Michele San Martin

Josiane Viboud

Virginie Charrière

Matthieu Jouvert

Mikhail Anikin

Odette Chaix

Hervé Roussel

Stephane Brochen

David Rojas-Munoz

Daniel Bellet

Thomas Coussuet

Viet Nguyen

Sara Aghazadehchors

Bruno Baert

Emille Fourneau

Jonathan Avila Osses

Keywords: aerosol assisted metal organic chemical vapor deposition AFM, atomic force microscopy ALD, Atomic Layer Deposition AZO, aluminum doped zinc oxide CB, conduction band CBM, conduction band minimum CIGS, copper indium gallium diselenide CuCrO2, Copper-Chromium Oxide Magnesium doped Cuprous Oxide CVD, chemical vapor deposition dye-sensitized solar cells EQE, external quantum efficiency EXAFS, Extended X-ray absorption fine structure Spatial Atomic Layer Deposition SEM, scanning electron microscopy TCM, transparent conductive materials TCO, transparent conductive oxides TFT, thin film transistors TEM, transmission electron microscopy UV, ultraviolet ix valence band VBM, valence band maximum XRD, X-ray diffraction XPS, X-ray photoelectron spectroscopy ZnO, Zinc Oxide x semi-empirical dielectric constant 𝑑, distance between crystallographic planes

Firstly, I would like to

Dans le travail sur Cu2O, nous avons réalisé l'incorporation de magnésium jusqu'à 17% dans des films minces par dépôt chimique en phase vapeur assisté par aérosol, entraînant des changements de morphologie. La résistivité électrique a été réduite jusqu'à des valeurs de 6,6 ohm.cm, en raison de l'augmentation de la densité de porteur de-charges jusqu'à 10 18 cm -3 . L'incorporation du magnésium a en outre eu un impact sur la stabilité de la phase Cu2O. En effet la transformation du Cu2O en CuO en conditions oxydantes est considérablement retardée par la présence de Mg dans les films, en raison de l'inhibition de la formation d'un type particulier de lacune de cuivre (split vacancy). L'intégration dans les jonctions pn a été réalisée avec succès en utilisant uniquement des voies de dépôt chimique en phase vapeur, en combinaison avec le ZnO de type n. Néanmoins, l'application de Cu2O dopé au Mg dans les cellules solaires présente un effet photovoltaïc très faible, loin des meilleures valeurs de l'état de l'art.

Dans le travail sur CuCrO2, nous démontrons la première fabrication d'hétérostructures de nanofils en configuration coeur/coquille ZnO/CuCrO2 utilisant des techniques de dépôt chimique adaptées pour des grandes surface, à faible coût, facilement implémentées à des températures modérées et leur intégration dans des photodétecteurs UV auto-alimentés. Une coquille conforme de CuCrO2 avec la phase de delafossite et avec une uniformité élevée a été élaborée par un dépôt chimique en phase vapeur assisté par aérosol sur un réseau de nanofils ZnO alignés verticalement, obtenu par dépôt par bain chimique. Les hétérostructures ZnO/CuCrO2 coeur-coquille présentent un comportement rectificatif significatif, avec un ratio de rectification maximal de 5500 à ± 1V, ce qui est bien meilleur que les dispositifs 2D similaires rapportés dans la littérature, ainsi qu'une absorption élevée supérieure à 85% dans la région UV. Lorsqu'ils sont appliqués en tant que photodétecteurs UV auto-alimentés, les hétérojonctions optimisées présentent une réponse maximale de 187 μA / W sous une polarisation nulle et à 374 nm ainsi qu'une sélectivité élevée avec un ratio de rejet entre l'UV-et le visible (374-550 nm) de 68 sous irradiance de 100 mW/cm 2 The main body of the thesis, composed by the following three chapters, explores the materials description and their application into functional devices. We start with chapter III where we present the preliminary tests on the Cu2O thin films doped with different cations, with a particular interest to Sn-doped Cu2O.

Chapter IV is based on Mg-doped Cu2O study with an extensive analysis on the crystallographic structure, microstructural topography, and optical and electrical response to understand the dopant effect. In depth study on the thermal stability is also reported. The chapter ends with a study on pn junction between Mg-doped Cu2O and ZnO, first with a numerical simulation and then by experimental deposition of these structures, and finally a solar cells performance results.

Chapter V starts with an optimization of CuCrO2 delafossite phase deposition by aerosol-assisted MOCVD, which is latter applied to a nanostructured UV photodetector based on ZnO and CuCrO2.

The last chapter serves as a conclusion of the work developed. It contains the main highlights in terms of results, the new scientific ideas suggested, the problems or open questions and finally, we suggest the future work to be developed.

The thesis' appendix presents the fitting of EXAFS data from Cu2O and Cu2O:Mg thin films. As supplemental materials, XRD PDF files and Raman reference files of the different oxides are also present in the appendix Chapter I: Introduction

Cuprous oxide (Cu 2 O)

The exploration of novel oxide semiconductors has been intensified in the last two decades, mostly driven by transparent electronics or solar cells applications. In the p-type semiconducting compounds group, cuprous oxide, Cu2O, presents promising electrical and manufacturing features that identify it as a suitable candidate for the electronics and energy devices 1 .

In the family of the oxide materials, the cuprous oxide has a crystalline cubic structure with a lattice parameter of 4.2696 Å. The crystal structure corresponds to the space group 𝑃𝑛3 ̅ 𝑚, which incorporate a full octahedral symmetry. In the case of considering the oxygen atoms as the origin of the body centred cubic (bcc) primitive cell, the 4 copper atoms are located in the diagonals in the positions: ( The oxygen atoms are tetrahedrally coordinated by the copper, while the copper ones are linearly coordinated by the oxygen ones. The distance between Cu-O neighbouring atoms is 1.85 Å, O-O is 3.68 Å and Cu-Cu is 3.02 Å. This structure creates an oxide with a density of 6.10 g.cm -3 , and a molar mass of 143.09 g.mol -1 . In stoichiometry conditions, the concentration of copper is twice the concentration of oxygen, 5.05x10 22 cm -3 and 2.52

x 10 22 cm -3 , respectively.

The Cu2O structure has another relevant characteristic, its dichotomy. There are two identical and interpenetrating networks of Cu and O atoms inside the crystal, which do not have direct chemical bond between them. Both network have an anti-SiO2 structure that are stable due to Van-der-Waals forces 3 . This is visible in the Figure I-1 b) by the two different tones of grey in each network.

The band gap of the cuprous oxide is 2.17 eV at 4 K, according to experimental results 4 . This is a direct gap at the centred of the Brillion zone and it is formed by the copper 3d 10 and 4s 0 orbitals which correspond to the valence and conduction band, respectively, both with the same parity 5 At room temperature, the predicted band gap is 2.096 eV 5 . On the effective mass on both bands, the curvature of the bands predicts a higher effective mass for electrons in the conduction band than for the holes in the valence band. Experimental values confirm this fact, presenting 0.99 of effective mass for electrons and 0.58 for holes 7 .

Conductivity in Cu 2 O

The p-type behaviour of cuprous oxide arises from the special configuration of the valence band, when compared to other oxides. In p-type transparent conductive oxides (TCO), the large electronegativity of oxygen creates a strong localized valence band edge formed by the 2p levels. Consequently, the hole effective mass is relatively high and the mobility of holes is weak. However, in Cu2O the top of the valence band is formed by 3d 10 levels of the Cu + cation, as visible in the Figure I-3. The existence of this different level contributes to a less localization of the holes, which improves the mobility of these charges.

Due to this fact, in Cu2O mobility values can reach up to 100 cm 2 .V -1 .s -1 in single-crystals 8 , making this material on of the few oxides with high p-type mobility. 10 and O 2p 6 are assumed to be equivalent Reproduced from H. Kawazoe et al 9 The conductivity in cuprous oxide appears by polaronic hoping, contrarily to classical semiconductors. In this mechanism and considering the p-type behaviour of Cu2O, the conduction of holes through the material distorts the crystal lattice, causing a holephonon coupling. Due to the interaction of electronic charges and phonons, traps appear that localize the hole. Then, the hole migrates to another site by thermal motion, via hopping mechanism 10 .

As an oxide with a band gap of 2.17eV, Cu2O is an insulator in a stoichiometric composition. Nevertheless, the origin of the Cu2O p-type conductivity surfaces from intrinsic defects present in the material 1 . As confirmed by experimental reports, the defects can vary in origin, from cation deficiency, interstitial or vacant oxygen, Schottky barrier or simply electron-hole defects 11,5 . Other types of defects as interstitial copper, Frenkel and anti-Frenkel defects seem to be highly unlikely to occur, based on theoretical predictions 3 .

Here below, in Table I-1, we present the most common types of intrinsic defects that can be present in Cu2O 12 . For the generation of holes, the copper vacancies, 𝑉 𝐶𝑢 ′ , are considered the most favourable defect, when compared to interstitial oxygen defects, 𝑂 𝑖 ′′ . These copper vacancies can be formed by oxidation of the Cu2O, which leads to the removal of a copper atom, leaving a negatively changed vacancy. As consequence, a hole is introduced in a valence band, creating a acceptor level between 0.3 -0.6 eV above this band 14,16 . The intrinsic p-type semiconductor behaviour appears from these stable copper vacancies, which can be high as 10 20 cm -3 in concentration, considering 1 to 3% of all copper atoms removed 13 . However, since the ionization of these particular defects is incomplete, the free holes concentration only reaches 10 18 cm -3 at room temperature 5 .

Still, the nature of these level is yet to be completely understand. Indeed, two possible copper vacancies can appear in Cu2O: the simple copper vacancy 𝑉 𝐶𝑢 ′ related to the removal of one copper atom, leaving two oxygen atoms with three copper neighbours, or the split copper vacancy 𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 in which the copper disappearance is followed by a neighbouring copper atom movement towards the vacancy. In the latter case, the copper atom moves into a tetrahedral site with four neighbouring oxygen atoms. This is coordination is similar to one in CuO phase. The split vacancy also promotes the change of the oxidation state of the dislocated copper from 1 + to 2 + 13 . A representation of both defects in visible in the Figure I-4. The formation energies of both copper vacancies are a controversy topic in the literature with different theoretical studies presenting distinct values. In the majority of the cases, simple vacancies are more stable than split vacancies, as reported by Raebiger et al 17 , Scanlon et al. 14 , Nolan et al. 13 , Soon et al. 18 , and Isseroff and Carter 16 . Nevertheless the formation energies are quite distinct: in Nolan's study, both vacancies energy formation are around 0.4eV, while in Scanlon case 𝑉 𝐶𝑢 is around 0.41eV and 𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 is 1.24eV. Soon et al. present 0.47 eV and 0.78 eV for 𝑉 𝐶𝑢 and 𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 , respectively. Isseroff and Carter 16 present different exchange-correlation functional methods used with a consistent difference between the two vacancies of 0.21 ± 0.03 eV, where the simple copper vacancy is more stable. In the position of these acceptors levels regarding the valance band, Nolan et al reports 0.47 eV for 𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 and 0.23eV for 𝑉 𝐶𝑢 13 .

Different studies report an increase of copper vacancies by post-deposition annealing treatments under oxidizing conditions 19,20,21 . At temperatures below 300ºC, the Cu2O thin films show a decrease of the resistivity due to the generation of holes 19 , as well as an increase of holes mobility 20 . Moreover, the transmittance can be further increased in these thermal treatments by a partial removal of defect band tail which enlarger the optical band gap 20 . At temperatures higher than 300 ºC, CuO parasitic phase starts to be formed, degrading the Cu2O optoelectronic properties 21 . Cuprous oxide thin films can also be annealed at higher temperatures, up to 700 ºC, in reducing atmospheres as vacuum 22 . In this case, even if the resistivity decreases due to a lower carrier density, the mobility of the thin films is improved to values higher than 10 cm 2 .V -1 .s -1 .

Cu 2 O deposition

The interest in Cu2O has generated a large number of studies to grow thin films using a large variety of techniques, from chemical to physical, under vacuum or at atmospheric pressure. In the Figure I-5, the number of publications on Cu2O obtained from Scopus are plotted regarding the year, from 2000 to 2017. An increase of "Cu2O" and "doped Cu2O" reports are clear in the last 17 years. The expression "copper oxide" shows a larger number of publications since it is a general term, including also CuO and Cu4O3.

