copper may get reduced from Cu 2+ , which could be present in small amounts, to Cu + , which is however unlikely as no changes are noticed in the Cu2p 3/2 satellite region. 22 Another possibility is the reduction from Cu + to Cu(0), which is also unlikely since then an extra signal at lower binding energy in the CuLMM spectra should appear. 22 Hence, just like for the line broadening observed after ITO deposition no clear explanation can be deduced for the main Cu2p 3/2 line width change. Both Fe2p 3/2 and Fe3p show that an extra signal appears at lower binding energy, which indicates the presence of Fe 2+ . 45,52 The Fe 2+ signal is also somewhat more intense than what was seen for the CuFeO 2 /ITO interface.

Again, the upwards Fermi level shift and coinciding reduction of Fe 3+ to Fe 2+ support the hypothesis of Fermi level pinning at the formation and occupation of the Fe 3+ /Fe 2+ polaron level at 0.7 eV vs RHE, which was introduced after the results of the BiVO 4 /ITO experiment were analysed.

It is surprising that Fermi level pinning seems to take place and Fe 2+ appears at an E F -E VB of 0.5 eV during ITO deposition and at 0. at the Fe 3+ /Fe 2+ polaron formation level, the E F -E VB positions should normally be the same for both experiments.

Conclusion

2H-CuFeO 2 hexagonal nanoplatelet heterostructures were effectively created through photodeposition. Specifically, silver, platinum and nickel(hydroxide) were deposited, whereby, silver and platinum were selectively found on the {0001} facets of the 2H-CuFeO 2 nanoplatelets. No selectivity was found for the nickel(hydroxide) deposits.

Despite the creation of anisotropic heterostructures with performant hydrogen evolution electrocatalysts, i.e. CuFeO 2 /Pt, only trace amounts of hydrogen were detected for the heterostructured CuFeO 2 powders after testing the photochemical sacrificial water splitting activity of the powders in a closed batch reactor coupled to an automatic burette.

CuFeO 2 substrates were prepared, by pressing CuFeO 2 powder into pellets, to perform interface experiments with ITO and water to elucidate why the CuFeO 2 powders showed a poor hydrogen evolution efficiency, while still being active for the photodeposition of Ag, Pt and NiO x OH y . The interface experiments showed that the Fermi level tunability was limited, with Schottky type barriers amounting to 0.5 and 0.8 for the CuFeO 2 /ITO interface experiment and the CuFeO 2 /H 2 O water exposure experiments, respectively. Additionally, for both type of experiments, the reduction of Fe 3+ to Fe 2+ could be observed in the Fe2p and Fe3p core level XP spectra.

These results indicate that for CuFeO 2 bulk Fermi level pinning takes place at the Fe 3+ /Fe 2+ elec-tron polaron level. Therefore, the photovoltage achievable with CuFeO 2 seems to be limited, inhibiting unbiased photochemical hydrogen evolution. Should this be true, no optimization would allow CuFeO 2 based photocathodes to even come close to its theoretically allowed photocurrent of 15 mA cm -2 in the absence of a bias, due to the intrinsic nature of the bulk Fermi level pinning phenomenon.

The results obtained in this Chapter also suggest that, when selecting future promising photocathode materials, the energetic positions of polaron levels have to be taken into account.

Conclusions and Perspectives

Concluding statements

In the broad context of the development of efficient materials for the renewable production of solar fuels, this work deals with the synthesis and interface analysis of heterostructured oxide photoabsorbers for the purpose of photochemical water splitting. We used bismuth vanadate (BiVO 4 ) and copper iron oxide (CuFeO 2 ) as main photoabsorbers for the water oxidation and water reduction, respectively.

We have shown, using a combination of electron microscopy techniques (SAM, TEM, (FE)SEM), that photodeposition was a reliable ways of creating anisotropic heterostructures.

Namely, metallic silver (Ag) and cobalt (oxy)(hydr)oxide (CoO x OH y ) could be selectively deposited onto the {010} and {110} facets, respectively, of a {010}/{110} dual facet exposed BiVO 4 microcrystalline powder. Metallic platinum (Pt) could evidently be deposited as well on the {010} facets, under the condition that the right pH, 2.5 -3, was used during the photodeposition procedure. Furthermore, the regioselectivity of the nickel (oxy)(hydr)oxide photodeposition was investigated for the first time. For the photodeposition of ± 10 wt% and 1 wt% NiO x OH y no selectivity could be obtained, whereas a 0.1 wt% NiO x OH y has apparently been deposited onto the {010} facets and the edges between the {110} facets in a preferential manner.

Another novelty in this work was the creation of heterostructured CuFeO 2 powders through the photodeposition method. Here, we successfully demonstrated that Ag and Pt could be selectively deposited onto the {0001} facets of the 2H-CuFeO 2 hexagonal nanoplatelets. Also, NiO x OH y could be effectively deposited onto the CuFeO 2 powder, however NiO x OH y seemed to be present over the entire CuFeO 2 particles, thus showing no particular regioselectivity. A possible reason for the non-selective NiO x OH y photodeposition on the CuFeO 2 nanoplatelets could be the simultaneous oxidation and reduction of the Ni 2+ ion.

The photocatalytic performance of the (heterostructured) powders was tested through sacrificial water splitting for which a closed batch reactor coupled to an automatic burette was used. In addition, the dye degradation performance of the (heterostructured) BiVO 4 powders was tested.

The results showed that a BiVO 4 heterostructured photocatalyst which performs well at dye degradation does not necessarily exhibit a good photochemical oxygen evolution rate. Quali-tatively, the photochemical water photo-oxidation performance of the heterostructured BiVO 4 powders are similar to the results of other researchers. However, with respect to those previously published results, we found a 21x higher absolute oxygen evolution rate with bare BiVO 4 , which could be due to our BiVO 4 powder having a higher quality. Unfortunately, the NiO x OH y contact material that we introduced for the first time on particulate BiVO 4 only yielded a 2x higher oxygen evolution rate compared to bare BiVO 4 , whereas Pt and CoO x OH y exhibited a 5x and 10x

higher water photo-oxidation rate, respectively. The low oxygen evolution rate with NiO x OH y as contact material is surprising, because NiO x OH y contact layers previously showed to have a positive effect on the water oxidation performance of BiVO 4 thin films. The low regioselectivity of the NiO x OH y deposit with respect to the other investigated contact materials, is the likely reason why we did not observe a high oxygen evolution rate with NiO x OH y as contact material.

The most disappointing result was obtained when testing the sacrificial water reduction performance of the heterostructured CuFeO 2 powders, because only trace amounts of hydrogen were found, even after prolonged reaction times and the use of highly intense UV light. Based on previous literature results, we assumed that the rather poor photochemical hydrogen evolution was due to Fermi level pinning which limits the photovoltage that can be generated with CuFeO 2 .

Through so-called interface experiments, whereby a thin film contact was stepwise deposited onto a photoabsorber substrate, the Fermi level tunability within CuFeO 2 and BiVO 4 was determined as well as the electronic and chemical structure of the in situ formed heterojunctions. For BiVO 4 we showed that the Fermi level could be tuned over a wide range of 1.1 eV, which may explain why high photovoltages have been reached previously with BiVO 4 based photoanodes.

Especially, by contacting BiVO 4 with CoO x , RuO 2 and NiO, which are materials with high work functions, values as low as 1.0 eV could be obtained for E F -E VBM . Thus, by contacting BiVO 4 with high work function materials, high upwards band bending can be reached. As a consequence enhanced charge carrier separation can be expected for these heterostructures. In contrast, when BiVO 4 thin films were contacted with ITO, a low work function material, a limited increase in E F -E VBM , up to 2.1 eV, was found, together with a reduction of vanadium and also bismuth when ITO was deposited at 400 ℃. These results indicate that Fermi level pinning occurs, which may be due to polaron formation.

To perform interface experiments with CuFeO 2 , the synthesized CuFeO 2 powder was pressed into pellets. These CuFeO 2 pellets were only applied for interface experiments with ITO, because the already very low 0.2 eV value of E F -E VBM , obtained for the p-type pellets, was not expected to change through contacts with high work function materials. Remarkably, only a rather low E F -E VBM increase, up to 0.5 eV, could be obtained during the ITO interface experiments, which suggests that Fermi level pinning likely occurs. Indeed, by taking a closer look to the Fe2p core level spectrum, indications of the formation of Fe 2+ were found. Furthermore the position of the Fe 3+ /Fe 2+ transition corresponds well to what others found in our group for other iron containing compounds. Thus, the reason why we did not observe any significant water splitting with any of the heterostructured CuFeO 2 powders is very likely because CuFeO 2 cannot reach high photovoltages due to bulk Fermi level pinning at the Fe 3+ /Fe 2+ electron polaron level.

The most remarkable results in this work were obtained through water exposure experiments, whereby surface cleaned CuFeO 2 and BiVO 4 substrates were exposed to water and analysed through photoelectron spectroscopy. Unexpectedly, the BiVO 4 surface was highly reduced after the contact with water was established, since features typical of V 4+ and Bi 0 could be detected.

Moreover, a decreased amount of lattice oxygen could be noticed. The reduction of the BiVO 4

surface could be partly reversed by introducing the sample into an oxygen atmosphere or completely reoxidised by subjecting the sample to an oxygen plasma treatment. For CuFeO 2 , the water exposure experiment demonstrated a similar behaviour as for the CuFeO 2 /ITO interface experiment. Namely, the E F -E VBM of CuFeO 2 only increased by a limited amount, which does not allow for water reduction, and Fe 3+ was partly reduced to Fe 2+ . Cu + was neither reduced nor oxidised. Furthermore, the UPS measurements, recorded after the CuFeO 2 pellets had been exposed

to water, showed a UPS valence band structure typical for physisorbed water and a relatively low work function of 4.4 eV. Therefore, the CuFeO 2 water exposure experiment can possibly be seen as an interface experiment. Indeed, physisorption of water may be feasible with the porous CuFeO 2 samples.

Perspectives

This dissertation has considered a few fundamental aspects regarding photochemical water splitting and, thus, contains a broad range of information, which may be used to guide future research.

First of all, it has been shown that the photodeposition technique is a versatile technique allowing the deposition of metals and metal(hydr)oxides on different kinds of photoabsorbers.

Additionally, deposits may exhibit a certain regioselectivity for powders wherein the particles have defined anisotropic shapes. The photodeposition method certainly has to be studied further to increase understanding why certain deposits show a particular regioselectivity, as well as to develop new anisotropic heterostructures. Third, the remarkable water exposure experiment on BiVO 4 suggests two things: 1) that the water reacts at the surface to form a volatile oxygen containing compound, leaving the surface reduced and 2) that the surface reduction could be reversed through oxidative treatments. These astonishing observations may indicate that the photochemical water oxidation mechanism at the BiVO 4 surface, follows a completely different path than what common knowledge dictates.

Namely, that in photocatalysis first light is absorbed, then charge carrier separation takes place and finally the surface redox reactions happen with the surface accumulated electrons and holes.

Instead, it may very well be that the surface reactions happen first after which the charge carriers at the surface need to be replenished through the absorption of light. This mechanism would also strongly agree with our assumptions on the role of a contact material in heterostructured BiVO 4 .

Namely, that a contact material primary helps to increase charge carrier separation and has little to no additional effect on the water oxidation reaction kinetics. We think that the research community should investigate these remarkable findings further in more detail. One should also check that the high energy holes generated during photoelectron spectroscopy measurements do not initiate the oxidation of the hydroxy groups at the water exposed BiVO 4 sample.

Furthermore, new photoabsorbers able to perform water oxidation and/or water reduction have to be identified, because BiVO 4 is a suboptimal light absorber and CuFeO 2 seems to have inherent limits, which inhibit its water reduction capacity. Off course, the GaInP/GaInAs combination has shown to have the ideal light absorption complementarity and demonstrated the state of the art 16.2% photoelectrochemical STH efficiency. However, their high cost and the insecurity about their long term stability in an electrolyte likely prevent III-V semiconductors in becoming the solution for particulate water splitting systems. Instead, new combinations of metal oxides have to be searched for, potentially partly mixing in nitrogen or sulphur to increase the valence band position, since now only binary, ternary and occasionally quaternary materials have been considered along the chemical space. Herein, computational studies could play a primary role. Preferably, iron should be excluded in the search towards photochemical water reducing compounds, due to the seemingly universal Fe 3+ /Fe 2+ polaron formation level which limits the potential photovoltage of iron containing compounds.

Although, strong doubts exist regarding the ability of photoelectrochemical water splitting with thin film stacks to be a viable alternative to steam methane reforming or photovoltaicelectrolysis in the near future, research should still be performed on these systems, because the information obtained may be transferred to the inherently cheaper but more difficult to study particulate systems. A particulate system which can achieve a 5% STH efficiency and has a stability of at least five years may be economically competitive to photovoltaic-electrolysis and even to steam methane reforming All in all, we can conclude that there is still a long road ahead to establish an efficient and stable particulate photochemical water splitting system that would yield economical parity to conventional hydrogen production techniques. La faible régiosélectivité du dépôt NiO x OH y comparativement aux autres matériaux de contact étudiés est probablement à l'origine des efficacités de l'oxydation photochimique peu élevées avec du NiO x OH y comme matériau de contact.

xix 2 .

 2 25 (a) Mechanism of selective deposition of reduction and oxidation cocatalysts on the {010} and {110} BiVO 4 facets, respectively, due to anisotropic photoinduced charge carrier migration and (b) FESEM images of a series of anisotropic BiVO 4 heterostructures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.26 Delafossite crystal structure for its two polytypes: (a) 3R and (b) 2H. . . . . . . . 2.27 Design of a PES instrument. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.28 Universal curvature of electron mean free path with respect to electron kinetic energy. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.29 V2p (and O1s) core level spectrum of a mixed vanadium oxide sample. . . . . . 2.30 Fe2p core level spectrum from the fractured surface of a Fe 0.96 O standard sample. 2.31 Fe2p core level spectrum from the fractured surface of a Fe 2 O 3 standard sample. 2.32 Co2p spectra for (bottom to top) Co metal, CoO, Co(OH) 2 and Co 3 O 4 . . . . . . . 2.33 Ni2p spectra for NiO, Ni(OH) 2 and NiOOH. . . . . . . . . . . . . . . . . . . . . . 2.34 Cu2p core level spectra for different copper species. . . . . . . . . . . . . . . . . 2.35 CuLMM core level spectra for different copper species. . . . . . . . . . . . . . . 3 Experimental Methods and Procedures 3.1 Set up for the truncated bipyramidal BiVO 4 powder synthesis. . . . . . . . . . . 3.2 Set up for the hexagonal nanoplatelet 2H-CuFeO 2 powder synthesis. . . . . . . 3.3 Photodeposition set up. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4 Spectrum solar simulator. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5 Magnetron sputtering set up. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.6 Schematic of DAISY-MAT. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7 (a) Magnetron cosputtering set up and (b) PLD set up for BiVO 4 thin film deposition.

3. 8

 8 Interaction of a primary electron beam with a solid state sample. . . . . . . . . .3.9Example of a survey spectrum measured on a cleaned BiVO 4 thin film surface. .

3. 10 2 , 1 4 Anisotropic BiVO 4 Heterostructures 4 . 1 6 Heterostructured CuFeO 2 6. 1 3 IUPAC:

 10214441213 Example of a high resolution Bi4f core level spectrum for which a Shirley and Tougaard function have been used to remove the background. . . . . . . . . . . 3.11 (a) Example of UP spectrum of BiVO 4 , (b) valence band region UPS and (c) valence band region XPS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xx LIST OF FIGURES 3.12 (a) Representation of an interface experiment, (b) Evolution of the binding energy with film thickness for the core levels, VBM and CBM of the substrate and the deposited contact film, (c) Evolution of the VBM of the substrate and the film with respect to the Fermi level position as a function of the film thickness. . . . 3.13 Schematic of DAISY-FUN. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.14 Experimental set up used to determine the sacrificial water splitting performance of the CuFeO 2 and BiVO 4 powders. . . . . . . . . . . . . . . . . . . . . . . . . . 4 Anisotropic BiVO 4 Heterostructures 4.1 (a) Powder X-ray diffractogram and (b) secondary electron image of monoclinic BiVO 4 synthesized according to the baseline parameters. . . . . . . . . . . . . . 4.2 Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized at different pH. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3 FESEM images of BiVO 4 prepared at pH 1.5 and pH 9. . . . . . . . . . . . . . . . 4.4 Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized at different autoclave temperatures. . . . . . . . . . . . . . . . . . . . . . . 4.5 Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized with different autoclave reaction times. . . . . . . . . . . . . . . . . . . . . 4.6 TEM image of BiVO 4 prepared with an autoclave reaction time of 64 h. . . . . . 4.7 (a) Secondary electron images and (b) powder X-ray diffractograms for BiVO 4 synthesized without SDS at different pH. . . . . . . . . . . . . . . . . . . . . . . 4.8 Secondary electron images and powder X-ray diffractograms for 5 BiVO 4 syntheses at pH 1.5. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.9 Secondary electron images and powder X-ray diffractograms for 2x upscaled BiVO 4 syntheses with different modifications to the baseline parameters. . . . . 4.10 Rietveld refinement on truncated bipyramidal BiVO 4 , irregular BiVO 4 and commercial BiVO 4 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.11 Secondary electron image of commercial BiVO 4 . . . . . . . . . . . . . . . . . . . 4.12 (a) UV-vis diffuse reflectance spectra and (b) Tauc plots for BiVO 4 with different morphologies. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.13 Survey X-ray photoelectron (XP) spectra for BiVO 4 with truncated bipyramidal, irregular, commercial and dendritic morphology. . . . . . . . . . . . . . . . . . . 4.14 High resolution Bi4f, V2p 3/2 and O1s core level XP spectra for BiVO 4 with truncated bipyramidal, irregular, commercial and dendritic morphology. . . . . . . . 4.15 Rhd B degradation profile for dendritic BiVO 4 . . . . . . . . . . . . . . . . . . . . 6.7 FESEM images of the batch of CuFeO 2 powder used for the photodeposition experiments. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.8 (a) FESEM image showing sides of hexagonal CuFeO 2 nanoplatelets, (b) common low-index facets of particles in the hexagonal crystal class and (c) possible indexing of exposed sides of the hexagonal CuFeO 2 nanoplatelets. . . . . . . . . . . . 6.9 Cu2p, Fe2p and O1s core level XP spectra and CuLMM Auger spectrum of the batch of CuFeO 2 powder used for the photodeposition experiments. . . . . . . . 6.10 FESEM images of CuFeO 2 /Ag (10 wt%). . . . . . . . . . . . . . . . . . . . . . . . 6.11 Left: FESEM images of CuFeO 2 /Ag (10 wt%) sample. Right: Ag M 4 VV Auger electron spectra. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.12 Cu2p, Fe2p, O1s and Ag3d core level XP spectra and CuLMM Auger spectrum of the CuFeO 2 /Ag (10 wt%) sample. . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.13 (a-c) FESEM and (d) BSE images of CuFeO 2 /Pt (10 wt%)-pH 2.9. . . . . . . . . . 6.14 Left: FESEM images of CuFeO 2 /Pt (10 wt%)-pH 2.9 sample. Right: Pt M 5 VV Auger electron spectra. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.15 Cu2p, Fe2p, O1s and Pt4f core level XP spectra and CuLMM Auger spectrum of the CuFeO 2 /Pt (10 wt%)-pH 2.9 sample. . . . . . . . . . . . . . . . . . . . . . . 6.16 FESEM images of CuFeO 2 /NiO x OH y samples. . . . . . . . . . . . . . . . . . . . . 6.17 Cu2p, Ni2p, Fe2p and O1s core level XP spectra and CuLMM Auger spectrum of the CuFeO 2 /NiO x OH y (± 10 wt%) sample. . . . . . . . . . . . . . . . . . . . . . . 6.18 (a) UV-Vis-NIR diffuse reflectance spectra and (b) Tauc plots for bare and heterostructured CuFeO 2 powder samples. . . . . . . . . . . . . . . . . . . . . . . . 6.19 Powder X-ray diffractograms for CuFeO 2 powder before and after calcination at 400 ℃ in air for 2 h. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.20 SEM images on CuFeO 2 powder (a) before and (b) after calcination at 400 ℃ in air for 2 h. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.21 Survey XP spectra of CuFeO 2 pellets subjected to different surface cleaning procedures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.22 Cu2p 3/2 , Fe2p, O1s core level and valence band XP spectra and CuLMM Auger spectrum of CuFeO 2 pellets subjected to different surface cleaning procedures. . 6.23 High resolution powder X-ray diffractograms for the CuFeO 2 powder used to prepare the pellets and a pellet after the thermally cleaning step at 200 ℃. . . . . 6.24 Cu2p, Fe2p, O1s, In3d 5/2 core level and VB XP spectra, and CuLMM Auger spectra for CuFeO 2 /ITO(RT) interface. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.25 UP spectra of the CuFeO 2 substrate and the ITO film. . . . . . . . . . . . . . . . CuLMM and Fe2p 3/2 difference spectra between bare CuFeO 2 and CuFeO 2 /ITO (4 s). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.27 Roughness profile of PLD BiVO 4 thin film and 2H-CuFeO 2 pellet. . . . . . . . . 6.28 (a) Energy band diagram of CuFeO 2 /ITO, (b) diagram showing the Fermi level tunability in CuFeO 2 and (c) a series of reduction potentials calculated at a certain pH on the RHE scale. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.29 Survey XP spectra of two thermally cleaned CuFeO 2 pellets before and after water exposure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.30 Cu2p 3/2 , Fe2p, O1s core level and valence band XP spectra and CuLMM Auger spectrum of two thermally cleaned CuFeO 2 pellets before and after water exposure. 6.31 UP spectra of two thermally cleaned CuFeO 2 pellets before and after water exposure. 6.32 Cu2p 3/2 , Fe2p 3/2 , CuLMM and Fe3p difference spectra between thermally cleaned CuFeO 2 pellet 1 before and after water exposure. . . . . . . . . . . . . . . . . . . A Regression BiVO 4 water oxidation A.1 Raw data obtained from the linear regression of the oxygen evolution curves and calculation of oxygen evolution for the sacrificial water oxidation experiments on the (heterostructured) BiVO 4 powders. . . . . . . . . . . . . . . . . . . . . . B Sommaire en français B.1 Morphologie des poudres BiVO 4 et CuFeO 2 . . . . . . . . . . . . . . . . . . . . . B.2 Poudres de BiVO 4 hétérostructurées. . . . . . . . . . . . . . . . . . . . . . . . . B.3 Poudres de CuFeO 2 hétérostructurées. . . . . . . . . . . . . . . . . . . . . . . . . B.4 Propriétés photocatalytiques des poudres de BiVO 4 . . . . . . . . . . . . . . . . . B.5 Alignements de bandes et accordabilité des niveaux de Fermi déterminés par des expériences d'interface avec des films minces BiVO 4 . . . . . . . . . . . . . . . . B.6 Alignement des bandes et accordabilité de niveau de Fermi déterminés par des expériences d'interface avec des granulés de CuFeO 2 . . . . . . . . . . . . . . . . B.7 Expériences d'exposition à l'eau sur film mince de BiVO 4 suivies par mesures XPS. B.8 Expériences d'exposition à l'eau sur des granulés de CuFeO 2 suivies par mesures XPS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Spectral fitting parameters for the V2p spectrum of various vanadium oxide species. 2.2 Cu2p 3/2 binding energy and modified Auger parameter literature values for copper species. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3 Ag3d binding energies for Ag, Ag 2 O and AgO. . . . . . . . . . . . . . . . . . . . 2.4 Pt4f binding energies for Pt, PtO and Pt(OH) 2 . . . . . . . . . . . . . . . . . . . . 2.5 Bi4f 7/2 binding energies for Bi(0) and Bi 2 O 3 . . . . . . . . . . . . . . . . . . . . . . 3 Experimental Methods and Procedures 3.1 Magnetron sputtering deposition parameters for RuO 2 , NiO, CoO x and ITO thin films. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2 Fitting parameters O1s and V2p 3/2 components. . . . . . . . . . . . . . . . . . . Reaction parameters and literature references for monoclinic truncated bipyramidal BiVO 4 powder synthesis procedures. . . . . . . . . . . . . . . . . . . . . . 4.2 Experimental details for the BiVO 4 based heterostructured powders fabricated through photodeposition. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 Band Alignment of Heterostructured BiVO 4 5.1 XPS parameters for a series of PLD BiVO 4 thin films. . . . . . . . . . . . . . . . 5.2 Summary of energetic values of all BiVO 4 based interface experiments. . . . . . xxviii LIST OF TABLES 5.3 Summary of BiVO 4 water exposure experiment. . . . . . . . . . . . . . . . . . . Experimental details for the CuFeO 2 based heterostructured powders fabricated through photodeposition. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Auger electron spectroscopy AFM: Atomic force microscopy ALD: Atomic layer deposition BE: Binding energy BoS: Balance of system BSE: Backscattered electrons CB: Conduction band CBM: Conduction band minimum CCD: Charge-coupled device CCS: Carbon capture and sequestration CoPi: Cobalt phosphate DAISY-MAT: Darmstadt integrated system for materials analysis DFT: Density functional theory DWE: Dual-working-electrode EDS: Energy-dispersive X-ray spectroscopy EDTA: Ethylenediaminetetraacetic acid EELS: Energy electron loss spectrum FEG: Field emission gun FESEM: Field emission secondary electron microscope FTO: Fluorine-doped tin oxide GC: Gas chromatograph GSPV-E: Grid supplemented photovoltaicelectrolysis HAADF: High-angle annular dark-field HER: Hydrogen evolution reaction ICSD: Inorganic crystal structure database IREC: Institut de recerca en energia de catalunya ITO: Tin-doped In 2 O International union of pure and applied chemistry JCPDS: Joint committee on powder diffraction standards KBi: Potassium borate LCAO: Linear combination of atomic orbitals LCH: Levelised cost of hydrogen production LCOE: Levelised cost of electricity MeOH: Methanol NHE: Normal hydrogen electrode NIST: National institute of standards and technology OER: Oxygen evolution reaction PEC: Photoelectrochemical PEM: Polymer electrolyte membrane PES: Photoelectron spectroscopy PLD: Pulsed laser deposition PV: Photovoltaic PV-E: Photovoltaic-electrolysis PXRD: Powder X-ray diffraction Rhd B: Rhodamine B RHE: Reversible hydrogen electrode SAED: Selected area electron diffraction SAM: Scanning Auger microscopy SDS: Sodium dodecyl sulphate SEC: Secondary electron cut-off xxx ACKNOWLEDGEMENTS SEI: Secondary electron imaging SEM: Secondary electron microscope SHE: Standard hydrogen electrode SMR: Steam methane reforming SPS: Surface photovoltage spectroscopy STEM: Scanning transmission electron microscope STH: Solar to hydrogen conversion efficiency TAS: Transient absorption spectroscopy TCD: Thermal conductivity detector TEM: Transmission electron microscope THF: Tetrahydrofuran TRMC: Time-resolved microwave conductivity UHV: Ultra-high vacuum UP: Ultraviolet photoelectron UPS: Ultraviolet photoelectron spectroscopy UV-Vis: UV visible UV-Vis-NIR: UV-visible-near infrared VB: Valence band VBM: Valence band maximum vol%: Volume percentage WF: Work function wt%: Weight percentage XP: X-ray photoelectron XPS: X-ray photoelectron spectroscopy ZPC: Zero point charge

10

 10

Figure 1 . 4 :

 14 Figure 1.4: Contour plot of LCH ($ per kg H 2 ) for the PV-E case (a) as a function of the plant efficiency and the areal capital expense of PV modules and electrolysers normalised by the necessary solar collection area ($ per square meter) and (b) as a function of the electrolyser and PV capital expenses ($ per watt). The yellow circle indicates the base case. 10

BiVO 4 .

 4 Next, the results of the interface experiments of BiVO 4 with the high work function (WF) materials, NiO, CoO x and RuO 2 , and the low work function material tin-doped In 2 O 3 (ITO) are described. Subsequently, the results of the interface experiments are thoroughly discussed to evaluate the interfacial properties of the analysed heterostructures, bulk Fermi level pinning phenomena and solid state band alignment.

Figure 2 .Figure 2 . 1 :

 221 Figure 2.1: Scheme showing how the 3sp 3 hybrid atomic orbitals of silicon combine into bands separated by a band gap.

Figure 2 . 2 :Figure 2 . 3 :

 2223 Figure 2.2: (a) n-type semiconductor, (b) intrinsic semiconductor, (c) p-type semiconductor. Left column: density of states D(E); Middle column: Fermi distribution f(E); Right column: carrier concentration (n,p). n and p are the concentrations of electrons and holes in the conduction and valence bands. E C : conduction band minimum, E V : valence band maximum, E F : Fermi energy level, E D : donor ionization energy, E A : acceptor ionization energy. 8

  4): 1) photons with an energy larger than the band gap are absorbed, leading to the creation of electron-hole pairs, 2) the electron-hole pairs are separated from each other, followed by charge carrier flow towards the semiconductor surface, 3) adsorbed reactants participate in redox reactions and consume the surface accumulated charge carriers.However, several other processes compete with photocatalysis (Figure2.5). As explained in the previous Section, the light induced electron-hole pairs can recombine again, whereby a distinction is made between surface and bulk recombination. Bulk recombination in indirect semiconductors mainly take place because of improper charge carrier separation and limited e-

Figure 2 . 4 :

 24 Figure 2.4: Photocatalysis scheme showing how light can be used by a semiconductor to initiate redox reactions at the semiconductor surface. e -: electron, h + : hole, VB: valence band, CB: conduction band.

Figure 2 . 5 :

 25 Figure 2.5: Common problems that may limit the efficiency of a photocatalytic material and solutions to overcome those problems.