A list of different studies developed in the last decade on intrinsic Cu2O are presented in Table I-2, with some important parameters highlights. These different studies show the versatility of Cu2O to be deposited from RT to 350 ºC using physical and chemical techniques. In terms of electrical properties, the resistivity values are in most of the cases close to 100 Ω.cm, with a decrease to 7.3 Ω.cm with air annealing, which leads to the formation of CuO at high temperatures 20 . The presence of oxygen seems to be one of main issues in the preparation of these films due to the formation of CuO parasitic phase. In the Figure I-6, the phase diagram of Cu-O for single crystals is represented. As one can observe, the Cu2O is unstable at room temperature under atmospheric pressure, 736 Torr. Thus, when one deposits thin films of Cu2O, the complete prevention of CuO phase is nearly impossible. This main issue can have large impacts on electric and optical properties of the films, since both materials present different band-gap and point defects. The appearance of an intermediate phase between CuO and Cu2O is also reported in the literature 30 , paramelaconite, Cu4O3, formed by stacks of two-fold coordinated copper atoms and four-fold coordinated copper atoms. Nevertheless, studies on the experimental deposition of Cu4O3 thin films are still lacking, as depositions by magnetron sputtering show a narrow oxygen flow rate process window 31,32,33 . Additionally, further optical and electrical properties measurements are required to access the possible applications of this material.

Dopants

The use of extrinsic elements on Cu2O has been widely tested by the literature in order to improve electric, optical and magnetic properties. In one first group of elements formed by non-metals, Cl -and N -3 are atoms that created a significant impact on the properties of Cu2O thin films. Those two elements can substitute the oxygen site and act as a donor for the Cl -, and as an acceptor for the N -3 . In the Cl -case, it was possible to obtain a n-type Cu2O, confirmed by photocurrent-potential measurements and presenting a low resistivity of 7 Ωcm 34 . Concerning acceptors, N doping seems to have also an effect on electrical properties with a reduction of resistivity down to 15 Ω.cm, and significantly high values of mobility, 20cm 2 .V -1 .s -1 35 .

Additionally, metals have also been tested as dopants for Cu2O, experimentally and in models. In the group of transition metals, there are different tested elements as: Cd 36 , Co 37 , Mn 38 , Ni 39 Ag 40 , Fe 40 , Cr 40 , Al 37 , Zn 37,41 and V 37 for different applications as ferromagnetic oxide, conductivity improvements or photocatalytic properties.

Na doping of Cu2O has been the most successful of doping cases with large enhancements in conductivity. In a study by Tadatsugu Minami et al. 42 resistivity values were decrease to 3.95 10 -2 Ω.cm in Cu2O films obtained by thermal oxidation of Cu sheets.

Mobility and charge carrier's values were also improved up to 100 cm 2 .V -1 .s -1 and 10 19 cm - 3 , respectively. In the same study, this doped system was implemented for photovoltaic applications as p-type absorber layer.

In 2008, based on first-principle calculations, Nolan et al. 2 suggested the doping of cuprous oxide with larger cations than Cu + in order to increase the band gap, while maintaining the cubic structure. These cations, such as Mg 2+ , Sn 2+ , Sr 2+ or Ca 2+ would distort the crystallographic lattice and consequently diminish three-dimensional Cu-Cu interactions, leading ultimately to a band gap increase. The p-type conductivity would appear consequently to the creation of a double copper vacancy. Since the dopants are divalent cations, one primary copper vacancy would compensate the dopant presence, while a secondary copper vacancy supplies an extra hole 2 . Additionally, the use of "electronically inert" dopants would avoid the hybridization of the valence band and conduction band edge states or the introduction of in-gap state 2 .

In a previous work, the incorporation of Sr had already been achieved, showing improvements in the conductivity of the films, however, without any change in the optical properties of the material 43 . Thin films of Cu2O:Sr with different strontium concentrations were deposited by MOCVD. As main results, Sr can be incorporated up to 14.9% with the presence of a parasitic SrCO3 phase for films with more than 5% of Sr. The major impacts were on the film morphology and resistivity, both represented in al. 43

CuCrO 2 -Delafossite

The Cu2O band gap of 2.17 eV is small for transparent electronics applications, since the transmittance of Cu2O films is low in the visible part of the light spectrum 1 .

Suggested by Kawazoe 44 , changing the crystallographic structure by introducing other cations in Cu2O would increase of the band gap and improve both transparency and electrical properties.

Consequently, a new family of TCO materials emerged, with a delafossite crystal structure, such as CuAlO2 44 , CuGaO2 [START_REF] Izaki | Electrochemically Constructed P-Cu2O/n-ZnO Heterojunction Diode for Photovoltaic Device[END_REF] , CuFeO2 46 , CuCrO2 47 , CuYO2 48 , or CuScO2 49 .

Additionally, other non-delafossite phases also appeared as possible materials for p-type TCO, such as Cu2SrO2 50 or (LaO)CuS 51 . However, none of these materials combined high conductivity, high transparency and low synthesis temperatures, as already achieved by ntype TCOs 52 . Among these two families of materials, CuAlO2 compound shows visible transparency around 70% for a 230 nm thick film, with a direct band gap of 3.5eV, yet with relatively high resistivity values up to 3 .cm [START_REF] Yanagi | Electronic Structure and Optoelectronic Properties of Transparent P-Type Conducting CuAlO2[END_REF] . In this cases, the temperatures of deposition were higher than 600 ºC, which is impractical for the majority of transparent substrates.

Other particular case is CuCrO2 that has received increasing interest as a promising p-type transparent semiconductor [START_REF] Nagarajan | P -Type Conductivity in CuCr1-xMgxO2 Films and Powders[END_REF] , owing to its wide direct band gap energy of 2.8 eV combined with its high conductivity, which is typically larger than 200 S.cm -1 when doping with Mg [START_REF] Nagarajan | P -Type Conductivity in CuCr1-xMgxO2 Films and Powders[END_REF][START_REF] Robertson | Limits to Doping in Oxides[END_REF][START_REF] Tripathi | Atomic Layer Deposition of Transparent Semiconducting Oxide CuCrO2 Thin Films[END_REF] . In comparison with other Cu-based delafossite phases, CuCrO2 is more conductive mainly due to the higher delocalization of holes produced by Cu vacancies. This is caused by a mixing of Cr d states with O 2p states in the valence band, which creates shallow energy transition levels for the holes [START_REF] Scanlon | Understanding the P-Type Conduction Properties of the Transparent Conducting Oxide CuBO2: A Density Functional Theoryanalysis[END_REF][START_REF] Scanlon | Understanding the P-Type Defect Chemistry of CuCrO2[END_REF] . In the last twenty years, CuCrO2 thin films have been deposited by a wide variety of physical and chemical deposition techniques, such as radio-frequency (RF) sputtering [START_REF] Nagarajan | P -Type Conductivity in CuCr1-xMgxO2 Films and Powders[END_REF][START_REF] Nagarajan | P-Type Conductivity in the Delafossite Structure[END_REF][START_REF] Yu | Applied Surface Science Characteristics of P-Type Transparent Conductive CuCrO2 Thin Films[END_REF] , pulsed-injection metal-organic chemical vapour deposition (MOCVD) [START_REF] Crêpellière | Transparent Conductive CuCrO2 Thin Films Deposited by Pulsed Injection Metal Organic Chemical Vapor Deposition: Up-Scalable Process Technology for an Improved Transparency/conductivity Trade-Off[END_REF] , aerosol-assisted CVD [START_REF] Sánchez-Alarcón | Ultrasonic Spray-Pyrolyzed CuCrO2 Thin Films[END_REF] , spray pyrolysis [START_REF] Farrell | Synthesis of Nanocrystalline Cu Deficient CuCrO2 a High Figure of Merit P-Type Transparent Semiconductor[END_REF] , sol-gel process [START_REF] Götzendörfer | Preparation of P-Type Conducting Transparent CuCrO2 and CuAl0.5Cr0.5O2 Thin Films by Sol-gel Processing[END_REF] , and atmospheric pressure plasma torch [START_REF] Chen | Characterization of Delafossite-CuCrO2 Thin Films Prepared by Post-Annealing Using an Atmospheric Pressure Plasma Torch[END_REF] . In particular, Farrell et al. [START_REF] Farrell | Synthesis of Nanocrystalline Cu Deficient CuCrO2 a High Figure of Merit P-Type Transparent Semiconductor[END_REF] reported the growth of undoped CuCrO2 thin films by spray pyrolysis with a high conductivity of 12 S.cm -1 combined with an optical transmittance higher than 55%. Interestingly, metal organic precursors that are compatible with aerosol-assisted CVD were used and the nanocrystallization process of the delafossite phase typically proceeds at temperatures below 400 °C, without any post-annealing treatments. Similar results were also reported by Sánchez-Alarcón et al. [START_REF] Sánchez-Alarcón | Ultrasonic Spray-Pyrolyzed CuCrO2 Thin Films[END_REF] In another recent study on magnesium-doped CuCrO2 [START_REF] Tripathi | Enhanced P-Type Transparent Semiconducting Characteristics for ALD-Grown Mg-Substituted CuCrO2 Thin Films[END_REF] , conductivity values achieved a record high of 217 S.cm -1 , with transmittance values of 70%, using a low deposition temperature of 250 ºC by Atomic Layer Deposition (ALD).

On physical techniques, the work from Barnabé et al. [START_REF] Barnabe | P-Type Conducting Transparent Characteristics of Delafossite Mg-Doped CuCrO2 Thin Films Prepared by RF-Sputtering[END_REF] using RF sputtering enable the increase of visible transparency, with a direct optical bandgap of 3.3 eV after annealing at 600°C. This finding in combination with an effective carrier concentration of about 10 21 cm -3 , lead to an high figure of merit of CuCrO2 thin films, up to 1.5 x10 -7 Ω -1 .

Additionally, CuCrO2 nanoparticles have been prepared by hydrothermal synthesis [START_REF] Xiong | Hydrothermal Synthesis of Ultrasmall CuCrO2 Nanocrystal Alternatives to NiO Nanoparticles in Efficient P-Type Dye-Sensitized Solar Cells[END_REF] . In terms of applications, CuCrO2 thin films have until now been integrated into amorphous thin film transistors (TFT) in electronics [START_REF] Sanal | Room Temperature Deposited P-Channel Amorphous Cu1-xCrxO2-δ Thin Film Transistors[END_REF] , ozone sensors [START_REF] Deng | Room Temperature Ozone Sensing Properties of P-Type Transparent Oxide CuCrO2[END_REF] , photocatalytic H2 production devices [START_REF] Ma | Photoelectrochemical Properties of CuCrO2: Characterization of Light Absorption and Photocatalytic H2 Production Performance[END_REF] and biological devices as antibacterial agent against E coli bacteria 72 .

Devices based on oxide semiconductors

Transparent electronics

The development of transparent electronics has increased drastically in the last years, being considered one of the trending research topics for new advanced materials [START_REF] Fortunato | Oxide Semiconductor Thin-Film Transistors: A Review of Recent Advances[END_REF] .

The possibility of creating transparent devices enables a new range of applications that starts in transparent displays for current everyday devices, such as phones, tablets and computers, but furthermore it can possibly reach smart surfaces as windows, tables or street furniture based on glass or other transparent substrates [START_REF] Fortunato | Oxide Semiconductor Thin-Film Transistors: A Review of Recent Advances[END_REF] . [START_REF]COST Action MP1308 -Towards Oxide-Based Electronics[END_REF][START_REF]Project ORAMA -Oxide Materials for the Post-Silicon Electronics Era[END_REF] .

Even though the first documented study on transparent conductive oxides (TCO)

was reported in 1907 with the production of CdO thin film [START_REF] Bädeker | Über Die Elektrische Leitfähigkeit Und Die Thermoelektrische Kraft Einiger Schwermetallverbindungen[END_REF] , just on the transition to the 21 th century the transparent electronic research gained a relevant importance. The success of n-type semiconductor oxides, as Sn-doped In2O3 (indium tin oxide, ITO), Al-doped ZnO, and Sb-doped SnO2, enable the implementation of these materials as transparent conductive oxides [START_REF] Ohta | Transparent Oxide Optoelectronics[END_REF] . The high transparency combined with a metallic-like conductivity lead to the application in devices as organic light-emitting diodes (OLEDs), liquid crystal displays (LCDs), or 3 rd generation solar cells [START_REF] Ohta | Transparent Oxide Optoelectronics[END_REF] . Furthermore, the semiconductor behaviour of these n-type materials allowed the production of transparent thin film transistors (TFTs), due to controllable electrical conductivity. Since 2003, when the first transparent TFT based on ZnO was discovered [START_REF] Hoffman | ZnO-Based Transparent Thin-Film Transistors[END_REF] , the developments lead to the integration in active matrix for organic light emitting diodes (AMOLED) technology [START_REF] Fortunato | Oxide Semiconductor Thin-Film Transistors: A Review of Recent Advances[END_REF] . Consequently, ultra-thin screens appeared in the last decade, allowing a semi-transparent device that consume 90% less of energy when compared to LCD and can ultimately use ambient light to increase its energy efficiency.