Figure 2 . 9 :

 29 Figure 2.9: Maximum theoretical STH efficiency for a dual absorber in (a) a tandem configuration and (b) a side-by-side configuration as a function of both photoabsorber band gaps.For the tandem configuration the top photoabsorber is illuminated by the whole solar spectrum, whereas the bottom photoabsorber only receives the photons transmitted by the top photoabsorber. For the side-by-side configuration both photoabsorbers are illuminated by the complete solar spectrum.19

Figure 2 .

 2 13 a band diagram for the semiconductor/metal junction is shown. In this band diagram, the indicated semiconductor electron affinity, work function and ionization potential are defined as the energy difference between the vacuum level on the one side and the conduction band, Fermi level and valence band, respectively, on the other side. The work function, herein, represents the minimum energy needed to bring the electron from the solid to a point in vacuum near the solid surface. Due to the metal having no band gap, the electron affinity, ionization potential and work function are identical for a metal. When the metal and semiconductor are brought into contact, the Fermi level is used as a reference instead, due to the equilibration of the electron electrochemical potential in both materials. The semiconductor/metal band alignment can be estimated through the Mott-Schottky rule, which state that the energy transfer between the metal and the semicon-

Figure 2 . 13 :

 213 Figure 2.13: Scheme illustrating the band alignment of an n-type semiconductor with a metal before and after contact, and with/without Fermi level pinning at surface states. I P,SC : ionisation potential of semiconductor, χ SC : electron affinity of semiconductor, φ SC : work function of semiconductor, φ M : work function of metal, Φ B,n : barrier height between Fermi level and CB, qV BB : band bending.

Figure 2 . 14 :Figure 2 . 15 :

 214215 Figure 2.14: Scheme showing charge carrier depletion, accumulation and inversion for an n-type semiconductor. The charges depicted in the scheme are the charges that are left behind near the semiconductor surface after the mobile charge carrier transfer.

Figure 2 . 18 :

 218 Figure 2.18: Band positions of semiconductor with respect to the absolute vacuum scale, SHE and RHE at flat band potential for pH ZPC and pH 0.

Figure 2 . 19 :

 219 Figure 2.19: Correlation of flatband positions with electron affinity calculation.57

Figure 2 .

 2 Figure 2.20: (a) Reaction pathway of HER at 0 V potential and (b) activity volcano plot of HER for metals and MoS 2 at the equilibrium potential of 0 V. 82,87

  is predicted at a ∆G O-OH of 1.6 ± 0.1 eV ((3.2 ± 0.2 eV)/2 = 1.6 ± 0.1 eV). Indeed, this assumption is confirmed by the volcano plot in Figure 2.21b.

Figure 2 .

 2 Figure 2.21: (a) Reaction pathway of associative OER mechanism at 0 V potential and (b) activity volcano plot of OER according to ∆G O -∆G OH for a series of oxides. 74

2. 23 )

 23 consists of VO 4 tetrahedra and BiO 8 dodecahedra. The BiO 8 dodecahedra are connected to each other and form arrays along the [1 10] direction, whereas the VO 4 tetrahedra are isolated from each other. The BiO 8 dodecahedra connect to the VO 4 tetrahedra by sharing an apex oxygen atom. Moreover, monoclinic clinobisvanite exhibits the characteristics of a layered structure, with the Bi and V atoms being alternated along the crystalline axes, having the -Bi-V-configuration along the c and a axis. This particular layered structure, in addition to the shorter Bi-O bond lengths along the b axis, 113 makes that the (010) plane is the principal cleavage plane.

Figure 2 .

 2 Figure 2.23: (a) Ball-and-stick and (b) polyhedron representation of the monoclinic clinobisvanite BiVO 4 crystal structure (with BiO 8 dodecahedron in purple and VO 4 tetrahedron in blue). (c) top and (d) side view of crystal structure. 113

Figure 2 . 24 :

 224 Figure 2.24: Scheme illustrating how Rh,La-codoped SrTiO 3 and Mo:BiVO 4 were immobilised ona conductive carbon film.37

Figure 2 . 2 . 3 . 2 2 CuFeO 2 Figure 2 . 26 :

 223222226 Figure 2.25: (a) Scheme showing the selective deposition of reduction and oxidation cocatalysts on the {010} and {110} BiVO 4 facets, respectively, due to anisotropic photoinduced charge carrier migration and (b) FESEM images of a series of anisotropic BiVO 4 heterostructures. Scale bar: 500 nm. 143,144

  photons are usually generated through X-ray anodes, whereby monochromatized Al Kα or Mg Kα X-rays with photon energies of 1486.6 eV and 1253.6 eV, and energy widths of 0.85 eV and 0.70 eV, respectively, are the most popular. The photon source in UPS is mostly a helium discharge lamp, which can emit photons with two characteristic energies 21.2 eV and 40.8 eV, denoted as He I and He II photons, respectively. Furthermore, PES systems can be coupled to a synchrotron, whereby the advantages of a synchrotron include the tunability of the photon energy, high brilliance, strong collimation, and high photon flux of synchrotron irradiation.

Figure 2 . 27 :

 227 Figure 2.27: Schematic of a PES instrument.©Phi Inc.

Figure 2 . 28 :

 228 Figure 2.28: Universal curvature of electron mean free path with respect to kinetic energy. 165

Figure 2 .

 2 Figure 2.29: V2p (and O1s) core level spectrum of a mixed vanadium oxide sample. 166

Figure 2 . 30 :

 230 Figure 2.30: Fe2p core level spectrum from the fractured surface of a Fe 0.96 O standard sample. 173

Figure 2 . 31 :

 231 Figure 2.31: Fe2p core level spectrum from the fractured surface of a Fe 2 O 3 standard sample. 173

Figure 2 . 33 .

 233 Generally, determining the composition of nickel hydr(oxide) mixtures through peak fitting generates large absolute errors due to the shape of the main Ni2p 3/2 and Ni2p 1/2 lines and their satellite structure being very similar for the different nickel hydr(oxide) species. However, NiO can be distinguished from Ni(OH) 2 and γ -NiOOH due to the main Ni2p 3/2 and Ni2p 1/2 lines being split up into two distinct signals separated by about 1 eV.

Figure 2 . 32 :

 232 Figure 2.32: Co2p spectra for (bottom to top) Co metal, CoO, Co(OH) 2 and Co 3 O 4 . 171

Figure 2 . 33 :

 233 Figure 2.33: Ni2p spectra for NiO, Ni(OH) 2 and NiOOH. 176

Figure 2 . 34 :

 234 Figure 2.34: Cu2p core level spectra for different copper species. 166

Figure 2 . 35 :

 235 Figure 2.35: CuLMM core level spectra for different copper species. 178

Figure 3 .

 3 1 summarizes the BiVO 4 synthesis.

Figure 3 . 1 :

 31 Figure 3.1: BiVO 4 synthesis: left: foaming aqueous mixture of bismuth precursor, vanadium precursor and sodium dodecyl sulphate (SDS); middle: treatment in pressure reactor; right: yellow BiVO 4 precipitate after removal from pressure reactor.

Figure 3 . 2 :

 32 Figure 3.2: Teflon beaker and pressure reactor used for the CuFeO 2 nanoplatelet synthesis.

Figure 3 . 3 :Figure 3 . 4 :

 3334 Figure 3.3: Photodeposition set up.

Figure 3 . 5 :

 35 Figure 3.5: Schematic of the magnetron sputtering set up. The grey and blue spheres represent the sputtered target atoms and the sputter gas atoms, representatively. 3

Figure 3 . 6 :

 36 Figure 3.6: Schematic of DArmstadt Integrated SYstem for MATerials science (DAISY-MAT).Adapted from[4].

Figure 3 . 7 :

 37 Figure 3.7: Diagrams of (a) the magnetron cosputtering set up 5 and (b) the PLD set up for the BiVO 4 thin film deposition. 6

3. 4 . 3 microscopy 3 . 4 . 3 . 1

 433431 Electron Scanning electron microscopy (SEM)Scanning electron microscopy (SEM) is a technique that allows obtaining chemical and morphological information of a surface by scanning a focused electron beam over the sample surface. Commonly, the electron beam can be generated through thermionic heating of a solid state crystal (mainly CeB 6 or LaB 6 ), resistive heating of a tungsten filament or field emission from a tungsten tip. The electrons released in this fashion are then accelerated by a voltage between 1 -40 kV and focused through an arrangement of condenser lenses from the source onto the sample surface. Electron beams generated by a field emission gun are narrower and brighter compared to thermionic devices and, thus, allow gaining a higher signal-to-noise ratio and an increased spatial resolution up to 0.5 nm. The electrons in the primary electron beam reach the sample surface with high kinetic energy, allowing them to interact with the sample up to a depth of 1 µm. Essentially, the electrons interact with the atoms in the sample through elastic and inelastic scattering processes, generating various depth-dependent signals including X-rays, back-scattered electrons and secondary electrons (Figure3.8).

Figure 3 . 9 :

 39 Figure 3.9: Example of a survey spectrum measured on a cleaned BiVO 4 thin film surface, whereby the detected BiVO 4 related core levels are indicated.

  B are the intensities of a particular core level corresponding to element A and B, respectively, and S CL A and S CL B are the atomic sensitivity factors of a particular core level corresponding to element A and B, respectively. The atomic sensitivity factors depend on the different photoelectron generation cross sections of different core level electrons as well as on the instrumental parameters, which normally do not vary for different core levels.

Figure 3 . 10 :

 310 Figure 3.10: Example of a high resolution Bi4f core level spectrum for which a Shirley and Tougaard function have been used to remove the background.

Figure 3 .

 3 Figure 3.11: (a) Example of UP spectrum of BiVO 4 , (b) valence band region UPS and (c) valence band region XPS. E SEC : secondary electron cut-off; E UPS VBM : valence band maximum from UPS valence band region; E XPS VBM : valence band maximum from XPS valence band region.

Figure 3 . 13 :

 313 Figure 3.13: Schematic of DArmstadt Integrated SYstem for FUNdamental research (DAISY-FUN). Adapted from [18].

  Figure 3.14: Experimental set up to determine the sacrificial water splitting performance of the CuFeO 2 and BiVO 4 (heterostructured) powders. (a) Effluent (b) vacuum/Ar (c) thermostat outlet (d) thermostat inlet (e) thermostatically controlled reaction vessel (f) thermostat inlet and out let (g) pressure sensor.

Figure 4 . 1 :

 41 Figure 4.1: (a) Powder X-ray diffractogram and (b) secondary electron image of monoclinic BiVO 4with irregular morphology synthesized according to the baseline parameters (pH 7, SDS, T autoclave : 180 ℃ and t autoclave : 24 h).

Figure 4 . 2 :Figure 4 . 3 :

 4243 Figure 4.2: Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized at different pH.

Figure 4 . 4 :

 44 Figure 4.4: Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized at different autoclave temperatures.

Figure 4 . 5 :

 45 Figure 4.5: Secondary electron images and powder X-ray diffractograms for BiVO 4 synthesized with different autoclave reaction times.

Figure 4 . 6 : 19 10

 4619 Figure 4.6: TEM image of BiVO 4 prepared with an autoclave reaction time of 64 h.

Figure 4 . 7 :

 47 Figure 4.7: (a) Secondary electron images and (b) powder X-ray diffractograms for BiVO 4 synthesized without SDS at different pH.

Figure 4 . 8 :

 48 Figure 4.8: Secondary electron images and powder X-ray diffractograms for 5 BiVO 4 syntheses at pH 1.5.

Figure 4 . 9 :

 49 Figure 4.9: Secondary electron images and powder X-ray diffractograms for 2x upscaled BiVO 4 syntheses with different modifications to the baseline parameters.

Figure 4 . 10 :Figure 4 . 11 :

 410411 Figure 4.10: Rietveld refinement on truncated bipyramidal BiVO 4 , irregular BiVO 4 and commercial BiVO 4 with the measured diffractogram in black, the fit in red and the difference in blue. In the refinement monoclinic clinobisvanite (blue, JCPDS: 75-7867), tetragonal clinobisvanite (red, JCPDS: 78-1534) and orthorhombic Bi 2 VO 5.5 (green, JCPDS: 88-870) were considered.

Figure 4 .

 4 Figure 4.12: (a) UV-vis diffuse reflectance spectra and (b) Tauc plots for BiVO 4 with different morphologies.

Figure 4 . 13 :

 413 Figure 4.13: Survey X-ray photoelectron (XP) spectra for BiVO 4 with truncated bipyramidal, irregular, commercial and dendritic morphology.

Figure 4 . 14 :

 414 Figure 4.14: High resolution Bi4f, V2p 3/2 and O1s core level XP spectra for BiVO 4 with truncated bipyramidal, irregular, commercial and dendritic morphology.

Finally, the photocatalytic

  oxidation performance of the differently shaped BiVO 4 powders was tested by evaluating at which rate the in water dispersed powders could degrade a dissolved organic dye in the presence of light. The whole dye degradation procedure is explained in detail in Section 3.4.5.1.

Figure 4 .

 4 15 shows how the absorbance of rhodamine B, the dye used in this work, decreases over time during a dye degradation experiment. Due to the dye absorbance being proportional to the concentration of the unreacted dye, the dye degradation efficiency can be determined by registering how the absorbance of the peak maximum, here at 554 nm, changes over time(Figure 4.16). Without light or in the absence of BiVO 4 , no decrease in absorbance of the dye was noticed, seemingly confirming the photocatalytic nature of the dye degradation process. Remarkably, the irregular shaped BiVO 4 powder had the best rhodamine B degradation efficiency, because only 10% of the original absorbance was left after a degradation time of 120 min. Worse efficiencies were obtained with the other powders. Over the photodegradation time span of 120 min the drop in dye absorbance (at 554 nm) for dendritic, commercial and truncated bipyramidal BiVO 4 amounted on average to 57%, 49% and 46%, respectively.

Figure 4 . 15 :

 415 Figure 4.15: Rhd B degradation profile for dendritic BiVO 4 . The reaction mixture was stirred for 30 min in dark, so that 0 min is the point in time at which the light was switched on.

Figure 4 . 16 :

 416 Figure 4.16: Photocatalytic degradation of Rhd B for BiVO 4 with different morphologies. Left: Absorbance (at 554 nm) and right: natural logarithm of absorbance versus the degradation time with light switched on at 0 min. The mean first order rate constants k mean and mean R 2 linear fitting parameter are denoted in the graph.

Figure 4 . 17 :

 417 Figure 4.17: FESEM images of BiVO 4 /Ag (10 wt%).

Figure 4 .

 4 Figure 4.18: (a) FESEM image of BiVO 4 /Ag (10 wt%) and EDS elemental maps: (b) composite, (c) Bi, (d) V and (e) Ag.

Figure 4 . 20 :

 420 Figure 4.20: FESEM images of (a) BiVO 4 /Pt (10 wt%)-pH 2.5, (b) BiVO 4 /Pt (10 wt%)-pH 5.8 and (c) BiVO 4 /Pt (1 wt%)-pH 3.6.

4 /

 4 Pt samples are shown in Figure 4.22. Considering all Pt4f spectra, three different Pt4f doublets can be noticed. Peak fitting was not performed because the pseudo-Voigt profile of the platinum doublets was somewhat lost due to inhomogeneous sample charging, which was especially noticeable for the BiVO 4 /Pt(10 wt%)-pH 2.5 sample due to prominent Bi4f line broadening. For BiVO 4 /Pt(10 wt%)-pH 2.5 only one doublet can be observed, in which the two signals at 71.2 eV and 74.6 eV likely correspond to Pt4f 7/2 and Pt4f 5/2 , respectively, of metallic platinum. 26 For the BiVO 4 /Pt(10 wt%)-pH 5.8 three doublets can be noticed, in which the signals at 71.2 eV, 72.5 eV and 75.0 eV likely

Figure 4 .

 4 Figure 4.21: Left: FESEM image of BiVO 4 /Pt(10 wt%)-pH 5.8 sample with the SAM measurement zones indicated in yellow (on {010} facets) and black (on {110} facets). Right: SAM sputter depth profile showing the 1 st derivative of Pt N 7 VV Auger signal for the different measurement zones.

Figure 4 . 22 :

 422 Figure 4.22: Pt4f, Bi4f, V2p 3/2 , O1s core level XP spectra of BiVO 4 /Pt samples. In the Pt4f spectra the Pt4f 7/2 features of Pt(0), Pt(OH) 2 and PtO are denoted.

Figure 4 .

 4 Figure 4.23: FESEM image of BiVO 4 /CoO x OH y (2 wt%).

Figure 4 .

 4 Figure 4.24: FESEM image and SAM elemental maps of BiVO 4 /CoO x OH y (2 wt%).

Figure 4 . 25 :

 425 Figure 4.25: Co2p 3/2 Bi4f, V2p 3/2 , O1s core level XP spectra of BiVO 4 /CoO x OH y (2 wt%).

Figure 4 . 26 :

 426 Figure 4.26: FESEM images: (a) bare BiVO 4 , (b) BiVO 4 /NiO x OH y (± 10 wt%) (c) + Ni EDS, (d) BiVO 4 /NiO x OH y (1 wt%) and (f,g) BiVO 4 /NiO x OH y (0.1 wt%).

Figure 4 . 27 :

 427 Figure 4.27: Left: FESEM image of BiVO 4 /NiO x OH y (1 wt%) sample with the SAM measurement zones indicated in yellow (on {010} facets) and black (on {110} facets). Right: SAM sputter depth profile showing the 1 st derivative of Pt N 7 VV Auger signal for the different measurement zones.

Figure 4 . 28 :

 428 Figure 4.28: Left: FESEM image of BiVO 4 /NiO x OH y (0.1 wt%) sample with the SAM measurement zones indicated in yellow (on {010} facets) and black (on {110} facets). Right: SAM sputter depth profile showing the 1 st derivative of Pt N 7 VV Auger signal for the different measurement zones.

Figure 4 .

 4 Figure 4.29: (a) TEM bright-field and (b) STEM dark-field image of BiVO 4 /NiO x OH y (1 wt%). (c) EDS of BiVO 4 microparticle and of (d) NiO x OH y nanodeposit.

Figure 4 . 31 :

 431 Figure 4.31: FESEM images: photodeposition of Ru(acac) 3 on BiVO 4 with FeCl 3 as electron scavenger in (a) acetic acid and (b) THF.

Figure 4 .

 4 Figure 4.33: (a) UV-Vis diffuse reflectance spectra and (b) Tauc plots for heterostructured BiVO 4 powders.

Figure 4 .)Figure 4 . 34 :

 4434 Figure 4.34 shows the photocatalytic degradation of Rhd B for the various heterostructured BiVO 4 photocatalysts. Here, BiVO 4 /NiO x OH y and BiVO 4 /Ag seem to be the most efficient, as the original absorbance drops by about 50% after 2 h of photodegradation. In contrast, BiVO 4 /CoO x OH y appears to have a similar Rhd B photodegradation efficiency as bare BiVO 4 .ln (A t A 0) vs t plots as well as the first order rate constants, determined from these plots, are shown in Figure4.34b. The first order rate model seems to fit more or less, as the linear regression yield R 2 values in the range of 0.95 -0.99. Overall, all heterostructures have rate constants which are only slightly higher than the one of bare truncated bipyramidal BiVO 4 . This is in part due to the already high first order rate constant of bare BiVO 4 , which is here 3.3 x 10 -3 min -1 , whereas Li et al. found a first order rate constant of 1.7 x 10 -3 min -1 .2 The authors also found rate constants of an equivalent order of magnitude, as the ones we found, for heterostructured truncated bipyramidal BiVO 4 covered by a single cocatalyst. Additionally, they report that only for BiVO 4

  2 ml amounts to half of the millimolar gas volume according to the ideal gas law at ambient conditions and because NaIO 4 scavenges two electrons per molecule according to: IO -

Figure 4 . 35 :Figure 4 . 36 :

 435436 Figure 4.35: Sacrificial water oxidation for several heterostructured BiVO 4 photocatalysts: gas evolution versus time. Gas evolution rates (mL h -1 O 2 ) are depicted as well. 1 mmol NaIO 4 was used as sacrificial agent and the amount of photocatalyst powder was 50 mg.

( 16 )

 16 Merupo, V.; Velumani, S.; Oza, G.; Makowska-Janusik, M.; Kassiba, A. Structural, electronic and optical features of molybdenum-doped bismuth vanadium oxide. Mater. Sci. Semicond. Process. 2015, 31, 618-623. (17) Jiang, H.; Dai, H.; Meng, X.; Zhang, L.; Deng, J.; Liu, Y.; Au, C. T. Hydrothermal fabrication and visible-light-driven photocatalytic properties of bismuth vanadate with multiple morphologies and/or porous structures for methyl orange degradation. J. Environ. Sci. (China) 2012, 24, 449-57. (18) Meng, X.; Zhang, L.; Dai, H.; Zhao, Z.; Zhang, R.; Liu, Y. Surfactant-assisted hydrothermal fabrication and visible-light-driven photocatalytic degradation of methylene blue over multiple morphological BiVO 4 single-crystallites. Mater. Chem. Phys. 2011, 125, 59-65.

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 19 Scherrer, P. Determination of the Size and Internal Structure of Colloidal Particles using X-rays. Nachr. Ges. Wiss. Göttingen 1918, 2, 98.

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 20 Tokunaga, S.; Kato, H.; Kudo, A. Selective preparation of monoclinic and tetragonal BiVO 4 with scheelite structure and their photocatalytic properties. Chem. Mater. 2001, 13, 4624-4628. (21) Cooper, J. K.; Gul, S.; Toma, F. M.; Chen, L.; Glans, P. A.; Guo, J.; Ager, J. W.; Yano, J.; Sharp, I. D. Electronic structure of monoclinic BiVO 4 . Chem. Mater. 2014, 26, 5365-5373.

  As described in the previous Chapter, BiVO 4 based devices should be carefully structured to overcome the inherent slow reaction kinetics and inefficient charge carrier separation of BiVO 4 .

  However, Koffyberg et al. found that the Butler-Ginley equation only gave reliable valence band edge positions for transition metal oxides whose valence bands were mainly built up by O2p orbitals. When the valence band contained primarily contributions from the transition metal d-orbitals, valence band positions were calculated which did not correspond to experimentally determined valence band positions.

V2p 3 /

 3 2 and O1s in Figure5.2b.

Figure 5 . 1 :

 51 Figure 5.1: Raman spectra of a BiVO 4 gradient thin film and of a Mo:BiVO 4 single crystal. The Raman spectra of the gradient thin film were measured at specific distances from the Bi rich edge.

Figure 5 . 2 :Figure 5 . 3 :

 5253 Figure 5.2: Survey and high resolution core level XP spectra of a gradient BiVO 4 thin film at specific distances from the Bi rich edge before and after oxygen plasma.

Figure 5 . 4 :

 54 Figure 5.4: Bi4f, V2p 3/2 , O1s core level and valence band XP spectra of a gradient BiVO 4 thin film at specific distances from the Bi rich edge along the entire sample width (22 mm).

Figure 5

 5 Figure 5.5: V/Bi surface ratio of a gradient BiVO 4 thin film using the tabulated sensitivity factors (crosses) and after applying the correction factor determined from the (010) exposed Mo:BiVO 4 single crystal (squares). In the inset the V/Bi ratio of the (010) Mo:BiVO 4 single crystal is shown after various treatments: 0: no treatment; 1: polished; 2: polished + 15 min O 2 plasma; 3: 2nd time polished + 15 min O 2 plasma.

Figure 5 .

 5 6a the Bi4f, V2p 3/2 and valence band X-ray photoelectron spectra of the polished and sputter cleaned Mo:BiVO 4 single crystal are shown as well as those of a sputter cleaned gradient BiVO 4 thin film at 13 mm from the Bi rich edge. The Bi4f and V2p 3/2 spectra of the single crystal are shifted by about 0.2 eV towards higher binding energies and its E F -E VBM difference is measured at 2.2 eV, compared to the measured E F -E VBM of 1.8 eV of the gradient thin film. The shift towards higher binding energies could be due to a shift in the Fermi level position, originating from the Mo doping in the single crystal, or due to surface charging, because of a too low electrical conductivity of BiVO 4 . The effect of surface charging can be clearly seen in the UP spectra of the corresponding samples (Figure5.6b), where the E F -E VBM

Figure 5 . 6 :

 56 Figure 5.6: (a) V2p 3/2 , Bi4f core level and valence band XP spectra and (b) UP spectra of the gradient BiVO 4 thin film at 13 mm from the Bi rich edge (green) and the polished (010) surface of the Mo:BiVO 4 single crystal (black). Both samples were cleaned by sputtering with an oxygen plasma.

Figure 5

 5 Figure 5.7: V/Bi ratio of gradient BiVO 4 thin film: determined from the Bi4f and V2p 3/2 XP core level areas after using the tabulated sensitivity factors and applying the correction factor determined from the (010) exposed Mo:BiVO 4 single crystal (squares) and determined through EDS after applying the correction factor (1.08) determined from the EDS measurements of (010) exposed Mo:BiVO 4 single crystal (triangles).

Figure 5 . 8 :

 58 Figure 5.8: (a) Survey and (b) high resolution core level XP spectra of a PLD BiVO 4 thin film before and after oxygen plasma.

Figure 5 . 9 :

 59 Figure 5.9: Summary of photoelectron spectroscopy analysis on BiVO 4 thin film substrates by plotting the determined binding energies of O1s, V2p 3/2 , Bi4f 7/2 and E F -E VBM relative to one another.

Figure 5 . 10 :

 510 Figure 5.10: Difference spectra of O1s, V2p 3/2 , Bi4f 7/2 and valence band of a sample with a highO1s binding energy (gradient BiVO 4 after 15 min oxygen plasma (yellow) and a sample with a low O1s binding energy (PLD BiVO 4 after 2 h of heating at 400 ℃ in oxygen atmosphere (blue). The spectra were normalized and shifted to the same binding energy.

Figure 5 . 11 :

 511 Figure 5.11: O1s, V2p 3/2 , Ru3d, Bi4f core level and valence band XP spectra for the BiVO 4 /RuO 2 interface. RuO 2 deposition times are denoted in the O1s spectra.

Figure 5 . 13 :

 513 Figure 5.13: O1s, V2p 3/2 , In3d 3/2 , Bi4f core level and valence band XP spectra for BiVO 4 /ITO(400 ℃) interface. ITO deposition times are denoted in the O1s spectra.

Figure 5 . 14 :

 514 Figure 5.14: Ni2p, O1s, V2p 3/2 , Bi4f core level and valence band XP spectra for the BiVO 4 /NiO interface. NiO deposition times are denoted in the O1s spectra.

Figure 5 . 16 :

 516 Figure 5.16: Co2p 3/2 , O1s, V2p 3/2 , Bi4f core level and valence band XP spectra for the BiVO 4 /CoO x interface. CoO x deposition times are denoted in the O1s spectra.

19 and Figure 5 .

 5 20 include the vacuum energy levels of the materials, which are derived from work function measurements using UPS. According to the measured work functions, the vacuum energies of BiVO 4 and RuO 2 are aligned, while there is a dipole potential step of about 1 eV between BiVO 4 and ITO (RT and 400 ℃), and a dipole potential step of 1.3 eV for the BiVO 4 /CoO x and BiVO 4 /NiO. The dipole potential step for the

Figure 5 . 18 :

 518 Figure 5.18: Shift of the V2p 3/2 and Bi4f 7/2 binding energies with increasing contact layer thickness (arrows) for all the interface experiments. The dashed blue line represents a slope of 1.

Figure 5 . 19 :

 519 Figure 5.19: Energy band diagrams for the BiVO 4 /RuO 2 , BiVO 4 /NiO, BiVO 4 /CoO x interfaces. All energy values are denoted in eV. The band gaps of BiVO 4 , NiO, CoO x were taken from literature.[START_REF] Tokunaga | Selective preparation of monoclinic and tetragonal BiVO 4 with scheelite structure and their photocatalytic properties[END_REF][START_REF] Irwin | p-Type semiconducting nickel oxide as an efficiency-enhancing anode interfacial layer in polymer bulk-heterojunction solar cells[END_REF][START_REF] Xu | The advances of Co 3 O 4 as gas sensing materials: A review[END_REF]

Figure 5 . 20 :

 520 Figure 5.20: Energy band diagrams for the BiVO 4 /ITO(RT) and BiVO 4 /ITO(400 ℃) interfaces. All energy values are denoted in eV. The band gap of BiVO 4 and ITO were taken from literature. 55,65

  2 and CoO x similar Schottky-type barrier heights between 1.0 -1.1 eV are obtained, despite the different work functions of the contact materials. The similarity of these Schottky-type barrier heights strongly indicate that a maximum shift of E F is reached. The reason for this could be Fermi level pinning. Another possibility is that the work function of CoO x , NiO and RuO 2 is not high enough to obtain further band bending. During the deposition of RuO 2 ,

Figure 5 .

 5 Figure 5.23 summarizes the interface experiments by showing the range over which the BiVO 4

Figure 5 . 23 :

 523 Figure 5.23: Left: diagram showing the Fermi level tunability in BiVO 4 and the limits based on the obtained information from the interface experiments. The BiVO 4 band positions are here taken from literature.20,43,78 Right: a series of standard reduction potentials related to bismuth and vanadium species are shown on the standard hydrogen electrode scale.82,83

Figure 5 . 24 :

 524 Figure 5.24: Conduction band and valence band alignment of a series of semiconductor oxides relative to the (a) H + /H 2 standard redox potential and to the (b) RuO 2 Fermi level position. The barrier heights of these oxides at the interfaces with RuO 2 (blue) and ITO (yellow) were determined through interface experiments. a-TiO 2 : anatase, r-TiO 2 : rutile, Φ B,RuO 2 : barrier height between substrate and RuO 2 , ∆Φ B : barrier height difference between substrate/RuO 2 and substrate/ITO.