As the n-type oxide semiconductors materials are already a well-established technology, the interest in p-type oxides is increasing significantly, in order to produce complementary metal oxide semiconductors (CMOS), enabling the application of extremely compact circuits with smaller power consumption needs, and even other type of electronic devices based on p-n junctions 1 . A p-type oxide material with a semiconductor high performance would improve as well the anode part of an organic light emitting device (OLED), because the drain current in the saturation mode would not be influenced by the TFT hole current supplies.

Nevertheless, the performance of these type oxides is fairly poor when compared with the n-type materials. The main reason of this difference lies in the low mobility of the charge carriers, holes, as a result of the lower mobility of the derived carries in the valence band when compared with the derived carriers of the conduction band 1 .

Thin Film Transistor

A fundamental component of the transparent electronics is the thin film transistor (TFT), a type of field-effect transistors (FET), composed of stacked layers of different materials thin films [START_REF] Aber | Thin Film Transistors & Flexible Displays[END_REF] . As it is a FET type, the TFT is composed of three different terminals: The working principle of the device is based on the controlled current flow between the source and drain, by the application of a bias voltage at the gate. In the off state of the TFT, the passage of current through the semiconductor body is difficult due to the low conductivity of this layer. When a positive bias is applied on the gate, the electric field created allows the flux of electrons from the source to the drain, due to the appearance of a negatively charge region close to the isolate layer. This field-effect is the mechanism responsible for the current flow, corresponding to the on-state of the transistor.

Cu 2 O based TFT

As copper oxide is one of the few oxides with high p-type mobility values, TFTs were tested using Cu2O as a channel semiconductor. The first publication reporting the production of a TFT based on Cu2O was achieved in 2008 by Hosono et al. 8 , where a single crystal of Cu2O was epitaxially grown on MgO at 700 °C The hole mobility presented a high value for Cu2O, 90 cm 2 V -1 s -1 . However, even if the TFT shown a p-channel operation although, the on-off ratio and the effective mobility were only 6 and 0.26 cm 2 .V -1 .s -1 respectively 8 . The transfer characteristics of the TFT is visible in Figure I-10a). In other follow-up research from 2010 [START_REF] Zou | Top-Gate Low-Threshold Voltage P-Cu2O Thin-Film Transistor Grown on SiO2/Si Substrate Using a High-κ HfON Gate Dielectric[END_REF] , the results were substantially improved for Cu2O deposited by pulsed laser deposition in pure O2 with a P(O2) = 0.6 Pa. The TFTs present a low threshold voltage of -0.8 V, an on-off current ratio of 3x10 6 , a saturation mobility of 4.3 cm 2 /V s, and a sub threshold swing of 0.18 V/decade, for films deposited at 500 °C.

The temperature increase during the deposition increase the concentration of Cu 2+ cations, which consequently result in the formation of the CuO phase. All the electric properties and transistor functions were diminished with the CuO phase increase, as visible in the

Figure I-10b).
Other study focuses on the annealing of Cu2O, at lower temperatures from 150 °C to 500°C for 7 minutes under a base pressure of 2x10 -6 Torr [START_REF] Sohn | Effects of Vacuum Annealing on the Optical and Electrical Properties of P-Type Copper-Oxide Thin-Film Transistors[END_REF] . Holes mobility values were 47.5 cm 2 V -1 .s -1 with a low concentration of charge carriers, 2.95x10 14 cm -3 and good transmittance of 55%, in the visible region. The morphology of the films is visible in the Thermal oxidation of Cu metallic films were used to prepare a p-type channel TFTs [START_REF] Figueiredo | P-Type CuO Thin-Film Transistors Produced by Thermal Oxidation[END_REF] . The treatment to obtain the Cu2O phase was performed at 200 °C in an oxygen rich atmosphere. The final device show a mobility value of 2.2 cm 2 V -1 s -1 for a thickness of 112 nm, and an on-off ratio of 60. In this study, the oxidation at higher temperature, 250 °C, origins the appearance of CuO phase. The same group, Fortunato et al [START_REF] Fortunato | Thin-Film Transistors Based on P-Type Cu2O Thin Films Produced at Room Temperature[END_REF] , also studied Cu2O thin films by reactive RF magnetron sputtering deposited at room temperature. The TFT show a field-effect mobility of 3.9 cm 2 V -1 s -1 and an on/off ratio of 200. In Figure I-11a) the transfer curve is plotted.

The first solution-based p-type transistor was achieved in 2013 by sol-gel using spin coating, followed by two annealing stages [START_REF] Kim | P -Channel Oxide Thin Film Transistors Using Solution-Processed Copper Oxide[END_REF] . During the first annealing in N2, a thin film of metallic Cu was formed, and then oxidized to Cu2O by the second annealing in O2. In this device, a CuO layer functioned as an insulator, resulting in a limited on-off ratio of 100, while the mobility presented to be 2.2 cm 2 V -1 s -1 , for films annealed at P(O2)= 0.04Torr.

Other possible configurations is a bilayer TFT, composed of two films, enabling the production of tuneable devices at low processing temperatures (170 °C) 85

Photovoltaics

This following chapter's part is based on a previously written chapter for the book "Materials for Photovoltaic Solar Cells: Materials for Sustainable Energy Applications:

Conversion, Storage, Transmission, and Consumption", (2017). It was written in collaboration between: Munoz-Rojas, D., Liu, H., Resende, J., Bellet, D., Deschanvres, J.

L., Consonni, V. and Zhang, S.

Recently, the scientific society has been concerned about the need for more energy sources. The energy dependence problem has stimulated the search for solutions and this has resulted in a significant boost to solar energy due to its future possibilities. So far, the silicon-based method has been the most common way of producing solar panels.

Nevertheless, new other materials have been tested to enhance the photovoltaic capacity to produce energy.

The interest on photovoltaics lead to new scientific projects supported by the Horizon 2020 initiative on low-cost and highly efficient solar cells. Two of these projects are specific for advanced technologies using new materials as: Cheops, for the development of the recent perovskites solar cells 86 , or CPVMatch for the research on multi-junction solar cell architectures aiming for practical performance of high concentrating photovoltaics 87 .

Solar Cells

A solar cell is a device capable of transforming light into electricity. Absorption of a photon causes the excitation of an electron within the absorbing material. In order to work, i.e. to collect the electrons to flow out of the cell generating a current and voltage, an asymmetry must be present in the device. This asymmetry is for instance created by having a pn junction, formed by two semiconductors with opposite electrical behaviour. As it is known, in p-type semiconductor material, the majority carriers are holes, while in ntype are electrons. In the case of inorganic semiconductors (silicon for example), when ntype and p-type are in contact so that an interface is created, the excess electrons in the ntype material will diffuse into p-type material, and the other way around for holes from the p-type into the n-type. The consequence of this spontaneous flow is the spontaneous generation of an electrical field near the interface, which opposes to a further diffusion of electrons into p-type, and holes into n-type material, until thermal equilibrium is reached To mathematically describe the depletion zone, charge neutrality and Poisson equation are used 89 . For simplicity, the discussion is limited to a 1D model. Thus, equations are given as follow:

- 𝑑 2 𝜓 𝑑𝑥 2 = 𝑑𝐸 𝑑𝑥 = 𝜌(𝑥) 𝜖 𝑠 𝑊 𝐷𝑝 𝑁 𝐴 = 𝑊 𝐷𝑛 𝑁 𝐷 (I.1)
ψ is the potential, E is the electric field, 𝜌 is the charge density, x is the distance from the junction and s is the dielectric constant. To simplify the model further, we assume an abrupt junction where ND, NA are both constant. A boundary condition assumes as well that at each end of the depletion zone there is no electrical field. Thus, the electrical profile can be solved by equations (I. Thus, the total potential across the device, also called built in potential ψbi, would be the sum of ψp and ψn. qVbi = ψbi= ψp+ψn. With the charge neutrality equation, the depletion width on each side can be determined.

In solar cells, when one or the two semiconductors of the p-junction has a gap suitable for visible light absorption, the illumination of the pn junction creates electronhole pairs in the depletion zone, a phenomenon called photovoltaic effect. Due to the builtin electrical field in this region, the holes will move to the p-type semiconductor, while the electrons move to the n-type material. If the circuit is open, the positive charges will accumulate in the p-type material and the negative ones in the n-type material, reducing the electric field in the depletion zone. However, if we introduce an external load in the system (i.e. the circuit is closed), the charges can move outside of the junction, generating an electrical current and increasing once again the electrical field in the depletion zone.

Consequently, the energy produced in the solar cell is dependent on the rate of electron-hole generation, bias voltage and all the conduction processes inside and at the borders of the cell [START_REF] Möller | Semiconductors for Solar Cells[END_REF] . In The performance of a photovoltaic solar cell can be evaluated using a variety of parameters [START_REF] Boyle | Renewable Energy: Power for a Sustainable Future[END_REF] . One of the first and most important types of analysis is the Power Conversion Efficiency (PCE), defined as the ratio of the power produced by the cell (Pprod) with regards to the power from the incident light on the cell (Pinc). It is given by:

𝜂 = 𝑃 𝑝𝑟𝑜𝑑 𝑃 𝑖𝑛𝑐 (I.6)
The values for the incident light power are usually standardized for light sources with an AM 1.5G spectrum and a light flux of 100 mW.cm -2 . This solar spectrum is a standard value taking into account for example the absorption by H2O or CO2 molecules in the atmosphere.

The produced power is the result of multiplying the obtained current for a particular applied voltage. For the calculation of the PCE, the maximum power is considered, as indicated in equation (I.7).

𝑃 𝑝𝑟𝑜𝑑 = 𝐼 𝑖 𝑉 𝑖 (I.7)
where Ii and Vi are the electric current and potential values that give maximum power. Thus, IV curves are measured to estimate the maximum power as well as other parameters. The IV curves are obtained by measuring the electrical current dependence on the bias voltage applied to the terminals of the cell, under illumination. When the applied voltage is zero, the current measured, ISC, corresponds to the short circuit current. This current is described as the electric flow produced by the illumination of the cell (photocurrent) with no external loads, which strongly depends on the number of photons absorbed and the efficiency with which carriers are collected.

When the bias voltage increases, the electrical current drops, due to the increase of where FF is defined as fill factor.

𝐹𝐹 = 𝑃 𝑚𝑎𝑥

𝐼 𝑠𝑐 𝑉 𝑜𝑐 (I.9)

In Figure I-15b) the power produced by the solar cell is directly related to the area of the square whose edges are equal to Vmax and Imax. Therefore, the fill factor can then be graphically interpreted as the ratio between the maximum power obtained from the cell and the product of the short circuit current and open circuit voltage (power that would be obtained from an ideal pn junction, i.e. diode).

The shape of the IV curve, which can be depicted in the FF value, is strongly dependent on the different resistances of the solar cell. The shunt resistance, RSh, is mainly responsible for the leakage current troughtout the cell. In this case, the resistance should be maximized to prevent recombination of electron-holes in the pn junction, reducing the leakage current on the cell. In the IV curve, this maximization is expressed by a smoother slope at lower bias voltages. While the series resistance, RS, is the component of resistance regarding to the electrodes contacts, bulk resistance or other interfacial barriers (i.e. series resistance). The reduction of this type of resistance leads to an increase of electrical current, which increases the slope of the IV curve at bias voltages near VOC , thus a more squared shape of the curve (i.e. closer to the ideal diode). The control of both resistances is fundamental to increase the final PCE of the cell.

In general, the equation that expresses the dependence of current and voltage can be defined as:

𝐼 = 𝐼 𝑆𝐶 -𝐼 0 (𝑒 𝑞(𝑉-𝐼𝑅 𝑆 ) 𝑛𝑘𝑇 -1) + 𝑉 -𝐼𝑅 𝑆 𝑅 𝑆ℎ (I.10)
where 𝑛 is the non-ideality factor. In practice, because measurements are made illuminating a particular cell area, short-circuit current density (JSC, i.e. ISC/illuminated area) is commonly quoted when evaluating IV curves. While the evaluation of IV curves under illumination is a basic evaluation of cell performance, dark IV curves (i.e. without any illumination) are also routinely performed. This measurement allows to evaluate the non-ideality factor of the pn junction in the absence of any light induced bias. The comparison of RS and RSh for light and dark measurements also provides useful information 92 .