Figure 5 . 25 :

 525 Figure 5.25: Survey XP spectra for the BiVO 4 water exposure experiment. O 2 : oxygen flow (2 sccm, 1 x 10 -3 mbar), H 2 O: water exposure.

3 MOFigure 5 . 26 :

 3526 Figure 5.26: O1s, V2p 3/2 , Bi4f core level and valence band XP spectra for the BiVO 4 water exposure experiment. The treatments are denoted in the O1s spectra. O 2 : oxygen flow (2 sccm, 1 x 10 -3 mbar), H 2 O: water exposure.

Figure 5 .

 5 Figure 5.30, the survey and C1s XP spectra still exhibit a signal related to extrinsic carbonaceous species. Most likely, the thermal contact between the heat source and the sample holder was not well established. The Bi4f, O1s, V2p 3/2 core level and valence band XP spectra recorded during the water desorption experiment are shown in Figure 5.31. The binding energies of V2p 3/2 , Bi4f

Figure 5 . 27 :Figure 5 . 28 :

 527528 Figure 5.27: Fitting of the V2p 3/2 and O1s core level components from the water exposure experiment on BiVO 4 . The treatments are denoted in the O1s spectra. The ratios of the integrated components are denoted as well.

Figure 5 . 29 :Figure 5 .

 5295 Figure 5.29: UP spectra (He I excitation) for the BiVO 4 water exposure experiment. Secondary electron cut-offs (SEC) are denoted in the graph. O 2 : oxygen flow (2 sccm, 1 x 10 -3 mbar), H 2 O: water exposure.

Figure 5 . 31 :

 531 Figure 5.31: O1s, V2p 3/2 , Bi4f core level and valence band XP spectra for the BiVO 4 water desorption experiment. The treatments and water coverage (in red: RT; in blue: 77 K) are denoted in the O1s spectra. O 2 : oxygen flow (2 sccm, 1 x 10 -3 mbar).

OH+CO O 2 -

 2

Figure 5 . 32 :

 532 Figure 5.32: Fitting of the V2p 3/2 and O1s core level components from the water desorption experiment on BiVO 4 . The treatments are denoted in the O1s spectra. The treatments in red and blue refer to the core level spectra measured at RT and at 77 K, respectively. Ratios of integrated components are denoted as well.

Figure 6 . 1 :

 61 Figure 6.1: (a) Secondary electron images and (b) powder X-ray diffractograms for 4x and 8x upscaled CuFeO 2 powder with the (hkl) reflections of 2H-CuFeO 2 indicated in the graph.

Figure 6 . 2 :

 62 Figure 6.2: High resolution powder X-ray diffractogram of 4x upscaled CuFeO 2 powder, whereby the 2H-CuFeO 2 (red, ICSD: 01-079-1546) and 3R-CuFeO 2 (blue, ICSD: 01-075-2146) phases are indicated.

Figure 6 . 3 :

 63 Figure 6.3: FESEM images of 4x upscaled CuFeO 2 powder.

6. 4 Figure 6 . 4 :

 464 Figure 6.4: SEM images of (a) almost not aggregated, (b) slightly aggregated and (c) strongly aggregated CuFeO 2 powder.

Figure 6 . 5 :

 65 Figure 6.5: Powder X-ray diffractograms of CuFeO 2 powders with almost no (blue), a very little amount of (green) and a little amount (yellow) of 3R-CuFeO 2 .

80 Figure 6 . 6 :

 8066 Figure 6.6: High resolution powder X-ray diffractogram of the batch of CuFeO 2 powder used for the photodeposition experiments, whereby the 2H-CuFeO 2 (red, ICSD: 01-079-1546) and 3R-CuFeO 2 (blue, ICSD: 01-075-2146) phases are indicated.

Figure 6 . 8 :

 68 Figure 6.8: (a) FESEM image showing sides of hexagonal CuFeO 2 nanoplatelets, (b) common lowindex facets of particles in the hexagonal crystal class and (c) possible indexing of exposed sides of the hexagonal CuFeO 2 nanoplatelets.

Figure 6 . 10 :

 610 Figure 6.10: FESEM images of CuFeO 2 /Ag (10 wt%).

Figure 6 .

 6 Figure 6.11: Left: FESEM images of CuFeO 2 /Ag (10 wt%) sample with the AES measurement points indicated in the top image. Right: Ag M 4 VV Auger electron spectra for the different measurement points.

Figure 6 .

 6 Figure 6.13: (a-c) FESEM and (d) BSE images of CuFeO 2 /Pt (10 wt%)-pH 2.9.

Figure 6 .

 6 Figure 6.14: Left: FESEM images of CuFeO 2 /Pt (10 wt%)-pH 2.9 sample with the AES measurement points indicated in the top image. Right: Pt M 5 VV Auger electron spectra for the different measurement points.

Figure 6 . 16 :

 616 Figure 6.16: FESEM images of CuFeO 2 /NiO x OH y samples with (a) ± 10 wt% and (b) 1 wt% NiO x OH y .

Figure 6 .

 6 Figure 6.18: (a) UV-Vis-NIR diffuse reflectance spectra and (b) Tauc plots for bare and heterostructured CuFeO 2 powder samples.

Figure 6 . 19 :Figure 6 . 20 :

 619620 Figure 6.19: Powder X-ray diffractograms for CuFeO 2 powder before (blue) and after calcination at 400 ℃ in air for 2 h (yellow). The crystalline phases 2H-CuFeO 2 (blue, JCPDS: 96-202-0231), Fe 3 O 4 (red, JCPDS: 96-900-6319) and CuO (light green, JCPDS: 96-101-1149) are indicated.

Figure 6 .

 6 Figure 6.21: Survey XP spectra of CuFeO 2 pellets subjected to different surface cleaning procedures: oxygen plasma (blue) and 3 h thermally cleaning at 200 ℃ in oxygen atmosphere (2 x 10 -3 mbar) (yellow).

Figure 6 . 22 :

 622 Figure 6.22: Cu2p 3/2 , Fe2p, O1s core level and valence band XP spectra and CuLMM Auger spectrum of CuFeO 2 pellets subjected to different surface cleaning procedures: oxygen plasma (blue) and 3 h thermally cleaning at 200 ℃ in an oxygen atmosphere (2 x 10 -3 mbar) (yellow).

10 Figure 6 . 25 :

 10625 Figure 6.25: UP spectra (He I excitation) of the CuFeO 2 substrate (blue) and the ITO film (yellow) after 240 s deposition time.

  .28. The band diagram shows that there is a staggered type II alignment between ITO and CuFeO 2 , where, upon light absorption, electrons would flow towards ITO due to the downward band bending. A vacuum level offset of 0.58 eV can be seen, which can be caused by a surface dipole or by Fermi level pinning. The possibility of a surface dipole at the CuFeO 2 surface can be cautiously neglected, because (as discussed above) the 2H-CuFeO 2 pellet ionization potential corresponds well to that of an {001} exposed 3R-CuFeO 2 single crystal determined at the point of zero charge under flat band conditions. 13 Contrarily, the 4.3 eV ITO work function is quite low, considering the value of 2.6 eV for E F -E VBM (Figure 6.24). However, this work function agrees well with what other people in our group have found for ITO films deposited at room temperature. Another possible reason for the observed vacuum level offset could be Fermi level pinning. Figure 6.28b shows the band positions of CuFeO 2 as they are commonly depicted, VBM at -1.0 --1.4 eV vs RHE and

13

 13

Figure 6 . 6 . 5 . 3

 6653 Figure 6.28: (a) Energy band diagram of CuFeO 2 /ITO with all energy values denoted in eV and the band gap of ITO taken from literature. 48 (b) Diagram showing the Fermi level tunability in CuFeO 2 based on the CuFeO 2 /ITO interface experiment. (c) A series of reduction potentials calculated at particular pH values on the RHE scale. i

Figure 6 .Figure 6 . 29 :Figure 6 . 30 :

 6629630 Figure 6.31 displays the UP spectra of CuFeO 2 before and after water exposure. First of all, the upward binding energy shift induced by the water exposure can also be noticed in the VB region of the UP spectra. Secondly, the work function of the CuFeO 2 pellets before water exposure, 5.2 eV, corresponds to what was found earlier and after water exposure the work function shifts towards 4.3 -4.4 eV We did not see this high work function shift for the BiVO 4 substrates. Most likely this is due to the CuFeO 2 substrate being a pellet, which has a higher porosity and is thus

Figure 6 . 31 :

 631 Figure 6.31: UP spectra (He I excitation) of two thermally cleaned (200 ℃, O 2 , 3 h) CuFeO 2 pellets before (blue) and after (yellow) water exposure.

Figure 6 . 32 :

 632 Figure 6.32: Difference spectra (green) of Cu2p3/2 , Fe2p 3/2 , CuLMM and Fe3p between thermally cleaned (200 ℃, O 2 , 3 h) CuFeO 2 pellet 1 before (blue) and after (yellow) water exposure. Shirley background subtraction was performed for Cu2p and Fe3p, and Tougaard background subtraction for Fe2p. The spectra were normalized and shifted towards the same binding energy.

Secondly, the interface

  experiments performed in this work have shown to give detailed information, not only on the in situ formed junction between substrate and contact material but also on the substrate itself. The interface experiments have shown that heterostructures may not only lead to the previously proposed benefits of improved reaction kinetics, enhanced surface passivation or corrosion protection, but may also yield increased band bending and, thus, better charge carrier separation. Future interface experiments could focus on the band alignment of contact materials with specific crystalline facets of BiVO 4 to elucidate whether other surface structures have a strong influence on the contact properties. Additionally, other promising photoabsorbers may be coated as well with high and low work function materials to investigate to what extent the Fermi level may be tuned in those compounds.

Pt4fFigure B. 2 :

 2 Figure B.2: Poudres de BiVO 4 hétérostructurées. À gauche: schémas illustrant la régiosélectivité du dépôt, milieu: images MEB et à droite: spectres caractéristiques XPS des dépôts.Les pourcentages en poids Ag, CoO x OH y , Pt, NiO x OH y correspondent à 10% en poids, 2% en poids, 1% en poids et 0,1% en poids, respectivement.

Figure B. 4 : 4 Figure B. 5 :

 445 Figure B.4: Propriétés photocatalytiques des poudres de BiVO 4 . (a) Dégradation photocatalytique de la rhodamine B (Rhd B) sous irradiation solaire AM1.5G avec illumination débutée à t = 0 min. La constante de vitesse moyenne du premier ordre, k mean , et le paramètre d'ajustement linéaire moyen, R 2 , sont indiqués sur le graphique. Pour chaque hétérostructure, il y avait une charge de 0,1% en poids du matériau de contact. (b) Oxydation de l'eau en présence d'un agent sacrificiel: dégagement gazeux en fonction du temps. Les taux de dégagement de gaz (mL h -1 O 2 ) sont également représentés. 1 mmol de NaIO 4 a été utilisé comme agent sacrificiel et la quantité de poudre de photocatalyseur était 50 mg.

Figure B. 6 :Figure B. 8 :

 68 Figure B.6: Alignement des bandes et accordabilité de niveau de Fermi déterminés par des expériences d'interface avec des granulés de CuFeO 2 . (a) Alignement de bande de CuFeO 2 /ITO, (b) accordabilité du niveau de Fermi dans CuFeO 2 sur la base de l'expérience d'interface CuFeO 2 /ITO et (c) une série de potentiels de réduction calculés à des valeurs de pH particulières sur l'échelle de l'électrode de référence à hydrogène (ERH).

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  21

  Morasch et al. found that by contacting Bi 2 O 3 with ITO and RuO 2 , the E F -E VBM barrier could be changed to 1.73 eV and 0.75 eV, respectively.49 Uddin et al. linked the increased photocatalytic dye decomposition efficiency for RuO 2 coated ZnO nanoparticles to the observation of a Schottky-type contact between ZnO and RuO 2 .50 Li et al.

	found that contacting ferroelectric (Ba,Sr)TiO 3 to the insulating Al 2 O 3 , led to a limited valence
	band discontinuity and a strong conduction band discontinuity of > 5 eV, resulting in a high
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Table 2 .

 2

		513.67	7.33
	V(III) oxide	515.29	7.33
	V(IV) oxide	515.84	7.33
	V(V) oxide	517.20	7.33

1: Spectral fitting parameters for the V2p spectrum of various vanadium oxide species.s. 166,167,169 Species V2p 3/2 (eV) V2p 1/2 -V2p 3/2 splitting (eV) V(0) 512.35 7.50 V(I) and/or V(II)

  In contrast, for the mixed Co 3 O 4 the satellite at 786 eV is less intense and an additional shake-up satellite at about 789.5 eV appears, which has been associated with Co 3+ .171,174,175

3/2 and Co2p 1/2 binding energies lie at about 780 eV and 793 eV, respectively, for cobalt hydr(oxide) species (Figure

2

.32). Thus, fitting the main Co2p 3/2 and/or Co2p 1/2 does not allow a reliable quantification of a mixed cobalt hydr(oxide) phase. However, the distinction between Co 2+ and Co 3+ can be made from the satellite structure between Co2p 3/2 and/or Co2p 1/2 . Co 2+ namely gives rise to a noticeable satellite at about 786 eV, as can be seen for CoO and Co(OH) 2 in Figure

2

.32.

Table 2 .

 2

	Species	Cu2p 3/2 (eV) Modified Auger parameter (Lit. Ave.)
	Cu(0)	932.61	1851.23
	Cu(I) oxide	932.43	1849.19
	Cu(II) oxide	933.57	1851.49
	2.4.2 Noble metal core level spectra
	2.4.2.1 Ag3d

2: Cu2p 3/2 binding energy and modified Auger parameter literature values for copper species.

166,177

Table 2 .

 2

	Ag	368.3	6.0
	Ag(I)O	367.9	6.0
	Ag(II)O	367.6	6.0
	Pt(OH) Species Pt4f 7/2 (eV) Pt4f 5/2 -Pt4f 7/2 (eV)
	Pt	71.1	3.3
	Pt(OH) 2	72.7	3.3
	PtO	74.0	3.3

3: Ag3d binding energies for Ag, Ag 2 O and AgO. 177,180 Species Ag3d 5/2 (eV) Ag3d 3/2 -Ag3d 5/2 splitting (eV) 2 can be found in Table 2.4. Table 2.4: Pt4f binding energies for Pt, PtO and Pt(OH) 2 . 177,180

Table 2 .

 2 5: Bi4f 7/2 binding energies for Bi(0) and Bi 2 O 3 .177,180

	Species Bi4f 7/2 (eV) Bi4f 5/2 -Bi4f 7/2 (eV)
	Bi(0)	156.9	5.3
	Bi 2 O 3	159.3	3.3

7/2 and Bi4f 5/2 . Bi 2 O 3 and Bi(0) can be easily distinguished from each other, as the Bi4f 7/2 and Bi4f 5/2 binding energies of Bi(0) are shifted downwards by about 2.4 eV with respect to the Bi4f 7/2 and Bi4f 5/2 binding energies of Bi 2 O 3 (Table 2.5).