Copper oxide based Solar cells

In order to have solar cells with a higher efficiency/cost ratio, the used materials should ideally be abundant and manufactured using low-cost and scalable techniques. In this context, there has been a recent drive towards the use of stable inorganic Those with better ratio than the currently used crystalline Si are located in the first quadrant 93 For application into solar cells, Cu2O is being heavily investigated since Cu is abundant, non-toxic and theoretical maximum obtainable efficiencies for Cu2O-based single junction cells are high, around 20% 94 .However, the low energy band-gap of Cu2O, 2.17eV, limits the abortion of the blue and green part of the solar spectrum, significantly limiting the potential of a single Cu2O-based solar cell. As a solution, copper oxide can be used as a PV material by combining several cells of different band gap stacked on top of each other, called a tandem cell. Each cell absorbs a different portion of the solar spectrum, with band gap decreasing in the direction of light incidence. Therefore each cell absorbs photons closer to its band gap, which also results in a lower thermalisation of the electronhole pairs generated. Tandem cells allow a theoretical efficiency limit of 42.5% for a twocell stack and 47.5% for a three-cell stack tandem cell, as opposed to the Schockley-Queisser limit of 31% for a single band gap cell 95 . For this purpose, a high efficiency Cu2O-based solar cell would be intended for tandem cell implementation.

The difficulty in obtaining n-type Cu2O is the main factor limiting the maximum efficiencies that can be obtained, since it has not been possible to make an efficient homojunction of the type n-Cu2O/p-Cu2O. Although there are some reports on n-type Cu2O fabricated [96][97][98] , the origin of the n-type character is not well established and the homojunction cells prepared with such materials presented very low efficiencies 99,100 .

This fact can be overcome by the use of a different n-type semiconductor that presents a band structure compatible with Cu2O to form a heterojunction. ZnO as an n-type semiconductor oxide can be combined with Cu2O which will function as p-type absorber material. The combination of these two oxide semiconductors has been widely investigated for a oxide based solar cell 101,102,103 . For this junction, a built-in bias of approximately 0.6 to 0.7 V is expected, which should approximate the achievable open-circuit voltage (VOC) 104 . In a Cu2O/ZnO solar cell, light is absorbed by Cu2O and the generated electrons are injected into the ZnO conduction band and extracted, thanks to the high electron mobility of ZnO, through the contact towards the external circuit, holes being transported through the Cu2O layer to the external circuits as majority carriers, and electrons towards the junction as minority carriers.

The maximum value achieved by functional devices is only 4.12% 105 for an intrinsic ZnO/Cu2O junction, which was recently increase to 8.1% in a MgF2/ZnO:Al/Zn0.38Ge0.62O/Cu2O:Na heterojunction 106 . Also, this record of efficiency was achieved by thermal oxidized Cu2O, which produce large grains, in the tens of microns range, reducing the fraction of grain boundaries and increasing the mobility. However, this technique is not suitable for industrialization, since it requires very high temperatures, 1000 ºC 107 . Thus, the values are still far from predicted efficiency, and using energy intensive methods not suitable for large scale production 108 .

There has been an intensive study on ZnO/Cu2O cells fabricated by electrodeposition. Although, the initial reports suffered from rather low VOC values, engineering of the ZnO/Cu2O and of the Cu2O/contact interface has had a huge impact on efficiency, clearly showing that in addition to the bulk properties of the different semiconductors used, the interfaces (surface chemistry and morphology) play a key role on device efficiency 109,110 . By engineering the interfaces, maximum efficiencies of 2.61 % and 5% have been achieved for cells prepared by low temperature and high temperature approaches, respectively 108,110 .

In other study, atmospheric pressure spatial atomic layer deposition (AP-SALD) was used to deposit Cu2O and applied to ZnO/Cu2O solar cells 103 . The low hole concentration of electrodeposited Cu2O imposes the use of thick Cu2O films in order to have a full depletion layer. The depletion width in a semiconductor depends mainly on the carrier concentration and, therefore, synthesizing Cu2O with a higher hole concentration would result in higher VOC values for a same cell thickness and higher electric field within the Cu2O layer 103 . The Cu2O films obtained using AP-SALD presented mobility values from 1.5 to 5.5 cm 2 V -1 s -1 , for deposition temperatures from 125 to 225 ºC. In terms of carrier concentration the films reach up to 10 16 cm -3 , which is three orders of magnitude higher than for electrodeposited films 24 .

Taking advantage of the high carrier concentration of the AP-SALD Cu2O films, a Back Surface Field (BSF) cell with the following structure Cu2OSALD/Cu2OED/ZnOED, where ED stands for electrodeposited, was implemented 103 . In a BSF cell, the top layer, with a high carrier concentration, dopes the layer underneath thus contributing to decrease the depletion width in the cell, and increasing the electric field within the electrodeposited Cu2O layer. In this way, record current densities were obtained for a ZnO/Cu2O cell deposited by low temperature, atmospheric methods, while using a Cu2O layer of around 1

UV Photodetector based on oxides

Considering the same principles as a solar cells, UV photodetectors can be created using high band-gap semiconductors, larger than 3 eV. Over the past years, increasing efforts have been dedicated to the development of self-powered nano-sensors thanks to the increasing demand of autonomous, wireless, nanoscale engineering devices with various functionalities, and to their extensive range of applications 111 . A promising path for the fabrication of self-powered UV photodetectors is to use the internal built-in potential difference of a pn heterojunction at zero bias to prevent the recombination of photogenerated electron-hole pairs and to generate the photocurrent 112 . Other strategies include photo-electrochemical cells 113,114 and Schottky metal semiconductor heterojunctions 112,115,116 .

Owing to its wide direct band gap energy of 3.37 eV and large exciton binding energy of 60 meV at room temperature as well as to its high electron mobility around 200 cm 2 V -1 s -1 , ZnO has been considered as a key n-type semiconductor for modern electronics and optoelectronics [START_REF] Robertson | Limits to Doping in Oxides[END_REF]117,118 . The combination of p-type CuCrO2 or Mg-doped CuCrO2 thin films in transparent pn diodes has previously been achieved with n-type ZnO thin films deposited by pulsed laser deposition 119,120 and magnetron sputtering 121 . However, the resulting transparent p-n diodes presented low rectifying behavior, with the highest reported rectification ratio of 75 between the forward current to the reverse current at ±1.5

V 119 . The poor characteristics for such pn heterojunctions were attributed to the large lattice mismatch between both oxides and to structural imperfections at grain boundaries in the delafossite phase 119,121 .

A promising path for improving the device performances is to introduce nanostructured architectures, such as one-dimensional nanowires (NWs), in order to benefit from their remarkable chemical and physical properties related to their large surface-tovolume ratio 122 . The possibility to form ZnO nanowires (ZnO NW) using various deposition techniques [123][124][125][126][127] , including the low-cost, low-temperature chemical bath deposition 123 , and to take advantage of their unique properties 122,128 has thus raised even more interest. Within the core-shell NW heterostructure configuration, and following the type-II band alignment, ZnO NW-based heterojunction devices offer efficient light trapping processes through optically guided and radiated modes [129][130][131] , efficient charge carrier separation over the small radius of the NW and efficient charge carrier collection owing to the high mobility electron pathways provided by the NWs 129,132 .

ZnO NW based core-shell heterostructures have successfully been integrated into gas sensors 133 , UV-photodetectors 134 , dye-sensitized 135 , and extremely thin absorber solar cells 136 . Among them, the conventional ZnO NW based-UV photodetectors rely on the photo-generation of electron-hole pairs upon UV illumination, and on the hole trapping at the NW surface through oxygen adsorption-desorption processes while unpaired electrons are collected using an applied bias voltage 137 . High photoconductive gains and responsivity were achieved following that approach, but the long relaxation time and the dependence on an external power source both limit their mobility and adaptability. Self-powered UV photodetectors made of ZnO NW core-shell heterostructures in combination with wide band gap p-type inorganic semiconductors such as GaN 138 , CuSCN

Conclusion

The state-of-the-art analysis of copper oxides-based thin films and devices lead us to foresee two scientific gaps in the literature. Firstly, the possible enhancements of Cu2O In this second chapter, we expose the details of the methodology used in the preparation of copper-based oxide thin films, as well as a description of the characterization techniques and the numerical simulations adopted throughout this work. The chapter is divided in three main parts, with a first section dedicated to material synthesis, annealing conditions and devices fabrication. Then, in a second part the characterization techniques for morphological, structural and functional properties are developed. The chapter ends with the development of the numerical simulations of a pn junction, focusing on the physics of the model and the algorithm used for the implementation.

Aerosol-assisted Metal-organic Chemical Vapour Deposition

In the Laboratoire des Matériaux et du Génie Physique (LMGP) in Grenoble, the expertise on different types of chemical vapour deposition (CVD) processes has started since the early stages of the laboratory. The chemical vapour deposition process is defined as a chemical reaction between reacting species, provided as vapours, in the neighbourhood of a heated substrate, which results into the deposition of solid by-products on the surface 1 . The possibility of producing different structures as films, powders or fibres, with a large variety of materials types as metals, oxides, carbides, silicon or nitrides stablished CVD as a powerful and flexible technique for materials production 1 .

A specific variant of the technique, the Metal-organic chemical vapour deposition (MOCVD), is based on the deposition of materials starting from metal-organic compounds as precursors. These precursor compounds are characterized by organic molecules with metallic elements presenting metal-carbon, metal-oxygen-carbon bonds that are decomposed by a reactive gas in a thermally-energetic environment, resulting in the formation of a solid material. In this case, the possible applications already harnessed in the industry are diverse, ranging from optoelectronics (lasers and diodes), telecommunications, solar cells or microelectronics (transistors), with a growth use in the last years due to its versatility 2 . The broad use of the MOCVD is motivated by the various advantages in materials processing, from high purity and molecular homogeneity of the films; particle size in sub-micron range with high reactivity and surface area of the structures produced. There are also economic advantages since it enables lower decomposition temperatures of the precursors, larger processed areas and relatively simple apparatus, when compared to MBE or other techniques. Nevertheless, some drawbacks do exist and mainly concern the risk associated to the use of chemicals, since the environmental impact of the toxicity of reactants and by-products as well as the safety lab operators are still an important issue to be addressed 2 .

From the different types of MOCVD techniques available we focused our research on Aerosol-assisted Metal-organic Chemical Vapour Deposition (AA-MOCVD) 3 . This type of CVD process uses an aerosol medium to transport the precursor into the deposition chamber. It was developed first and patented at Grenoble by the CEA with the commercial name Pyrosol (Pyrolysis of aerosol) 4,5 . AA-MOCVD enables the growth of high-quality and homogeneous thin films from fundamental studies to low-cost and large-areadeposition for industrial applications 6 . The use of aerosol-assisted generation places this procedure within the liquid source delivery CVD processes and offers the possibility to produce films without any vacuum system, at conditions close to atmospheric pressure and even in open air 7 .

The process uses a solution-based precursor that is sprayed into micro-size droplets, forming the aerosol mist. This solution can be obtained from a pure liquid precursor, or by dissolving solid or liquid precursors in a solvent 3 . The formation of an aerosol from a liquid solution avoids the use of highly volatile precursors since the first requirement is their solubility in the solvent. Moreover, for the deposition of multicomponent materials, different precursors can also be mixed, as long as solubility and stability in liquid phase is guaranteed. Compared to other types of CVD processes, this precursor's versatility enables the use of more thermally stable and lower price precursors, reducing the final cost of materials production. Specifically, in the case used for this work, the aerosol mist is generated by an ultrasonic vibration of a piezoelectric transducer at a specific resonance frequency (close to 800 kHz) that maximizes the mist quantity. The size of the created droplets d can be estimated by the following expression:

𝑑 = 𝑘 ( 2𝜋𝛾 𝜌 𝑙 𝑓 2 ) 1 3 (II.1)
Where f is the excitation frequency of the transducer, k is semi-empirical constant, ρl is the density of the liquid and γ the correspondent surface tension 3 . The equation (II.1)

shows that by changing the solvent or the concentration of the solution, the size of the droplets can be varied through the value of the viscosity, even if the resonance frequency is kept constant. An image of the mist generated by mechanical vibration is shown in Figure

II-18.
. After the creation of the aerosol, the mist is mixed with a carrier inert gas that transports the aerosol to the reaction chamber. In the path from the aerosol generation and the reaction chamber, a reaction gas is added to the aerosol and heating of the mist can be induced by external elements. When the mist arrives to the heated substrate, its temperature increases and the solvent is evaporated. Then, the precursor in its gaseous state reaches the heated substrate, evaporates, decomposes and deposition occurs 4,6 . In the Figure II-19, a scheme shows the different steps in a AA-MOCVD process 3 , where all the different steps in the process of a film deposition from the atomization of the aerosol droplets, the transport and heating, evaporation of solvent and precursor, and absorption on the substrate. Other possible phenomena are also represented, as powder production or direct spray-pyrolysis deposition.

Figure II-19 Schematic diagram of the AACVD process for the deposition of films and powders from

Hou and Choy 3 .

The nucleation and growth of the film occurs by thermally-activated surface reaction of the precursor with the reactive gas. During the adsorption on the heated substrate, the bond between the metal and organic components of the precursor is broken, leading to the subsequent interaction between the metal and the reaction gas, which is oxygen in the case of this study. This leads to the formation of small nuclei on the substrate, while the organic species are released and eliminated from the surface. Consequently, the number and size of the adsorbed nuclei increase, which induces the formation of the film due to a continuous arrival of metal precursor and reactive gas. From Carmalt et al 6 .