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3.1 Synthesis heterostructured particulate photocatalysts

3.1.1 Chemicals used

Bismuth(III) nitrate pentahydrate (98%, Sigma-Aldrich), ammonium metavanadate ( 99.0%, Sigma-Aldrich), nitric acid (65%, Sigma-Aldrich), ammonium hydroxide (28 -30%, Sigma-Aldrich ), sodium dodecyl sulphate ( 98.5%, Sigma-Aldrich), iron(III) chloride hexahydrate ( 99%, Sigma-Aldrich), copper(I) iodide ( 99.5%, Sigma-Aldrich), sodium hydroxide (extra pure, Acros Organics), silver nitrate (>99.8%, Sigma-Aldrich), cobalt(III) nitrate hexahydrate ( 98%, Sigma-Aldrich), chloroplatinic acid hexahydrate ( 37.50% Pt basis, Sigma-Aldrich), sodium iodate (> 99%, Alfa Aesar), methanol ( 99.9%, Sigma-Aldrich), rhodamine B ( 95%, Sigma-Aldrich) were used as received. Pure (Type 3) and ultra pure water (Type 1) were provided by a Millipore Milli-Q Direct 8 water purification system.

Table 3 .

 3 1: Magnetron sputtering deposition parameters for RuO 2 , NiO, CoO x , ITO(RT) and ITO(400 ℃) thin films.

		RuO 2	NiO	CoO x ITO(RT) ITO(400 ℃)
	T sub (℃)	RT	RT	RT	RT	400
	Pr (Pa)	1	0.5	0.5	0.5	0.5
	O 2 /Ar ratio (%)	7.5	20	10	0	0
	P (W)	10 (DC) 40 (DC) 40 (RF) 25 (RF)	25 (RF)
	Flux (sccm)	10	20	20	6.6	6.6
	d (cm) R (nm.min -1 )	9.1 3	8.0 3	8.5 1.3	9.6 5	9.6 5

In the above Table T sub is substrate temperature, RT is room temperature, Pr stands for pressure, P is the power applied to the sputter target, d stands for the target-to-substrate distance and R stands for the deposition rate of the film.

  To correct for the spectrometer work function Φ sp , PES measurements of a sputter cleaned silver sample were recorded . In short, a piece of silver foil mounted onto a stainless steel sample holder was first sputter cleaned for at least ten minutes using an argon ion gun. An ion kinetic

	Bi4f		-9 mbar. For the
	XPS measurements a monochromatized Al Kα X-ray source (1486.6 eV) was used, while UPS was
	measured with He I irradiation (21.2 eV) emitted from a helium discharge lamp. UPS and XPS measurements were recorded at an emission angle (angle between sample and analyser) of 90°a nd 45°, respectively. Additionally, for UPS a bias of -4 V was applied to the sample in order to accelerate low kinetic energy photoelectrons towards the hemispherical analyser. O1s V2p Bi4d 5/2 Bi4d 3/2 Intensity (arb. units) V2s
	Bi5p	Bi5d V3p V3s
	600 500 400 300 200 100	0
	Binding energy (eV)
	In the other measuring mode, the high resolution mode, a certain energetic range of inter-
	est is sweeped multiple times to acquire a detailed photoelectron spectrum. High resolution XP
	spectra are mainly carried out to acquire detailed chemical and electronic information by mea-
	suring specific core photoelectron, Auger electron and valence band regions. High resolution XP
	spectra, in this work, were measured by performing a number of sweeps in the energetic range

energy of about 1 keV, a sample current of 0.9 µA and a sputtering area of 4 x 4 mm 2 were used as sputtering parameters. Afterwards, a survey (description see below) XP spectrum was measured to confirm the absence of carbonaceous species and water on the silver surface. Then, the silver Fermi edge as well as the Ag3d 5/2 core level spectrum were measured with the high resolution mode (description see below). The measured Ag3d 5/2 binding energy can, namely, also be used to calibrate the spectrometer as the Ag3d 5/2 binding energy of metallic silver should lie at about 368.26 eV. When performing UPS measurements, an additional silver calibration needs to be performed to calibrate the UP spectra, due to UPS being performed under slightly different conditions than XPS (see above). For the binding energy calibration, the silver Fermi edge was measured under the same conditions as the UP spectra.

Photoelectron spectra, in this work, have been measured according to one of two modes, survey and high resolution, whereby the spectra are denoted as survey spectra and high resolution spectra, respectively. In XPS, survey spectra are photoelectron spectra that are recorded rapidly over a broad range of binding energies to assess which chemical species are present at the sample surface, mainly to detect whether the sample surface is contaminated with undesired species. Survey XP spectra, in this work, were measured by performing 1 sweep from -1 eV to 1400 eV with a pass energy of 187.75 eV, an energy step size of 0.800 eV/step and a step duration of 100 ms. Survey XP spectra include signals from valence electrons, core electrons and Auger electrons, and have a background due to inelastically scattered electrons, which increases with increasing binding energy. Figure

3

.9 shows how the survey XP spectrum of a cleaned BiVO 4 surface looks like. of interest with a pass energy of 5.85 eV, an energy step size of 0.050 eV/step and a step duration of 100 ms.

  5/2 core level signals slightly better. The Bi4f 7/2 and Bi4f 5/2 line maxima are not influenced by the background removal functions.

	(arb. units)	Bi4f spectrum Shirley function Tougaard function Bi4f 5/2	Bi4f 7/2			Bi4f 5/2	Bi4f 7/2
	Intensity
		168	164	160	156	168	164	160	156

  3/2 and O1s core level spectra measured during the water exposure and water desorption experiments, peak fitting was performed to determine the relative amount of components in the V2p 3/2 and O1s spectra. All components were fitted using pseudo-Voigt profiles. The fitting parameters of V 5+ and O 2-(in BiVO 4 ) were determined by fitting the V2p 3/2 and O1s spectra, BiVO 4 sample calcined in DAISY-FUN there was a second O1s component due to leftover carbonaceous species, which was fitted as well. The fitting parameters of the other components V 4+ , OH (DAISY-MAT) and H 2 O (DAISY-FUN), found after water exposure/water adsorption/water desorption were fitted while keeping the previously determined fitting parameters for V 5+ , O 2-, OH/CO (DAISY-FUN) fixed. All fitting parameters can be seen in Table3.2.

		Table 3.2: Fitting parameters O1s and V2p 3/2 components.
				DAISY-MAT			DAISY-FUN
	Line V2p 3/2 V2p 3/2 O1s	Signal S1(V 5+ ) S2(V 4+ ) S1(O 2-)	L/G FWHM ∆BE (eV) L/G FWHM ∆BE (eV) 0.45 1.1 -0.64 0.9 -0.32 1.7 -0.76 0.32 1.7 -0.76 0.12 1.1 -0.12 1.0 -
	O1s	S2(OH/CO) 0.12	1.9	1.05	0.12	2.8	1.25
	O1s	S3(H 2 O)	-	-	-	0.12	1.6	3.3

relatively, that were measured after the calcination of the BiVO 4 sample. The 0.12 Lorentzianto-Gaussian ratio (L/G) determined for O 2-was also used for all other O1s components. For the Sx: signal x, L/G: Lorentzian-to-Gaussian ratio, ∆BE: Binding energy between Sx and S1, FWHM: full width at half maximum.

  Most catalytic tests were run for 20 h. (14) Anton, R.; Wiegner, T.; Naumann, W.; Liebmann, M.; Klein, C.; Bradley, C. Design and performance of a versatile, cost-effective microwave electron cyclotron resonance plasma source for surface and thin film processing. Rev. Sci. Instrum. 2000, 71, 1177. (15) Bein, N. The Fermi Level in Bismuth Ferrite., Ph.D. Thesis, TU Darmstadt, 2018, p 117. BiVO 4 /MnO x , BiVO 4 /Co 3 O 4 could boost the water oxidation performance of bare BiVO 4 and that the creation of heterostructures with both a reduction and oxidation cocatalyst on different facets, like BiVO 4 /Pt/Co 3 O 4 , yielded an even higher water oxidation performance. We wanted to first reproduce the BiVO 4 /Pt, BiVO 4 /Ag and BiVO 4 /Co 3 O 4 heterostructures to gain knowledge on how to create these kind of structures, so that we could create similar heterostructures with the other photoabsorber studied in this work, CuFeO 2 . Additionally, a new heterostructured BiVO 4 powder was prepared through photodeposition of NiO x OH y . NiO x OH y was chosen, because nickel based contact materials have already shown their merit for water photo-oxidation. Mccrory et al. showed that Ni based compounds in a basic medium have electrocatalytic oxygen evolution overpotentials which are among the lowest and on par with state of the art noble metal catalysts. 3 Furthermore, nickel based compounds have already shown to increase the water oxidation performance of BiVO 4 thin films, whereby

	(16) Weidner, M. Fermi Level Determination in Tin Oxide by Photoelectron Spectroscopy: Re-lation to Optoelectronic Properties; Band Bending at Surfaces and Interfaces; Modulation Chapter 4
	Doping., Ph.D. Thesis, TU Darmstadt, 2016, p 427.
	(17) Jaegermann, W.; Kaiser, B.; Ziegler, J.; Klett, J. In Photoelectrochem. Sol. Fuel Prod. Springer International Publishing: Cham, 2016, pp 199-280. Anisotropic BiVO 4 Heterostructures
	(18) Tengeler, S. Cubic silicon carbide for direct photoelectrochemical water splitting., Ph.D.
	Thesis, TU Darmstadt, 2017, p 183.
	4.1 Introduction

(19) Silversmit, G.; Depla, D.; Poelman, H.; Marin, G. B.; De Gryse, R. Determination of the V2p XPS binding energies for different vanadium oxidation states (V 5+ to V 0+ ). J. Electron Spectros. Relat. Phenomena 2004, 135, 167-175.

(20) Demeter, M.; Neumann, M.; Reichelt, W. Mixed-valence vanadium oxides studied by XPS. Surf. Sci. 2000, 454-456, 41-44.

(21) Moulder, J. F.; Stickle, W. F.; Sobol, P. E.; Bomben, K. D., Handbook of X-ray photoelectron spectroscopy: a reference book of standard spectra for identification and interpretation of XPS data; Physical Electronics: Eden Prairie MN, 1992, p 261. As mentioned in the main Introduction in Section 2.3.1, photodeposition is an effective way to create dissymmetric heterostructured powders. The scientific articles published by Li et al. are essentially groundbreaking, as they showed that through photodeposition, oxide and reduction cocatalysts could be loaded onto different facets of a {010}/{110} dual facet exposed BiVO 4 microcrystalline powder. 1,2 Furthermore, they found that the creation of heterostructures with a single cocatalyst on a particular facet, like BiVO 4 /Pt,

Table 4 .

 4 1: Reaction parameters for the synthesis of monoclinic BiVO 4 powders with a truncated bipyramidal morphology, including the literature references to the synthesis procedures.

		Reaction mixture		Autoclave	Reference
	Bi precursor	V precursor	Surfactant	pH T (℃) t (h)
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	SDS	7	180	24	[7]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	EDTA	7	180	24	[8]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	-	2	200	24	[2]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	-	2	200	24	[9]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	-	2	200	24	[1]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	-	2	180	24	[10]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	-	0.6	180	24	[11]
	Bi(NO 3 ) 3 .5H 2 O	NH 4 VO 3	TiCl 3	2	180	24	[12]
	BiCl 3	NH 4 VO 3	ethanolamine 6.15	160	12	[13]
	Bi 2 O 3

4 123Table 4 .

 4 2: Experimental details for the BiVO 4 based heterostructured powders fabricated through photodeposition.

	Heterostructure	m BiVO 4 (mg)	Precursor	Hole/electron scavenger	pH	V (mL)	t illum. (h)
	BiVO 4 /Ag	100	AgNO 3	MeOH (25 vol%)	7	20	5
	BiVO 4 /Pt BiVO 4 /CoO x OH y	100 100	H 2 PtCl 6 Co(NO 3 ) 2	-NaIO 3 (0.01 mol/L)	Various 7	20 20	4.5 3
	BiVO 4 /NiO x OH y	100	NiCl 2	-	9.2 (KBi buffer)	20	2
	BiVO 4 /RuO x	100	Ru(acac) 3	FeCl 3	-	20 (THF/ CH 3 COOH)	3

t illum. : solar simulator illumination time. MeOH: methanol, KBi: potassium borate; THF: tetrahydrofuran, acac: acetylacetonate.

  Therefore, we opted to denote the deposited cobalt species as CoO x OH y .

		Co2p 3/2		O1s		V2p 3/2		Bi4f
	Intensity (arb. units)	Co 2+ Co 3+
		790 785 780 775	536	532	528	520	515	170	160
				Binding energy (eV)

  2+ reduction for the 0.1 wt% nickel solution, which has a lower reduction potential, could be thermodynamically restricted. Remarkably, because nickel was preferably found on the {010} facets for the 0.1 wt% nickel solution, this would mean that the oxidation of Ni 2+ at a pH of 9.2 happens preferably on the {010} facets, which are normally considered to be the more reducing facets of truncated bipyramidal BiVO 4 particles.1 It could possibly be that the regioselectivity of the photodeposition on BiVO 4 is pH sensitive. Figure 4.30 shows the high resolution XP core level spectra for the BiVO 4 /NiO x OH y powders.The sample with ± 10 wt% NiO x OH y is not included, as it was charging to a degree that the Ni2p spectrum could not give any accurate information any more Inhomogeneous surface charging can still be noticed in the Bi4f spectrum for the 1 wt% NiO x OH y sample with extra bumps being present at higher binding energies. The XP core level spectra show that more nickel is present in the BiVO 4 /NiO x OH y (1 wt%) sample since the Ni2p-to-Bi4f intensity ratio is higher for the BiVO 4 /NiO x OH y (1 wt%) sample compared to the BiVO 4 /NiO x OH y (0.1 wt%) sample. Similar

	pH	(4.10)
	Again, for a 1 wt% nickel solution, which contains 0.85 mmol/L nickel, and a pH of 9.2:
	E NiOOH/Ni 2+ ,1wt%Ni (RHE) = 2.05 V -0.118 V x log(0.00085) -0.118 V x 9.2	(4.11)
	E NiOOH/Ni 2+ ,1wt%Ni (RHE) = 1.33 V	(4.12)
	For a 0.1 wt% nickel solution, which contains 0.085 mmol/L nickel, and a pH of 9.2:
	E NiOOH/Ni 2+ ,0.1wt%Ni (RHE) = 2.05 V -0.118 V x log(0.000085) -0.118 V x 9.2	(4.13)
	E NiOOH/Ni 2+ ,0.1wt%Ni (RHE) = 1.44 V	(4.14)
	When considering the calculated reduction potentials, and the VBM and CBM of BiVO 4 at 0.1
	eV and at 2.5 eV, respectively, 56 it can be evaluated whether the reduction and/or oxidation of
	Ni 2+ by BiVO 4 is thermodynamically allowed. The NiOOH/Ni 2+ oxidation reaction can take place
	for both 1 wt% and 0.1 wt% nickel solutions since the BiVO 4 VBM is much more positive than the

calculated reduction potentials of both solutions. Contrarily, the BiVO 4 CBM is situated closely to the reduction potentials calculated for the Ni 2+ /Ni reduction reaction. Thus, the reduction of Ni

2+

when utilizing a 1 wt% nickel solution could still take place, whereas the Ni

  Figure 4.30: Ni2p, Bi4f, V2p 3/2 and O1s core level XP spectra of BiVO 4 /NiO x OH y samples.