The density and the growth rate of each nucleus is the main factor to control the size of the grains and stoichiometry of the phase. During all of the process, a group of parameters have to be controlled with the aim of enhancing the deposition growth rate and the film physical and chemical properties. Regarding the film composition, the nature of the precursor solution, the substrate temperature and the oxygen pressure have important roles, while the aerosol parameters such as piezoelectric frequency and power, density of flux, speed, total pressure and nature of the gases have a more direct effect on the growth rate and thickness of the films.

According to Hou and Choy 3 , there are five main advantages to the use of aerosol assisted MOCVD, namely (i) larger number of available precursors at a lower price, (ii) simplicity of the technique for transport and vaporization of the precursors, (iii) high volume transport of the precursor which can increase the deposition rate, (iv) flexibility in terms of atmosphere, since it can work from low pressure to open atmosphere and (v) the accurate stoichiometric control of multicomponent materials. Nevertheless, the stabilization of the aerosol flux and the geometric conditions of its delivery in the chamber are large issues for the efficiency of this procedure 7 .

Metal-organic Precursors

The precursors used in the development of this work were all metal-organic species, where a metal atom is combined with an organic group, which stabilizes it. Since the organic part is unstable under high temperatures, these precursors have to be decomposed using a heated source. In our case, it is the heating plate below the substrate which fulfils this requirement. Different types of precursors where tested, from acetates to 2,2,6,6tetramethyl-3,5-heptanedione (TMHD) precursors, all of them betadiketonates. The main changes are in the number of organic bonds presents in the precursor and the type of organic compound, as well as, the type of bonding, monodendate or bidendate. For all the results

shown in this work, we will focus on acetylacetonate precursors which derive from the There were various reasons motivating the choice of this type of metal-organic precursor, starting from there solubility in different solvents, water, butanol or ethanol, with or without using additives. In the case of Cu(acac)2, the decomposition temperature of 284 °C is slightly lower than the formation of copper oxide phases and lower than the melting point of glass, which allows the deposition process to occur in a range between 300 and 500°C. This precursor presents a complete decomposition into a Cu atom and acac-ligands 9 , which enables the formation of metallic copper, Cu2O or CuO depending on the oxygen partial pressure. Additionally, the precursor has a low moisture sensitivity 10 , it is generally used in the literature, which simplifies the determination of ideal conditions and finally, it has a reduced cost (2.96€/g in Sigma-Aldrich) allowing the use of large amounts of precursor for dense solutions and long times of deposition.

Description of an AA-MOCVD experiment

The preliminary step in the deposition process is the preparation of the solution.

Usually based on solvents as butanol or ethanol, the solution is produced by simply mixing the precursor and the solvent, followed by a constant stirring. As the precursors used were acetylacetonate, ethylendiamine was added to increase the solubility, at a double concentration in relation with the total molar concentration of precursors 11 . In case of a doped materiel, the ratio of precursors indicated is always the fraction between the dopant molar concentration and the total molar concentration of all precursors, being presented as X/Cu+X. The stirring lasted 15 hours until a homogenously solution was obtained. The temperature of the solution is also increased to 50 ºC to promote the precursor dissolution. The substrate cleaning procedure consists in mechanical brushing using paper and a sequence of lubricant solvents: acetone, isopropanol and de-ionized water, followed by an ultrasonic bath in isopropanol for 15 minutes. In order to complete the cleaning step, they are rinsed with de-ionized water, dried using azote and placed in the holder At this moment, the holder with the substrates is attached to the heating plate, Argon is used as carrier gas injected on the top of the solution vessel, while O2 is used as reaction gas, which enters in the system at a higher position, visible in At the end of the deposition, the mist generation is stopped, the gas flux is decreased and the heating of the plate is switched off. After 1 hour, the films can be removed, and the different parts of the equipment are cleaned.

All the different steps previously described have a direct impact on the film properties, in terms of phase, stoichiometry, thickness, homogeneity, grain size, microstructural morphology, dopant absorption and carbon residual content. In Table II-3 we present a summary of the different deposition parameters and the consequent impact on the properties of the deposited materials.

Modification of the reactor configuration

In the middle of the development of this thesis, an incident occurred that lead to significant changes in the deposition system. As a consequence, two types of reaction configurations were tested.

The initial configuration, named "vertical configuration", visible in Figure II-27, had a vertical inlet on the aerosol flux, combined with a tilted substrate holder. In this configuration, mainly butanol was used as solvent. The problems of this configuration were mainly associated to the risk for a macroscopic drop of solvent or a large non-dissolved precursor particle to hit the heated substrate and to create an explosion. Additionally, the use of butanol in a heated oxygen rich atmosphere could cause auto inflammation, since the auto-inflammation point is 343 ºC and the low flammability limit (LFL) is 1.7%. Therefore, a second configuration, named '"horizontal configuration" was tested, presented before in the deposition description, where the flux enters in the chamber horizontally and the heated substrate holder is kept in a vertical configuration. Moreover, the solvent was also changed to ethanol with an auto inflammation point at 363 ºC and a LFL at 3.3%. However, this configuration can impact the laminar flow of the aerosol and increase the presence of organic contamination in the films.

Annealing treatment

The improvements in optical and electric performance of oxide-based materials can, in general, be achieved by post deposition annealing treatments 12 . In our case, different annealing studies were conducted in a traditional box furnace under air. Two main parameters were varied: temperature variation from 200°C to 450 ºC and anneal duration from short plateaus of 5 minutes up to 5 hours. Previous tests were conducted under O2 atmosphere, however, they show similar results as the ones under air. Therefore, the annealing treatments here presented were developed under air conditions. Additionally, annealing treatments were conducted in a 2-probe stage that allowed the measurement of resistance while heating, once again under air.

Metals Thermal evaporation

The finalization of some devices as well as the electric analysis of thin films required metallization of contacts in order to improve and simplify the measurements.

Usually, gold or silver contacts were thermally evaporated by Joule effect heating of the source in a vacuum chamber using highly pure wires of both metals. The process is based on the increase of temperature of these metals by applying an electric current to a molybdenum crucible. The metal passes from a solid to liquid and then vapour state that travels to the surface of the sample, previously covered by a designed mask. The thermal

Techniques of characterization

In the course of the thesis, a detailed characterization of the synthesized materials was conducted using a wide range of standard analysis techniques. Separated in two different axes, the first group of techniques relied on thin films characterization, focusing on crystallographic, microstructural, optical and electrical properties. In parallel, the analysis techniques of the second group were centred on device performance both in electrical and optical characteristics.

Morphological, chemical and structural characterization

Visual aspect and Optical Microscope (OM)

The first elemental analysis of the films was conducted by visual inspection of the grown specimen. After the removal of the film from the chamber, different characteristics can be detect by the naked eye: thickness of the film and respective gradient, presence of oxide powders, dust or other centres of inhomogeneity. Specially, the estimation of thickness is accurate from the fringes of refraction in the silicon substrate, covered by native SiO2. As a colour chart of Cu2O on SiO2 could not be found in the literature and the refractive light is dependent on the refractive index of the material, a colour chart of Si3N4 was used to estimate the thickness, visible in Figure II-30a). This material has a refractive index of 1.98 at 1000nm 13 , while Cu2O presents a value of 2.4 14 at the same wavelength.

Scanning Electron Microscope (SEM), Energy dispersive spectroscopy (EDS) and Wavelength dispersive X-ray spectroscopy (WDS)

The imaging of the microstructure of the samples and the determination of their physical thickness were obtained using a scanning electron microscope (SEM). This technique has a large domain of application especially important in physics, chemistry and materials engineering but also in biology or medical science. SEM allows to obtain images with high resolution and magnification, showing the topography of the sample, through the interaction of electrons emitted by an electron gun with the surface of the material 16 The energy dispersive spectroscopy (EDS) was additionally performed to analyse the chemical composition of the sample at the microscopic scale. In this case, the energy of the incident electron beam excites electrons located on the atom inner shells, causing their injection. Due to the displacement from a higher energy shell electron and consequent relaxation of the exited charge, the energy difference is release by an X-ray photon. These energies depends on the shell of the electron in each element and can be used to detect different elements on the specimen. In general, an energy beam of 5 keV was used to quantify the different elements in the film without a significant signal from the substrate.

The Cu Lα line is at 0.930KeV, while in the case of Mg Kα line is at 1.253 KeV.

Additionally, elemental maps can be created to observe segregation of elements. The SEM and EDS were conducted in a FEI Quanta 250 MEB FEG ESEM tool visible in the Figure

II-31c).

Complementary to EDS, Wavelength dispersive X-ray spectroscopy (WDS) was also performed in a selected number of samples. This technique, based on the same principle as EDS, provide additional information about the elements present in the film in a volume of 1µm 3 . By performing analysis at different acceleration voltages (12, 16, 22KeV) it is possible, via software, to determine the composition and the surface mass of the film by using software the Stratagem. The equipment used was a CAMECA SX50.

Transmission electron microscope (TEM)

The transmission electron microscope (TEM) operates on the same basic principles as the transmission optical microscope but uses electrons as in the SEM case. However in this case, the information is obtained through the electrons that cross the sample. Moreover, this technique allows to observe the fine structure of the specimen (observation mode), as well as the crystallographic characterization (diffraction mode) 18 . In the observation mode, INCA Energy TEM 100 X-Max 65T. Cross-section samples were prepared from films deposited on silicon and glued to a copper grid by tripod polishing resulting in a sample thickness of about 10 μm. Argon ion beam milling was used until perforation of the interface. Automated crystal phase and orientation mapping (ACOM) with precession (ASTAR) system implemented in the JEOL 2100F FEG microscope was also used to obtain the crystal phase and orientation maps. This method is based on the precession of the primary electron beam around the microscope's optical axis at an angle of 1.16° while collecting the electron diffraction patterns at a rate of 100 frames per second with a step size of 1 nm.

Atomic force microscope (AFM)

In the group of microscopy analysis, another technique used in this work was atomic force microscopy to measure the roughness of the sample. This technique uses mechanic interaction of a small tip with the atoms at the surface of the specimen. Located at the end of a cantilever, the tip is deflected by peaks and valleys on the film, due to forces of attraction or repulsion at the nano-newton scale 19 . Two modes were used to scan the thin films. In the contact mode, the tip is pressed on the surface of the film and it moves through

X-Ray diffraction (XRD)

The analysis of the crystallographic structure of the thin film was performed by Xray diffraction (XRD). The interaction of X-rays with the atoms of a crystal allows a suitable probing of the atomic positions and consequentially their arrangement in the crystal structure. The interference of reflected X-rays is possible due to the small wavelength of these electromagnetic waves, which are comparable to the lengths between crystallographic planes, producing diffraction patterns that can be subsequently detected. The variation of the incident beam in the sample will satisfy different diffraction conditions determined by Bragg's equation (II.2):

2𝑑𝑠𝑖𝑛(𝜃) = 𝑛𝜆 (II.2)
where 𝑑 corresponds to the distance between crystallographic planes, 𝜃 is the diffraction angle of the light, 𝑛 is the order of diffraction and 𝜆is the wavelength used, in this case the Cu K-edge (0,1519 nm). A scheme of the interplanes distances and X-Ray diffraction is visible in Figure II-34 a). The detection of the XRD peaks can then be related to the lattice parameter and structure of the crystalline material 21 , enabling the detection of different phases in the sample and differences in lattice size. with a scanning rate of 3 degrees per minute. In order to evaluate the grain growth of the Cu2O films, the crystallite sizes in the films was estimated using Scherrer's equation 23 :

𝐿 = 𝐾𝜆 𝐵 cos (𝜃) (II. 3)
where L is the mean crystallite size, B is the full width at half maximum (FWHM) of the intensity peak, K is the Scherrer constant, λ is X-ray wavelength and θ is the Bragg angle.

Spectroscopic and spectrometry techniques

Raman spectroscopy and Photoluminescence

Concerning the chemical analysis of the films, Raman spectroscopy is a powerful tool to access vibrational, rotational, and other low-frequency modes in a material. This type of spectroscopy technique relies on the Raman scattering (inelastic) occurring within the samples upon illumination with a monochromatic light source. When a high intensity monochromatic light, in this case a laser, impinges a material, most photons are scattered with the same energy as the incident light. This elastically scattering process is called Rayleigh scattering. However, some of the scattered photons, around 1 in 10 millions, have a shift of energy in comparison with the original energy. In case of a lower or higher energy scattering, this inelastic process is called Raman scattering, with a scheme visible in Figure II-35 a). The intensity of photos regarding each energy difference are directly linked to chemical bonds, therefore, a spectrum of Raman light can be obtained 24 . In this technique, it is usual to use wavenumber, cm -1 , instead of wavelength, since we always present the shift from the original light source. In the context of Cu2O thin films, this technique was key to for the detection of the CuO phase. -1 , with the use of different monochromatic lights. In the majority of the analysis, the 488 nm blue line of an Ar + laser with a power of 0.4 mW on the substrate surface was used. The laser was focused to a spot size close to 1 µm 2 by using a 50x long working distance objective.

Moreover, photoluminescence measurements were also performed in the same equipment. However, in this case, the emission of light from the sample after the excitation occurs due to different phenomena, mostly excitons and defects in the material 26 . The interest of PL measurements for Cu2O thin films is mainly connected to the detection of vacancies of copper or oxygen.

Fourier transform infrared spectroscopy (FTIR)

Complementary to Raman spectroscopy, Fourier transform infrared spectroscopy working under vacuum. Spectra were recorded with a resolution of 4 cm -1 by accumulating 64 scans in transmission mode using silicon substrates, transparent to infrared light.

X-ray photoelectron spectroscopy (XPS)

X-ray photoelectron spectroscopy is a surface technique that analyse the energy of bonds between atoms. A monochromatic beam of X-Ray bombards the sample, causing the ionization of the atoms. Each ejected electron will be emitted from the atom and by measuring its kinetic energy, Ek, and with previously known energy of the X-ray beam, hv, the bond energy, Eb, required for the ionization can be determined:

𝐸 𝑏 = ℎ𝑣 -𝐸 𝑘 (II. 4)
The binding energy distribution provides quantitative information about the atomic composition of the film. Moreover, spectra obtained with the highest resolution can additionally probe the nature of the chemical bonds and oxidation state of the atoms 28 .

Figure II-37 X-Ray interaction with atoms during XPS analysis b) K-alpha spectrometer at SIAM -Synthesis, Irradiation and Analysis of Materials in the Université de Namur

Integrating the electrons at different energy bonds produces a spectrum where the peaks correspond to characteristic orbitals of each element. In the context of this study, XPS was fundamental to observe differences in copper oxidation state, in order to detect the presence of CuO in the films. This technique is very sensible to the surface (several Angströms).

The analysis was conducted in SIAM -Synthesis, Irradiation and Analysis of Materials in the Université de Namur, using a K-alpha spectrometer, from Thermo Scientific, with a X-ray source Al Kα1,2 (1486.6 eV), represented in Figure II-37 b).

Extended X-Ray Absorption Fine Structure (EXAFS)

In order to probe the chemical environment of a single element and obtain the interatomic distances, Extended X-Ray Absorption Fine Structure (EXAFS) was performed with synchrotron at a large scale instrumental facilities. In a simple approximation, EXAFS measures the energy dependence of X-rays absorption by an atom near a defined core level. These spectra show oscillations near the band edge that can be correlated to different near-neighbour coordination shells and respective distance to these neighbours 29 . In this case, the analysis was performed at ESRF, in Grenoble, in the BM31 line, where the Cu K edge line was studied in detail. The use of EXAFS was fundamental to observe the neighbouring environment of the Cu + cation by measuring distances to neighbour atoms.

UV-VIS-IR transmittance

The optical characterization of the films was in a large part based on optical spectroscopy analysis, where the transmittance was the main analysed property. The transmitted part of the light is then collected in an integrated sphere and the detector compares the intensity obtained with a previously measured baseline. A representation of this technique is shown in Figure II-38 a). Additionally, in the system used, we can discriminate between total and direct transmittance. nm to 2500 nm, presenting a wavelength step of 5nm. The reported average transmittance in the visible is defined as the ratio between the integrated transmittance between 390nm and 700nm and the wavelength interval. The spectra were used to calculate the band gap by tracing the Tauc plot using the PARAV-V2.0 software, with a given uncertainty of 0.1eV.

Electrical and functional properties of materials and devices

4-Probe

Combined with optical inspection, the determination of the sheet resistance by 4probe measurement is one of the first analyses to be performed in the sample. In this approach, the applied current flows between the outer probes while the voltage is measured in the inner probes, providing a resistance value based on Ohm's law. The combination of two different conditions allow the conversion of the measured total resistance to sheet resistance and finally to resistivity. The first requirement relies on the film thickness to be less than half of the constant spacing between the probes; while the second important requirement imposes the measuring point to be away from the edges of the film, at least 4 times the spacing between probes. Under those conditions, the sheet resistance can be given by:

𝑅 𝑆 = 𝐹 𝑉 𝐼 (II. 5)
where RS is the sheet resistance, V the voltage measures, I the current applied and F a constant defined as:

𝐹 = 𝜋 𝑙𝑛(2) = 1.442 (II. 6)
Finally, the resistivity is determined by following equation:

𝜌 = 𝑅 𝑠 𝑒 (II. 7)
where ρ is the resistivity and e the thickness of the film. At LMGP, the measurements were performed using a Por While applying the current between two contacts in one direction, the voltage is measured in parallel contacts. It is demonstrated that the two measured resistances are related to RS by the following expression:

𝑒 - 𝜋𝑅 𝐴 𝑅 𝑆 + 𝑒 - 𝜋𝑅 𝐵 𝑅 𝑆 = 1 (II. 8)
The numerical resolution of this equation can provide the sheet resistance and converted into resistivity, by the expression previously described.

Hall Effect measurements

Complementary to van der Pauw measurements, Hall-effect measurements were used to determine the density and the mobility of free charge carriers of the film. The Hall Effect is based on the charge build-up in a medium due to magnetic forces. A magnetic field, B perpendicular to the film surface, produces a force on moving charges, in an electric current I, those charges are deflected in a direction perpendicular to the original movement and the magnetic field. The caused force is called Lorentz force, FL. The magnetic part of this force can be expressed as:

𝐹 𝐿 = 𝑒 𝑣 x B (II. 9)
where e (1.602 x 10 -19 C) is the elementary charge, v the velocity of the charge. The perpendicular movement of the charges creates a transversal potential to the electric field sense, which can be represented as: 10) where VH is the Hall potential, I is the current, p the carriers density and t is the sample thickness. The measured Hall potential can then be used to determine the nature and concentration of charge carrier, as well as their mobility µ, as follows: 12) In Figure II-40b), a scheme of VH measurement is represented. The electronic transport properties were obtained at room temperature using a Van der Pawn configuration and a RH2035 Hall effect measurement setup from Phys Tech equipped with a magnetic field of 0,43T.

𝑉 𝐻 = 𝐼 𝐵 𝑒 𝑝 𝑡 (II.
𝑝 = 𝐼𝐵 𝑑 𝑒|𝑉H| (II. 11) µ = |𝑉H| 𝑅 S 𝐼𝐵 = 1 𝑒 𝑝 𝑡 𝑅 S (II.

IV and PV measurements

The characterization of the fabricated devices first involved current-voltage (I-V) measurements, which are fundamental to evaluate the quality of the pn junctions. By applying a bias voltage between 2 electrodes and sweeping different voltage values, the resistance between the two contact points is measured simultaneously, resulting in an electric current calculation. Additionally, if this analysis is performed under a light source, we can access to the efficiency of a solar cell, a key performance indicator for such a device.

Responsivity

The responsivity analysis was used to evaluate the UV selectivity of the photodetector. The technique measures the amount of electrical output from the device due to an optical environmental input. In this specific case, we measure the current generated in the detector per incident radiant power of a specific wavelength. Therefore, the responsivity is expressed in Amperes per Watt. The analysis was performed at room temperature in the range of 300 to 800 nm using a xenon arc lamp. The short-circuit current (ISC) from the sample was recorded every 10 nm across the spectrum and compared to the ISC from a calibrated Si-photodiode with a known responsivity.

Simulation

Device modelling was performed to simulate the electrical characteristics of pn junctions formed with Cu2O and ZnO. Numerical analysis is a powerful tool to analyse a large quantity of different phenomena, providing valued information to improve the experimentally grown materials and consequently the applied devices. This study was used for a direct comparison with the experimentally fabricated junctions, in order to access carrier's concentrations distributions, size of depletion region and ideal diode behaviour.

This part of the study was developed at the Université de Liège, using a Fortran program originally created within the SPIN laboratory. In this section we explain the basic semiconductor equations used in the model, as well as some technical details on the procedure leading to the computer program.

Basic equations

In a first approach, Maxwell's equations are rewritten to obtain the distribution of charges and the electrical potential. In the differential form, we use Gauss's law to relate the electric flux leaving a certain volume the electric charge density. In an isotropic material, the displacement vector (D) is proportional to the electric field (E) by a scalar quantity, the semiconductor permittivity 𝜀 𝑠 . Therefore, the electric flux equation can be simplified as:

𝜀 𝑠 ∇ • 𝑬 = 𝜌 (II. 13)
Subsequently, we have to consider Faraday-Henry's law of electromagnetic induction and the vector potential A relation to the magnetic field B:

∇ x 𝑬 = - 𝜕𝑩 𝜕𝑡 (II. 14) ∇ x 𝐀 = 𝐁 (II. 15) ∇ • 𝐀 = 0 (II. 16)
The combination of the previous equations considered leads to:

∇ x (𝑬 + 𝜕𝑨 𝜕𝑡 ) = 0 (II. 17)
This last expression results in a creation of a gradient field, -∇𝜓, due to the fact that the curl of a vector field is equal to zero. Therefore, the electric field is equal to:

𝐄 = - 𝜕𝑨 𝜕𝑡 -∇𝜓 (II. 18)
As a consequence, Gauss's law can be expressed as:

∇ • 𝜀 𝑠 𝜕𝑨 𝜕𝑡 + ∇ • 𝜀 𝑠 ∇𝜓 = -𝜌 (II. 19)
Considering the precious definition of A, ∇ •A=0, we obtain a final expression for the gradient field of:

𝜀 𝑠 Δ𝜓 = -𝜌 (II. 20)
The second side of the equation, the space charge density 𝜌, can also be considered in terms of the sum of all charged defects multiplied by the elementary charge q. In the group of total defects we count as contributions to the total positive charge the hole concentration 𝑝 and the ionized shallow donor concentration 𝑁 𝐷 , while the negative charge incorporates the electron concentration 𝑛, the ionized shallow acceptor concentration 𝑁 𝐴 , and the occupied trap concentration 𝑛 𝑡 * . Putting all terms together, we reach the following expression for the total charge density :

𝜌 = 𝑞(𝑝 + 𝑁 𝐷 -𝑛 -𝑁 𝐴 + 𝑛 𝑡 * ) (II. 21)
The next part of the model development is connected to how the space charge density varies with respect to time which leads to the continuity equations. The application of a divergence to the Ampère's circuital law leads to the relation between current density and the time derivative of the charge densities.

∇ • (∇ x 𝑯) = ∇ • (𝐉 + 𝜀 𝑠 𝜕𝑬 𝜕𝑡 ) (II. 22)
Since the application of the divergence operator to the curl of a vector field is equal to zero, we can extract that:

∇ • 𝐉 = -∇ • (𝜀 𝑠 𝜕𝑬 𝜕𝑡 ) (II. 23)
And with some extra manipulation of the (II. 23) equation, we obtain:

∇ • 𝐉 = - 𝜕𝜌 𝜕𝑡 (II. 24)
We can separate the current density in the sum of the different components, electrons and holes and by considering the donor and acceptor dopant as time independent we obtain: where Fn, Fp and Ft are the quasi-Fermi energies of electrons, holes and trap states. Ec and Ev are the position of the conduction and valence band respectively, while gt is the degeneracy factor of the trap states.

∇ • 𝐉 𝐧 -𝑞 𝜕𝑛 𝜕𝑡 = 𝑞 𝑅 𝑛 (II.
The numerical simulation was then based on this group of equations: the relation between the different current densities and the charge carriers and electric field, as well as the charge carriers' density distributions.

Model's algorithm

The set of equations presented here above is a system of nonlinear, coupled, partial differential equations. Therefore, linearization and discretization of the initial problem into a step-sized mesh are required. The system of equations is linearized by a Newton approach and the initial values are introduced using a guessed solution. The algorithm chosen to solve this model is based on iterative methods that involve a series of corrections to the initial guess. These iterations continue until a defined set of convergence criteria is reached. These methods are usually quite efficient for large linear systems, converging rapidly and using memory efficiently. Though, the results are inaccurate within the round-off error and the convergence is not always obtained.