	Ni2p			O1s	V2p 3/2		Bi4f
	Intensity (arb. units)	1 wt% NiO OH x 0.1 wt% NiO OH y x	y
	880 870 860 850	540	530	520	515	170 165 160 155
				Binding energy (eV)
	Ni(OH) 2 (s) + 2 e -	Ni(s) + 2 OH -
	Due to the very low standard reduction potential of Ni(OH) 2 /Ni, -0.714 V, 55 the photo-
	oxidation of Ni is thermodynamically allowed and likely to happen.

  Before analysing the photocatalytic performance of the different heterostructured BiVO 4 samples, the differences in light absorption were studied through UV-Vis diffuse reflectance spectroscopy. Remarkably, the Tauc plot in Figure4.33 shows that the deposited contact materials do not have a strong influence on the optical band gap of BiVO 4 , because all heterostructures appear to have an optical band gap in the range of 2.48 -2.52 eV.

	4.3.6 Photocatalysis
	4.3.6.1 UV-Vis diffuse reflectance
						1 wt% Ni	(a)	1 wt% Ni	(b)
	F(R) (arb. units)	0.5 1.0 2.0 1.5				2 wt% Co 10 wt% Pt Bare 10 wt% Ag	[F(R)hν] 2 (arb. units)	0.5 1.0	2 wt% Co 10 wt% Pt Bare 10 wt% Ag
		0.0	400	500	600	700	800	2.40 2.45 2.50 2.55 2.60 2.65 2.70 0.0
				λ (nm)			hν (eV)
		Fe2p			V2p 3/2	C1s + Ru3d	Bi4f
	Intensity (arb. units)							Ru3d
	730 720 710 700	520		515	290	280	170	160
							energy (eV)
	Figure 4.32: Fe2p, V2p 3/2 , C1s, Ru3d and Bi4f core level XP spectra of BiVO 4 sample after pho-
	todeposition of Ru(acac) 3 with FeCl 3 as electron scavenger in THF.

y (0.1 wt%)dark 4 BiVO /NiO

	BiVO 4	/	Bare	Pt	NiO x OH y	CoO x OH y
	mL h -1 O 2	0.19	1.10	0.36	0.37	1.88 1.61
	μmol h -1 O 2	7.7 45.1	14.6	15.2	76.9 66.0
	V (mL)
	120	180	240	300	360
			Time (min)

x OH

  Ye, F.; Li, H.; Yu, H.; Chen, S.; Quan, X. Constructing BiVO 4 -Au@CdS photocatalyst with energic charge-carrier-separation capacity derived from facet induction and Z-scheme bridge for degradation of organic pollutants. Appl. Catal. B Environ. 2018, 227, 258-265. (11) Liu, Q.; Liu, Y.; Gao, B.; Chen, Y.; Lin, B. Hydrothermal synthesis of In 2 O 3 -loaded BiVO 4 with exposed {010}{110} facets for enhanced visible-light photocatalytic activity. Mater. Wang, D.; Jiang, H.; Zong, X.; Xu, Q.; Ma, Y.; Li, G.; Li, C. Crystal Facet Dependence of Water Oxidation on BiVO 4 Sheets under Visible Light Irradiation. Chem. -A Eur. J. 2011, Xi, G.; Ye, J. Synthesis of bismuth vanadate nanoplates with exposed {001} facets and enhanced visible-light photocatalytic properties. Chem. Commun. 2010, 46, 1893. (14) Tan, H. L.; Tahini, H. A.; Wen, X.; Wong, R. J.; Tan, X.; Iwase, A.; Kudo, A.; Amal, R.; Smith, S. C.; Ng, Y. H. Interfacing BiVO 4 with Reduced Graphene Oxide for Enhanced Photoactivity: A Tale of Facet Dependence of Electron Shuttling. Small 2016, 12, 5295-5302. (15) Tan, H. L.; Wen, X.; Amal, R.; Ng, Y. H. BiVO 4 {010} and {110} Relative Exposure Extent: Governing Factor of Surface Charge Population and Photocatalytic Activity. J. Phys.

	Res. Bull. 2017, 87, 114-118.
	(12) 17, 1275-1282.
	(13)

Chem. Lett. 2016, 7, 1400-1405.

Table 5 .

 5 1: Areas of Bi4f and V2p 3/2 core levels, calculated surface V/Bi ratios, and V2p 3/2 to Bi4f 7/2 binding energy difference for a series of PLD BiVO 4 thin films.Bi4f tab. and V2p 3/2,tab. are the respective Bi4f and V2p 3/2 areas divided by the tabulated sensitivity factors 7.632 and 1.275, respectively, whereby the surface V/Bi ratios were calculated from these areas. V/Bi corr. ratio was calculated by multiplying the V/Bi ratio with the correction factor 1.25 determined from the XPS measurements on the (010) exposed Mo:BiVO 4 single crystal.

	n 0	Bi4f Bi4f tab. V2p 3/2 V2p 3/2,tab. xBi xV V/Bi V/Bi corr. ∆BE V2p 3/2 ,Bi4f 7/2 (eV)
	1 48488	6353	6068	4759	0.57 0.43 0.75	0.94	357.68
	2 43087	5646	5422	4253	0.57 0.43 0.75	0.94	357.69
	3 63905	8373	7361	5773	0.59 0.41 0.69	0.86	357.80
	4 67541	8850	7754	6082	0.59 0.41 0.69	0.86	357.64
	5 43777	5736	5694	4466	0.56 0.44 0.78	0.97	357.72
	6 40231	5271	5244	4113	0.56 0.44 0.78	0.98	357.67
	7 47049	6165	5369	4211	0.59 0.41 0.68	0.85	357.69
	-	-	-	-	Mean	0.58 0.42 0.73	0.91	357.70

  3.2.2. In this Subsection the results of the BiVO 4 /RuO 2 and BiVO 4 /ITO interface experiments, carried out with the gradient BiVO 4 thin film substrates, and the BiVO 4 /CoO x and BiVO 4 /NiO interface experiments, performed with PLD BiVO 4 thin films, are described.

Table 5 .

 5 2: Summary of important energetic values of all BiVO 4 based interface experiments.

		BiVO 4 /RuO 2	BiVO 4 /ITO (RT)	BiVO 4 /ITO (400 ℃)	BiVO 4 /NiO BiVO 4 /CoO x
	BE Bi4f 7/2 ,0 s BE Bi4f 7/2 ,final BE V2p 3/2 ,0 s BE V2p 3/2 ,final VBM BiVO 4 WF BiVO 4 VBM CL	159.18 158.46 516.78 516.08 1.8 5.7 0	159.08 159.22 516.74 516.92 1.8 5.6 2.5	159.00 159.31 516.71 517.00 1.8 5.8 2.8	159.03 158.25 516.71 515.93 1.8 5.9 0.4*	158.65 158.22 516.29 515.82 1.5 6.6 0.3
	WF CL	6.5	4.4	4.5	5.4	5.7

  Figure 5.22: Bi4f 7/2 spectra for the BiVO 4 /RuO 2 , BiVO 4 /NiO and BiVO 4 /CoO x interface. First, the background of the Bi4f core level spectra was subtracted using a Tougaard function.

	Bi4f 7/2	Bi4f 7/2	Bi4f 7/2
	RuO BiVO 4 / 2	NiO BiVO4 /	BiVO 3 4 Co O 4 /
	Deposition	Deposition	Deposition
	time	time	time
	120 s	30 s	135 s
	60 s	10 s	45 s
	30 s	5 s	15 s
	15 s	0 s	0 s
	0 s
	162 161 160 159 158 157 156	162 161 160 159 158 157 156 162 161 160 159 158 157 156
		Binding energy (eV)

  3/2 , O1s and valence band XP spectra of BiVO 4 is shown and in Table5.3 a summary of the most important information obtained from this water exposure experiment is given. There are a few remarkable observations one can draw from the water exposure experiment on BiVO 4 . First of all, an in-

	crease in binding energy of about 0.3 -0.5 eV can be noticed for the Bi4f, V2p and VBM spectra
	between a cleaned sample (after calcination/oxygen plasma) and the same sample exposed to
	water. After the water exposure a Bi4f 7/2 and V2p 3/2 binding energy of 159.6 eV and 517.3 eV, re-
	spectively, is retrieved, which is slightly higher than the respective 159.2 eV and 517.0 eV binding
	energies, observed after ITO was sputtered onto BiVO 4 . However, the cleaned BiVO 4 substrate,
	before water exposure, had already high Bi4f 7/2 and V2p 3/2 binding energies of 159.2 and 516.8

Table 5 .

 5 3: Summary of BiVO 4 water exposure experiment. : Integrated intensity of core level spectrum x, I V 5+ /I V 4+ : ratio of integerated intensities of fitted V2p 3/2 components, I O 2-/I OH : ratio of integerated intensities of fitted O1s components, O 2 : oxygen flow (2 sccm, 1 x 10 -3 mbar), H 2 O: water exposure.

	Treatment	as is	2 h, 400 ℃, O 2	H 2 O	30 min, RT, O 2	15 min O 2 plasma	H 2 O
	BE Bi4f 7/2 (eV) BE V2p 3/2 (eV) BE O1s (eV)	159.5 517.1 530.2	159.2 516.8 529.9	159.6 517.3 530.5	159.6 517.3 530.4	159.3 517.0 530.1	159.6 517.3 530.5
	I V2p 3/2 /I Bi4f I O1s /I V2p 3/2 I O 2-/I OH I V 5+ /I V 4+	0.78 4.4 --	0.90 3.9 --	0.94 3.3 1.96 0.39	0.89 3.7 4.10 0.61	0.94 4.0 --	0.99 3.2 1.91 0.38
	BE: Binding energy, I

x

  2-component increases during the water adsorption and decreases after water desorption. This could mean that water already

	14.7 eV SEC			2 O Plasma 15 min
	15.8 eV			30 min	O 2
	16.0 eV			H O 2
	15.7 eV			2 h, 400 10 mbar °C O 2 -3
	15	10	5	0

  1 eV through the formation of junctions with high and low work function materials. This remarkably high Fermi level tunability convincingly demonstrates that bulk Fermi level pinning is limited within BiVO 4 and explains why rather high photovoltages can be obtained with BiVO 4 . -E VBM , as probed by contacting BiVO 4 at elevated temperature with a low work function material, ITO, Fermi level pinning seems to take place at E F -E VBM > 2.0 -2.1 eV since a reduction of bismuth and vanadium was observed when the BiVO 4 /ITO contact was established. This Fermi level pinning state has most probably no severe implications for reduction reactions in photoelectrochemical cells since BiVO 4 is usually not subjected to cathodic potentials. However, it may limit the effective n-type doping concentration which can be attained by incorporating donor impurities. Furthermore, the assumed transitivity of the RuO 2 and ITO interface experiments allowed to align BiVO 4 relative to other oxides studied for photo(electro)chemical water splitting, showing that the alignment based on solid/electrolyte and solid/solid interfaces differ from each other.

	The high upwards band bending, 0.45 eV, 0.71 eV and 0.78 eV observed for the BiVO 4 /CoO x ,
	BiVO 4 /RuO 2 and BiVO 4 /NiO heterostructures, respectively, strongly indicates that high work
	function transition metal oxides very likely enhance the separation of charge carriers within
	BiVO 4 upon light absorption. Thus, the charge carrier separation effect of a contact material
	may absolutely not be disregarded when assessing the photo(electro)chemical water splitting
	efficiency of a heterostructure. It may even be the main reason for improved water splitting
	activities.
	For increasing E

F

Table 6 .

 6 Fe2p and O1s core level XP spectra and CuLMM Auger spectrum of the batch of CuFeO 2 powder used for the photodeposition experiments. 1: Experimental details for the CuFeO 2 based heterostructured powders fabricated through photodeposition. The low silver coverage can be related to the rather large size (50 -200 nm) of the deposited silver particles, indicating that growth instead of nucleation was preferred during the photodeposition process. The low coverage and particle size of the silver deposit agree quite well with what was found for the silver photodeposition experiments on BiVO 4 (Section 4.3.1).

	Cu2p			Fe2p		CuLMM	O1s
	Intensity (arb. units)			Fe 3+	In3p 1/2
	960 950 940 930	740 730 720 710 700	580 575 570 565 560	535	530	525
				Binding energy (eV)
	Figure 6.9: Cu2p, Heterostructure	m CuFeO 2 (mg)	Precursor	Hole/electron scavenger	pH	V (mL)	t illum. (h)
	CuFeO 2 /Ag	100		AgNO 3	MeOH (25 vol%)	7	20	5
	CuFeO 2 /Pt	100		H 2 PtCl 6	MeOH (40 vol%)	2.9	20	5
	CuFeO 2 /NiO x OH y	100		NiCl 2		-	9.2 (KBi buffer)	20	2
	AES was carried out to see whether more information on the regioselectivity of the silver
	deposit could be obtained. Due to the previously observed problems with AES mapping on BiVO 4
	samples (Section 4.3.2), i.e. sample degradation and long measurement times, we opted to not
	perform any mapping and instead did several rapid point measurements on certain positions
	on the CuFeO 2 /Ag (10 wt%) sample (Figure

t illum. : solar simulator illumination time facets than on the sides, which is however difficult to conclude with high certainty because the silver coverage appears to be quite low despite a 10 wt% Ag solution being used in the photodeposition process.

  so that inhomogeneous sample charging during XPS measurements on CuFeO 2 should be less prevalent. Indeed, in the CuFeO 2 /Ag core level spectra no indications of inhomogeneous sample charging can be found, so that the recorded Ag3d spectrum can be reliably used to define the silver species. The Ag3d spectrum here consists of two symmetric signals at 368.2 eV and 374.3 eV which belong to the characteristic binding energies of Ag3d 5/2 and Ag3d 3/2 , respectively, of metallic silver,24 thus proving that the silver photodeposition involves the reduction of Ag + to Ag(0).

	Cu2p	Fe2p	CuLMM	O1s	Ag3d
	Intensity (arb. units)	In3p 1/2
	960 950 940 930 740 730 720 710 700 580 575 570 565560	535 530 525	380 376 372 368
		Binding energy (eV)
	Figure 6.12: Cu2p, Fe2p, O1s and Ag3d core level XP spectra and CuLMM Auger spectrum of the
	CuFeO 2 /Ag (10 wt%) sample.

  O1s and Pt4f core level XP spectra and CuLMM Auger spectrum of the CuFeO 2 /Pt (10 wt%)-pH 2.9 sample. some of the nanoplatelets of the batch of bare CuFeO 2 (Figure6.7). The nickel deposit seems to be distributed all over the CuFeO 2 nanoplatelets, just as was seen for the BiVO 4 /NiO x OH y (10 wt%) and BiVO 4 /NiO x OH y (1 wt%) sample.