The whole problem size depends on the number m of unknown quantities per node and the total number of mesh points N, which provides a total number of unknowns that can be expressed by the product m×N. In the case of Newton's method, which is used in this framework, the coefficient matrix contains (m×N) 2 elements and exhibits a blockdiagonal shape, for which dedicated algorithms are known. The quality of the numerical solution to the equation system depends thus on an optimal trade-off between precision and size of the linear system. If we consider a large number of mesh points, the quality and precision of the solution rises, with an increase in computation time and memory capacity requirement. For more specific details, the reader can consider the works by Baert 32 and

Nguyen 33 where the authors present a full description of the physical considerations and the algorithm behind the implemented model.

Chapter III: Cation-doped Cuprous Oxide thin films In order to improve the optical and electrical properties of Cu2O films as a p-type semiconductor, possible cations as dopant candidates were screened from a large group of atoms. This general study allows to narrow down the list of potential dopants and to identify the most promising elements to be used in a more in-depth experimental work.

As it is well known, the properties of a thin film are strongly dependent of the synthesis method, therefore, we established the "background line" by performing a study on the properties of Cu2O thin films deposited by AA-MOCVD and correspondent annealing process as the starting point of the study. Then, using a mixed solution of precursors, including the dopant source, the incorporation of the doping species into the copper oxide films was investigated. Here, we present in more detail the use of several cations as dopant which had an impact on the structural properties as well as on the electric response, with a special consideration to Tin as a dopant due to resistivity improvements.

The study containing the incorporation of Mg is comprehensively reported in the next chapter (IV), fully devoted to this cation doping system.

Intrinsic Cu 2 O thin films

The deposition of intrinsic Cu2O by aerosol-assisted MOCVD was preliminarily studied in order to obtain the required crystallographic phase, homogeneity, thickness, electrical and optical properties. In this first study, the aerosol-assisted apparatus and the home-made reactor were assembled as described in configuration 1 in Chapter II. An important optimization work on the deposition of undoped Cu2O deposition by Pulsed injection MOCVD (PIMOCVD) was developed previously by L. Bergerot 1 . The results of that study were used in this research as supporting information for the parameters variation in the system. Although the techniques differ substantially, as PI-MOCVD operated under vacuum and aerosol-assisted MOCVD is developed at atmospheric pressure; notably, they share a few common features such as the composition of the precursors used in solution, as well as their concentration and range of temperature during deposition. This comparison also provides interesting conclusions for the understanding of the influence of the oxygen partial pressure.

The main studied parameter was the substrate temperature during deposition. Deposition temperature was increased from 300 °C to 375 °C by steps of 25°C. The The investigation will therefore aim at optimizing the Cu2O with other parameters.

It was also observed from this study that most of the effect on the optical properties is due to the variation of the film thickness.

In order to reduce furthermore the sheet resistance of intrinsic Cu2O films, postdeposition annealing treatments under air were studied. The duration of the post-deposition anneal was 1 hour, which is in the same range of time as other previous studies 2 , 3 . The temperature varied from 225 °C to 300 °C. The use of pure oxygen instead of air was also tested. However, no significant difference was observed in the oxygen case when compared From these preliminary results we can infer that for the optimized films of pure Cu2O, with a thickness of 100 nm, the lowest value of sheet resistance is around 20 MΩ/sq, corresponding to a resistivity value of 200 Ω.cm and a transparency of 50% in the visible electromagnetic spectrum.

Cation doped Cu 2 O thin films: screening different elements

In the perspective of improving the optical and electrical properties of Cu2O films, the incorporation of different cations as potential dopants was systematically studied. The elements used in these trials were chosen based on different characteristics, mainly theoretical predictions from Nolan 6 , atomic radius of the cation, compatibility of the required precursor with AA-MOCVD. Therefore, we present in Table III-5, the list of cations tested, with the respective atomic radius and precursor used. The deposition were conducted using a 10% of the dopants precursor molar concentration in relation to the total concentration of precursor in solution. The rest of deposition parameters were maintained constant as previously described in Table III-4, with the deposition temperature fixed at 350°C. The films were also annealed at 250 °C for 1 hour under air.

The preliminary characterization of these films was based on EDS analysis, to detect the presence of the dopant, and on sheet resistance measurements, to confirm the functional enhancement of the film resistivity. Among the total group of cations, only Mg, Co, Ni and Sn were detected by EDS analysis. However, it is important to state that the Li cation detection is under the limit of EDS due to the low atomic weight. Nevertheless, no significant differences in sheet resistance were observed with the use of Li.

From the point of view of electrical properties, the presence of Mg and Sn lowered the sheet resistivity to values below 10 MΩ/sq, which present significant improvements comparing to 20MΩ/sq of intrinsic Cu2O. In terms of optical appearance and transmittance, no significant changes were visible in the transparency with the dopants used at 10% of concentration. Therefore, we continue the study focusing on Sn and Mg-doped systems, since both of these dopants show improvements of the electric transport properties; they were detected by EDS in a relevant amount and their radius were similar to Cu + cation.

Nonetheless, the deposition of Cu2O with cations had an effect on the microstructure morphology even in the case of cations not detected or weakly detected by EDS. In  Even when no detection by EDS could be achieved, Li precursor creates a film with a granular shaped surface and visible porosity between some grains.

 The presence of Ca(acac)2 during the Cu2O deposition seems to increase the porosity of the film; however, these features have been also observed in some intrinsic Cu2O films.

 Both Co and Ni were detected by EDS however with different effects on the microstructure. While Ni seems to create a rough surface similar to the intrinsic film, the Co presence forms flower-like structures on the top of the surface. These structures seem to be Co-rich by EDS analysis, and generally dispersed throughout the whole film.

 The presence of Zn leads to an observable granular structure, with small grains better revealed when compared to the Cu2O films.

 At last, Sr incorporation yields a strongly rougher surface with grain clusters visible, which are similar to the previous work by PI-MOCVD 1 and also to the films doped with Mg, represented as well in Figure III-46. However, the presence of Sr could not be detected by EDS analysis, even if the experiment was reproduced several times with different concentrations of the precursor solution.

Further analysis of the Cu2O growth with the presence of these elements could be relevant to understand the origin of these morphological changes. The effect on the microstructure could be attributed to different growth mechanisms, not explored in this work. From these results, tin and magnesium were selected as doping elements. In the next section we present the characterization of Cu2O films co-deposited with tin. The combination of Sn with Cu2O was not found in the literature, which additionally contributed for its interest. The study on magnesium doped Cu2O films is carried out in Chapter IV. A recent study from 2016, Kardarian et al. 8 , presented the incorporation of Mg into Cu2O and their application to solar cells. However, a comprehensive study of the films'

properties with the dopant concentration was missing, providing a motivation for a systematic work reported in the next chapter (IV).

Tin doped Cu 2 O thin films

Tin oxide presents different oxidation states with different semiconductor behaviours: SnO2, a well-known n-type semiconductor oxide, and SnO that behaves as a p-type semiconductor, with resistivity values close to 1 Ω.cm 9 . The formation of each phase is dependent on the oxygen partial pressure during deposition or annealing treatments 10 , where the transformation from SnO to SnO2 common under oxidizing atmosphere 11 .

The improvements in sheet resistance with the Sn incorporation lead us to yield some focus on this cation as a dopant for Cu2O. To incorporate the dopant in the films, we added Di-n-butyltin diacetate to the precursor solution in different concentrations, 0.002, 0.01 and 0.02 mol.l -1 , which corresponds to a Sn/(Cu+Sn) atomic ratio in the solution of 9%, 33% and 50%, respectively. The other deposition parameters were maintained as presented in Table III-4, as in the previous depositions.

The increase in concentration from 9% to 33% of Sn precursor in solution induced a large impact in the optical appearance, as noticeable in The presence of SnO2, a well-known n-type semiconductor oxide 12 , lead to lower resistivity, however with a change of charge sign confirmed by Hall Effect. In this sample, the resistivity values were around 0.24 .cm, and the respectively mobility and freeelectrons density was 1.6 cm 2 .V -1 .s -1 and 1.7 x 10 19 cm -3 .

As brief conclusions of this dopant, Tin was selected as a promising candidate to improve the optical and electrical properties of Cu2O, however, low mobility and formation of n-type SnO2 phase were detrimental for the applicability as a p-type transparent semiconductor. In a future work, the analysis of this system using lower content of Sn can provide an additional route for the enhancement of Cu2O thin films.

of

Magnesium-doped cuprous oxide (Cu 2 O:Mg) thin films

Introduction

As previously discussed in Chapter I and III, cuprous oxide (Cu2O) presents promising electrical and versatile processing features that establish it as a highly researched p-type transparent semiconductor. However, its high absorbance in the visible range reduces the application in practical devices. In 2008, based on first-principle calculations, Nolan et al. 1 suggested the doping of cuprous oxide with cations larger than Cu + in order to increase the band gap, while maintaining the cubic structure. These cations, such as Mg 2+ , Sn 2+ , Sr 2+ or Ca 2+ would distort the crystallographic lattice and diminish the threedimensional Cu-Cu interactions. As a consequence, the Cu 3d and 4s partial electronic density of states would be changed which could ultimately led to a band gap increase, depending on the cation used. The p-type conductivity would simultaneously be enhanced, due to the creation of a double copper vacancy. Since the dopants are divalent cations, one primary copper vacancy compensates the dopant presence, while a secondary copper vacancy supplies an extra hole 1 . Additionally, the use of "electronically inert" dopants would avoid the hybridization of the valence band and conduction band edge states or the introduction of in-gap states 1 . Both optical and electrical consequences of Nolan's hypothesis would suggest improvements on the Cu2O application into devices.

In a previous work, the incorporation of strontium (Sr) had already been achieved, showing improvements in the resistivity of the films, up to 1 Ω.cm; but without any change in the optical gap energy of the material 2 . The lower resistivity of Cu2O:Sr was associated to an increase of simple copper vacancies (VCu) -, attributed to a decrease of its formation energy with strontium presence 2,3 . Moreover, the incorporation of strontium also leads to the emergence of a coexisting shallow acceptor level, which is tentatively assigned to large size impurity-vacancy complexes 3 .

In this study, we deposited magnesium-doped cuprous oxide (Cu2O:Mg) thin films by aerosol-assisted metal-organic chemical vapour deposition (AA-MOCVD). The choice of magnesium as dopant was motivated by theoretical predictions on the band gap changes 1 , the ionic radius and the set of deposition conditions which is comparable to that of Cu2O 4 . As discussed extensively in Chapter II, AA-MOCVD is a well-established deposition technique for copper-based oxides, using cheap and efficient metal organic precursors and allowing the growth at atmospheric pressure, avoiding high vacuum.

Deposition of Cu 2 O:Mg thin layers by AA-MOCVD

We used a butanol-based solution for the deposition of Mg doped Cu2O thin films by AA-MOCVD. The copper precursor was copper acetylacetonate (Cu(acac)2), and the molar concentration was fixed at 0.02 mol.l -1 all along this study on magnesium doping.

To incorporate magnesium in the films, we added magnesium acetylacetonate, Mg(acac)2, in different molar concentrations: 0.001, 0.002, 0.004 and 0.013 mol.l -1 , corresponding to Mg/(Mg+Cu) atomic ratio of 0%, 5%, 9%, 17% and 40%, respectively. All the other parameters were maintained as described in the Chapter III, where Cu2O thin deposition was optimized.

The deposition was performed at 350°C in a homemade cold-wall chemical vapour deposition chamber on substrates of alkaline earth boroaluminosilicate glass (Corning 1737) or p-type single crystal silicon wafer (001) oriented. Argon was used as carrier gas, with a flow of 6 l.min -1 , while O2 was used as reaction gas, with a flow of 2.5 l.min -1 . These flow-rate values allowed us to maintain a solution consumption rate of 2 ml.min -1 . The deposition lasted for 60 minutes, resulting in a consumption of 120 ml. After the deposition, the films were annealed at 250°C under air for 1 hours to improve physical properties. These conditions were previously optimized on intrinsic cuprous oxide (Chapter III), in order to obtain monophasic films with no change in colour.

Structural characterization

As a primarily qualitative analysis of the samples, the optical appearance of an Mg-

The Mg ratio indicated was quantified by EDS analysis (Table IV-6). The different elements X-Ray emission lines are also indicated.

The magnesium incorporation in each film was quantified from the EDS spectra, using the P/B-ZAF method provided by Bruker. The Mg/(Mg+Cu) ratio is presented in to the absorption edge. Nevertheless, the increase in transparency above 500nm is more abrupt for the highly doped Cu2O film than for the others. An improvement of 6% in transparency is observed, even when the thicknesses are comparable, around 80 nm.