	Cu2p	Fe2p	CuLMM	O1s	Pt4f
	Intensity (arb. units)	In3p 1/2		PtO
	960 950 940 930 740 730 720 710 700	580 575 570 565560	535 530 525	80 76 72 68
		Binding energy (eV)
	Figure 6.15: Cu2p, Fe2p,

  The Ni2p spectrum is quite similar to the Ni2p spectrum measured for BiVO 4 /NiO x OH y (Figure4.30 in Section 4.3.4), indicating that the nickel species consists of oxides and hydroxides, which is why the nickel deposit herein is denoted as NiO x OH y . Despite the presumed photoreduction of Ni 2+ to metallic nickel, no sign of metallic nickel, which has a characteristic Ni 2p 3/2 binding energy of 852.6 eV, can be seen in the Ni2p spectrum. The reason is that metallic nickel would readily oxidise to Ni(OH) 2 , as explained in Section 4.3.4. Figure6.17: Cu2p, Ni2p, Fe2p and O1s core level XP spectra and CuLMM Auger spectrum of the CuFeO 2 /NiO x OH y (± 10 wt%) sample.

	Cu2p	Ni2p	Fe2p	CuLMM	O1s
	Intensity (arb. units)		In3p 1/2
	960 950 940 930	890 880 870 860 850	740 730 720 710 700	580 575 570 565560	535 530 525
		Binding energy (eV)

  carbon species. However, after the oxygen plasma treatment extra signals in the Cu2p region of the survey spectrum can be seen, which indicate a change in the CuFeO 2 surface chemistry. Additionally, an extra signal at 1071.2 eV can be identified, which is characteristic for the Na1s core level electrons of sodium. Thus, sodium seems to segregate to the pellet surface after the heat treatment. Most likely, the small amount of sodium is a leftover from the sodium hydroxide that was used in the 2H-CuFeO 2 powder synthesis. The presence of sodium was unexpected, because no sodium was detected while measuring the XP spectra of the heterostructured CuFeO 2 powders. Possibly, sodium is incorporated into the bulk of the CuFeO 2 particles and may explain the bad hydrogen evolution efficiencies, since bulk incorporated impurities can act as charge carrier recombination centres. However, photodeposition of Ag, Pt and NiO x OH y took place readily, despite those light-induced reactions being negatively influenced by charge carrier recombination as well. The small amount of sodium at the surface was not expected to interfere with the

	interface experiments.
		Pellet 2 3 h, 200 °C, O 2	Na1s Cu2s	OKLL	Cu2p	FeLMM	Fe2p	O1s
	Intensity (arb. units)	Pellet 2 As is Pellet 1 O plasma 2					CuLMM	C1s	Fe3p Cu3p Fe3s Cu3s
		Pellet 1 As is
		1400	1200		1000	800	600	400	200	0
							Binding energy (eV)

  As the 2H-CuFeO 2 pellets already showed a low E F -E VBM value of 0.19 eV after the thermally

		Cu2p 3/2	Fe2p		CuLMM		O1s	VB
								(eV) E VBM
						Pellet 2
	Normalized intensity	(arb. units)				Pellet 1 Pellet 2 As is O plasma 2 3 h, O , 200 °C 2	0.41 0.19
								0.63
						Pellet 1 As is	0.72
		950 940 930 740	720	700	570	560	535 530 525	10 8 6 4 2 0	-1
				Binding energy (eV)

cleaning step, interface experiments with high work function materials, to lower E F -E VBM further, were not expected to give any new information, so that only interface experiments with a low work function material, ITO, were performed. Testing by how much E F -E VBM in 2H-CuFeO 2

  Figure 6.23: High resolution powder X-ray diffractograms for the CuFeO 2 powder used to prepare the pellets (blue) and a pellet after the thermally cleaning step at 200 ℃ (yellow). The (hkl) reflections of 2H-CuFeO 2 are indicated in the graph. From the secondary electron cut-offs of the measured UP spectra (Figure 6.25), work functions of 5.2 eV and 4.3 eV were determined for the CuFeO 2 pellet and the ITO film, respectively. The ITO work function, here, is similar to the ITO work functions measured earlier for the BiVO 4 /ITO interface experiments (Section 5.3.1). When combining the 5.2 eV CuFeO 2 work function with the value of 0.2 eV for E F -E VBM , an ionization potential of 5.4 eV is retrieved for the CuFeO 2 pellet. Ginley equation correlate well with experimental values when the valence band consists mainly of O2p character, but not when the valence band has mainly metal 3d character, as is the case for example for CuO and NiO. 43 Hence, the strong discrepancy in calculated and experimentally determined ionization potentials most likely has to do with the Cu3d character of the CuFeO 2 valence band. More recently, Omeiri et al. determined through Mott-Schottky flat band measurements that the ionization potential of the (001) surface of a 3R-CuFeO 2 single crystal amounts to 5.4 eV, 13 which is quite similar to the 5.2 eV ionization potential determined here for the [001] oriented CuFeO 2 pellet. Due to the Mott-Schottky measurements being performed at the point of zero charge, the similarity of the CuFeO 2 surfaces and the similarity of the ionization potentials, the presence of a strong surface dipole at the 2H-CuFeO 2 pellet surface is regarded as being highly unlikely.Difference spectra of Cu2p 3/2 , CuLMM and Fe2p 3/2 were made (Figure 6.26) to verify whether the ITO interface experiment led to any chemical changes in CuFeO 2 . In the Cu2p 3/2 core level spectrum a broadening of the main Cu2p 3/2 line can be seen, which is at first glance a somewhat strange phenomenon, because oxidation state changes normally lead to the appearance of new signals with a shifted binding energy. Anyway, the different copper oxidation states cannot be easily recognized by analysing the main Cu2p 3 line, as the Cu2p spectrum of Cu + only differs

	Cu2p		Fe2p In3p 1/2	(004) CuLMM			O1s		In3d 5/2	VB * 3R-CuFeO 2
												E VBM
		(002)	Sn3p 3/2		(101)	(102)	(103)	(104) (006)	(105) 240 s (106) (110) 32 s	(008) (107) (114) (201) (202) (110)	pellet 200 °C 2 + 3 h, O , 2.62 (eV)
												0.49
	Intensity (arb. units)	10	20	30	*	40	50	8 s 16 s	60	70	80 powder 90 0.46 0.50
										4 s		0.34
												0.18
										0 s
	line shape. For the VBM and Cu2p 3/2 a complete upwards binding energy shift of 0.32 eV and 740 720 700 580 570 560 534 530 448 446 444 442 10 8 6 4 2 0 960 950 940 930 -1
	0.37 eV, respectively, is obtained. A second interface experiment showed very similar shifts of Binding energy (eV)
	0.37 eV and 0.44 eV for VBM and Cu2p 3/2 , respectively.
	Figure 6.24: Cu2p, Fe2p, O1s, In3d 5/2 core level and valence band XP spectra, and CuLMM Auger
	spectra for CuFeO 2 /ITO(RT) interface. ITO deposition times are denoted in the O1s
	spectra.
	Contrarily, Benko et al. calculated an ionization potential of 4.95 eV from the flatband potential, derived from photocurrent onset potential measurements at pH ZPC , of a polycrystalline CuFeO 2 I P = (χ (Cu) * χ (Fe) * χ (O) 2 ) 1 6 + 1 2 E g (6.1)
	pellet, which was prepared through conventional solid-state synthesis. 25 However, the authors with E g being the electronic band gap, herein given by the 1.24 eV indirect optical band gap:
	report optical band gaps, 1.15 eV (indirect) and 2.03 eV (direct), which differ strongly from the
	optical band gaps, 1.24 eV (indirect) and 1.43 eV (direct), determined for our 2H-CuFeO 2 pow-der. This band gap dissimilarity could be due to structural and electronic differences, which can I P = (4.48 * 4.06 * 7.54 2 ) 1 6 + 1 2 1.24 = 6.29 (6.2)
	explain the difference in ionization potentials. Another way of estimating the ionization poten-This calculated ionization potential differs largely from what we obtained using photoelectron
	spectroscopy and from the CuFeO

tial is through a semi-empirical calculation developed by Butler and Ginley,

41

which involves the Mulliken electronegativity of CuFeO 2 , calculated from the absolute electronegativities χ of copper (4.48 eV), oxygen (7.54 eV) and iron (4.06 eV):

42

2 valence band positions mentioned in literature.

1,10

Koffyberg et al. found out that the ionization potentials of metal oxides calculated with the semi-empirical Butler-

  Figure 6.26: Difference spectra of Cu2p 3/2 , CuLMM and Fe2p 3/2 between bare CuFeO 2 and CuFeO 2 /ITO (4 s). Shirley background subtraction was done for Cu2p and Tougaard background subtraction for Fe2p. The spectra were normalized and shifted towards the same binding energy.Figure 6.27: Roughness profile of PLD BiVO 4 thin film and 2H-CuFeO 2 pellet, measured with a Dektak XT Advanced System profilometer tentials are shown on the RHE scale. i As was shown above, Fe 3+ is evidently partly reduced to Fe 2+ after establishing the CuFeO 2 /ITO interface. The Fe 3+ /Fe 2+ (pH 7) reduction potential, at about -1.18 eV vs RHE, cannot explain the induced Fe 3+ reduction at the CuFeO 2 /ITO contact, as this reduction potential would mean that the iron would already be in the Fe 2+ state for bare CuFeO 2 , with a Fermi level position at about -1.02 eV vs RHE. However, the Fe 3+ coordination in solution and in the 2H-CuFeO 2 crystal lattice differs from one another, meaning that the actual Fe 3+ /Fe 2+ charge transfer level in 2H-CuFeO 2 probably does not correspond to the Fe 3+ /Fe 2+ aqueous reduction potential. Recently, Lohaus et al. demonstrated that the Fe 3+ /Fe 2+ polaron level for hematite, lies at -0.3 eV vs NHE, 46 which on the RHE scale will be -0.7 eV vs RHE (pH 7). Now, for hematite the iron sites are octahedrally coordinated by oxygen just like the iron sites in CuFeO 2 . Therefore, the Fe 3+ /Fe 2+ polaron level in CuFeO 2 can expected to be situated around -0.7 eV vs RHE (pH 7) as well, a position which would agree nicely with the Fermi level shift feasible by ITO deposition. Furthermore, Prévot et al. also observed Fermi level pinning in 3R-CuFeO 2 at

			Cu2p 3/2			Fe2p 3/2				CuLMM
				bare CuFeO	2
	Intensity (arb. units)			4 s ITO Difference
	0
	950 945 940 935 930 925	720	716	712	708	580	575	570	565	560
					Binding energy (eV)
		100 120 140							4 PLD BiVO thin film 2H-CuFeO pellet
		80
		60
	Profile (nm)	-40 -20 0 20 40
		-60
		-80
		-100
		-120
		-140
		0.00	0.05	0.10	0.15		0.20	0.25		0.30	0.35
					Scan length (mm)

2

  7 -0.8 eV after water exposure. For Fermi level pinning

		Cu2p 3/2		Fe2p 3/2		CuLMM		Fe3p
	Intensity (arb. units)							Difference CuFeO 2 CuFeO + H O 2 2
	0
	950	940	930	720	710	580	570	560	60	50
					Binding energy (eV)

  /Fe 2+ correspond bien à ce que d'autres chercheurs ont trouvé dans notre groupe pour d'autres composés contenant du fer. Ainsi, la raison pour laquelle nous n'avons pas observé de craquage de l'eau significative avec les poudres de CuFeO 2 hétérostructurées est très probablement due au fait que CuFeO 2 ne peut pas atteindre des phototensions élevées en raison du blocage du niveau de Fermi au niveau du polaron des électrons Fe 3+ /Fe 2+ .

	E (eV)	CuFeO 2	ITO		E (eV) vs RHE	CuFeO 2		E (eV) vs RHE
										CBM
	E	vac			0.58			0			H /H + 2	0 Ni /Ni(s) (pH 9.2)
	E		CB	5.2			4.3
	E	F VB E	1.5 0.18 0.32	0.50	2.6	2.8	-1	VBM F E	0.18 0.32	O /H 2 2 O	Cu /Cu(s) (pH 7) Fe /Fe (pH 7) 3+ 2+ -1 + PtCl /PtCl (pH 3) 2-6 2-4	Ag /Ag(s) (pH 7) + 3+ 2+ Fe /Fe polaron level (pH 7)

d . Depletion region widths between 10

nm

+ 2 H + + 2 e - IO -

Acknowledgements

which crystalline facets can be associated with the exposed faces of the hexagonal CuFeO 2 nanoplatelets. Considering that the hexagonal nanoplatelets belong to the hexagonal 2H-CuFeO 2 crystalline phase, the basal planes of the hexagonal nanoplatelets can be confidently indexed as the {0001} facets. The indexing of the side faces of the nanoplatelets is more complicated because the sides of the nanoplatelets appear to be rounded, as can be seen in Figure 6.8a. In Figure 6.8b the most common facets of the hexagonal crystal family are shown. Depending on the relative size of the facets, particles, belonging to the hexagonal crystal family, exhibit different morphologies including rods, 17 prisms, 18 bipyramids, 19 truncated bipyramids 20 and discs/platelets. 11,21 The hexagonal CuFeO 2 nanoplatelets synthesized in this work most likely don't have a prism morphology, because the sides would then have to be flat and correspond solely to {10-10} facets.

Instead, the hexagonal truncated bipyramidal morphology (Figure 6.8c) would be more likely, where the nanoplatelet sides certainly consist of {10-11} facets and possibly {10-10} facets. The presence and relative size of {10-10} facets could, however, not be proven and, possibly, the relative size of {10-10} facets varies within the batch of CuFeO 2 powder. Therefore, no specific index is associated with the sides of the hexagonal CuFeO 2 nanoplatelets in the remainder of the work.

XPS was performed on the CuFeO 2 batch to figure out whether the main oxidation state of the copper and iron atoms was actually Cu + and Fe 3+ , respectively (Figure 6.9). The oxidation state of copper cannot be solely determined from the Cu2p 3/2 core level spectrum since the Cu2p 3/2 binding energy for the different oxidation states is quite similar and the whole Cu2p 3/2 spectrum for metallic copper and copper(I) oxide is indistinguishable. 22 However, Cu 2+ does not seem to be present in large amounts since that would have given rise to a strong satellite signal in the range of 940 -945 eV. To distinguish metallic copper from copper(I) oxide, the CuLMM Auger spectrum should be inspected, whereby the CuLMM binding energy here is found at 569.8 eV, typical for Cu + . 22 No signal is present at a binding energy of 568.0 eV, the CuLMM binding energy characteristic for metallic copper, thus excluding the presence of metallic copper. Determining the oxidation state of iron from the main Fe2p Nous pensons que les chercheurs devraient approfondir ces découvertes remarquables plus en détail. Il conviendrait également de vérifier que les trous à haute énergie générés lors des mesures de spectroscopie de photoélectrons ne déclenchent pas l'oxydation des groupements hydroxyles dans l'échantillon de BiVO 4 exposé à l'eau.