Additionally, the optical band gap values as calculated from the Tauc representation, are given in Table IV-8, which also confirm a small increase of 0.2 eV from the intrinsic to the highly doped film. .cm 2 . Nevertheless the minimum resistivity value is still higher than for N-doped Cu2O thin films, with a resistivity value of 2x10 -1 .cm 10 . decreasing behaviour with Mg content, and varies from 6.8 cm 2 .V -1 .s -1 in the Cu2O film to 1.1 cm 2 .V -1 .s -1 in the 17% Cu2O:Mg film. On the other hand, the charge carrier density increases, ranging from 4.5x10 15 cm -3 up to 8.1x10 17 cm -3 for an increasing in Mg content from 0% to 17%. The two combined properties lead to a general decrease of resistivity, as

visible in Figure IV-62.
The values of carrier concentration obtained in this work are higher than those found in the literature, which range from 10 13 to 10 16 11 . This fact can be explained by the substitution of copper by magnesium, creating copper vacancies 12 . Indeed, two possible copper vacancies can appear in Cu2O: the simple copper vacancy (𝑉 𝐶𝑢 ) related to the removal of one copper atom, leaving two oxygen atoms with three copper neighbours; or the split copper vacancy (𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 ) in which the copper disappearance is followed by a neighbouring copper atom movement towards the vacancy 12 . In the latter case, the copper atom moves into a tetrahedral site with four neighbouring oxygen atoms. The split vacancy is usually associated to highly localized hole, responsible for the change in oxidation of copper cation from 1 + to 2 + , similar to CuO phase.

Magnesium as a divalent cation would be incorporated in a tetrahedral position, similar to a split copper vacancy, as described by Isseroff and Carter 13 . This defect would be electrically inert, since the Mg 2+ cation would replace a Cu + cation and one additional Cu + would be removed, visible in Figure IV-64. However, this incorporation would reduce the formation of Cu based split vacancies and therefore inhibit the nucleation of the CuO phase 13 . Therefore, during the annealing, more simple copper vacancies are created, leading to an increase of the hole concentration without forming the extra CuO 13 . Further analysis on the stability of Mg-doped Cu2O during thermal treatments would be required to confirm the occurrence of this mechanism. The presence of magnesium in the film also has an additional role, both on grain shape and consequently on the carrier mobility of Cu2O. As the amount of dopant inside the film is fairly high (17%), it is reasonable to think that a fraction of the magnesium content is present at the grain boundary. Therefore, hole scattering is more enhanced at the boundaries, reducing the material mobility. As the exact position of Mg in the films was not possible to be clarified by the techniques used until now, other analysis as EXAFS and XPS are fundamental to probe the chemical environment of the cations (Mg and Cu). These results are present in the following section of this chapter.

Summary

The incorporation of magnesium in cuprous oxide thin films was successfully achieved by aerosol-assisted metal-organic chemical vapour deposition under atmospheric pressure condition at 350ºC. The use of physicochemical analysis techniques such as FTIR and EDS confirmed the presence of magnesium up to 17% in cation concentration, while the analysis of XRD and Raman spectra led to the detection of cubic Cu2O material, without the appearance of any magnesium-related phase. Therefore, we can conclude that the incorporation of the dopant atom is localized inside the structure or at the grain boundaries of the cuprous oxide. However, the exact location of the dopant is still unclear based on SEM and TEM micrographs. The presence of Mg affected the microstructure of the film, forming agglomerates of grains, reaching sizes in the hundreds of nanometres, which led to a higher roughness of the films surface.

The incorporation of magnesium in Cu2O films had a strong effect on the electrical properties, reducing the electrical resistivity to 6.6 .cm, by mainly increasing the charge carrier density up to 8.1x10 17 cm -3 . The doping effect can be attributed to the increase of simple copper vacancies, probably by the magnesium presence in a tetrahedral position, leading to an increase of the hole concentration, as proposed by Nolan et al. 1 . This work reports, to our knowledge, the first experimental evidence of a hole concentration increase correlated to the addition of a divalent cation in Cu2O, complementary to previous reports on lower resistivity 2 .

However, the assessment of the impact of the doping on the optical transparency was inconclusive since the bandgap (2.4eV) and the total transparency (51%) only improved slightly when comparable with intrinsic Cu2O.

In conclusion, transparent conductive properties of magnesium-doped cuprous oxide were improved when compared to intrinsic Cu2O thanks to a strong reduction of the resistivity. Ultimately, the control of the charge carrier density can provide a way to use Cu2O in transparent electronic devices in the future. Additionally, the observed reduction of electrical resistivity, associated with an increase of charge carrier concentration, can enhance the efficiency of Cu2O-based solar cells 11 .

Stability of Magnesium-doped cuprous oxide (Cu 2 O:Mg) thin films under thermal treatments

In the second section of this chapter, we will continue the study of Mg incorporation in Cu2O, focusing on the dopant impact on the properties of Cu2O under thermal treatments.

As suggested before, the use of annealing treatments on both intrinsic and Mg-doped Cu2O thin films can clarify the influence of the dopant on the copper vacancy generation mechanism, both simple and split vacancies, as well as the transformation of Cu2O into CuO parasitic phase.

Introduction

The use of post-deposition annealing treatments on Cu2O have been widely studied as a method to increase the free carrier concentration 14 . It has been stablished that annealing in oxidizing environments promotes the creation of copper vacancies, 𝑉 𝐶𝑢 , which is the main mechanism for the increase of p-type conductivity in the material by the generation of holes. The concentration of charge carriers can reach up to 10 16 cm -3 15 , leading to resistivity values as low as 100 Ω.cm in intrinsic Cu2O 16 . While when annealing doped films, the values can increase up to 10 18 cm -3 , which reduces de resistivity to the 1 Ω.cm range 3,17 .

Nevertheless, the use of higher temperature promotes the transition from Cu2O thin films into tenorite oxide, CuO, which usually starts at temperatures ranging from 250 to 300 ºC in an oxidizing atmosphere 18,19 . Even if Cu2O is considered an unstable oxide at room temperature under air, in fact, the transformation to CuO takes place but it is extremely slow, that the system can be considered as thermodynamically stable for the main applications 20 . At higher temperatures, the conversion of Cu2O into CuO is attributed to the diffusion of copper to the surface, which consequently generates copper vacancies.

As explained in the previous section, simple copper vacancy (𝑉 𝐶𝑢 ) and split copper vacancy (𝑉 𝐶𝑢 𝑠𝑝𝑙𝑖𝑡 ) can be formed during thermal treatments in Cu2O. In the split vacancy, the shifted copper becomes 4-fold coordinated, combined with a highly localized hole, which is associated to a nucleation centre for CuO phase. Additionally, the diffusion of copper to the surface of the film in an oxidizing atmosphere can also lead to the formation of CuO due to the direct oxidation of the cation, from Cu + to Cu 2+ . In the case of bulk Cu2O, the Table IV-9, and follow similar trends as observed before for this system. The films were then annealed and characterized from different points of view. The difference in appearance can be evaluated even by the naked eye, as the phase Initially both films present a yellow colour as-deposited. When the films are annealed at 250ºC, no clear difference is visible. However a colour change is detected after 350ºC, in both cases, especially in the undoped case. At 450ºC, we can see a darker film in the intrinsic copper oxide film, while the Mg-doped one present a brown colour, however not completely dark and opaque. This fact led us to infer that the magnesium presence has an impact in the phase transformation from Cu2O to CuO as compared to undoped Cu2O films, since the darker phase is not yet completely visible in the Mg-doped films.

To follow the phase transformation, the structural properties of the films were evaluated by grazing incidence X-ray diffraction (GI-XRD), more sensitive to a phase The GI-XRD spectra of the second set of samples corresponding to films of Cu2O doped with magnesium are presented in Figure IV-67. In this case and similarly to the intrinsic films, the as-deposited and annealed at 250ºC films present only diffraction peaks corresponding to the Cu2O phase. In the film annealed at 350ºC, a small diffraction peak attributed to the CuO phase is detected (111) while the intensity of the Cu2O peaks decreases. For the film annealed at the higher temperature, 450ºC, the intensity of the CuO diffraction peaks is higher, however, the (111) Cu2O diffraction peak at 36.4° is still the treatment is performed at 350°C, small grains start to appear on the top of the original viewed morphology hinting at a size smaller than 100 nm and apparently only formed at the surface. At the sample annealed at 450°C, a widespread of grains occurs with approximate size below 100nm, but in a different way compared to the undoped Cu2O films, as they cover here all the surface. The rough surface of the film is maintained even with the presence of the recently formed grains. The direct comparison between the effects of the annealing on the two sets of samples establishes differences in terms of growth of the secondary phase, CuO. As already stated before in the Mg content study, the main differences in the as-deposited intrinsic and doped Cu2O films are related to the roughness of the film, visible in top view but also in the cross-section. Combining SEM images with XRD results, the formation of CuO is associated to the formation of smaller and brighter grains at the surface. The formation of these grains occurs in a unlike way for the two samples annealed at 350ºC, since in the intrinsic case the newly formed grains seem to be generalized on the surface of the film, while in the Mg doped films the grains are only formed in a fraction of the visible surface.

For the Cu2O:Mg films it is necessary to perform the thermal treatment at 450ºC to obtain a look like that of the undoped annealed at 350ºC, consisting in small grains covering completely the surface of the film. Moreover, the Cu2O undoped films annealed at 450ºC show different shape and larger size of the CuO formed grains, consistent with a higher degree of phase transformation, previously visible in the GI-XRD results.

STEM was used to analyse the grains on the surface of Cu2O:Mg film annealed at 350°C. However, the identification of CuO phase by the diffraction pattern on the formed grains was not obtained. Moreover, EDS was performed punctually using spot analysis in The comparison of the two oxides is based on the appearance of a peak at 348 cm -1 , which is attributed to the Bg 1 mode of CuO without correlation to Cu2O. This is the only peak independent of Cu2O, since the others at 298 cm -1 (Ag) and 628 cm -1 (Bg 2 ) are also visible in the Cu2O. For the intrinsic Cu2O, the Bg 1 mode is barely detected in the film annealed at 350°C, but the most intense are that attributed to Cu2O phase. When annealing at 450°C, the Raman spectrum changes drastically and only peaks of CuO are visible, since all the distinctive modes of Cu2O disappear. On the Mg doped Cu2O set of samples, the detection of CuO is only visible at 450°C, and even at this temperature, the majority of Cu2O peaks are still observed. The FTIR measurements were performed on films deposited on Si substrate to guarantee a transparency to infrared. In this case, only the as-deposited and the annealed (350°C) films were probed. accompanied by satellites at 940-945 eV 23 . The Cu +1 state or Cu 0 presents a doublet at 932 eV with a low FWHM 23 . For the Cu2O as-deposited sample, the satellites at 940-945 eV indicate the presence of Cu +2 . After the thermal treatment at 250°C, the intensity of Cu 2+ satellite increases and the Cu 2p3/2 peaks are additionally deformed, indicating an increase of CuO at the surface of the material. For the samples containing Mg, both as deposited and annealed at 250°C, satellites are not visible confirming that the films at the surface are formed only of deposited and annealed at 250°C, the peaks observed are only attributed to the presence of Cu2O. A complementary technique to XPS that allows to probe the full thickness of the sample and gives information about the chemical environment is EXAFS. This technique is available thanks to Synchrotron facilities in ESRF Grenoble. The two sets of Cu2O samples were analysed and the Cu-K edge was registered. Two extra samples of Cu2O and Cu2O:Mg were annealed at 500°C for 30 minutes and analysed exclusively by this method, in order to have a complete CuO film. Athena software enables the treatment of the X-ray absorption near edge structure (XANES) raw data into EXAFS 24 . After applying a background removing, the representation of the Fourier transformations in k 2 from 2 to 14 Å -1 are presented in Figure IV-74. Each of the visible peaks can be attributed to a direct neighbour of Cu atom and the respective distance. Figure IV-74 a) shows the spectra related to Cu2O undoped samples, annealed up to 500°C. Several changes are visible between the as-deposited and the annealed sample at 500°C in the neighbour distance comprised between 1 and 4 Å. We will present here the hypothesis of the change induced by the oxidation of Cu2O into CuO from the changes in the EXAFS spectra. In the first peak corresponding to a neighbour distance between 1 and 2 Å and identified as the oxygen first neighbour (Cu-O 1 st ) for Cu2O and CuO, we observe a change of intensity of the peak combined with a small shift to the right with the high temperature annealing, above 450°C. These modifications are connected with the increasing number of O neighbouring atoms between the Cu2O and CuO, changing from 2 to 4, as well as a small increase of Cu-O distance, from 1.85Å in Cu2O to 1.95Å in CuO. In the second peak corresponding to a neighbour distance between 2 -4 Å and attributed to cooper neighbouring atoms (Cu-Cu 1 st ), small changes start to appear as a decrease of intensity of the peak with the lower temperature annealing. At temperatures higher than 450°C, with the transformation into CuO, this peak splits into two parts, possibly attributed to copper neighbours of higher order (1 st , 2 nd and 3 rd ), as well as 3 rd order oxygen neighbours.

In the case of Cu2O:Mg, presented in Figure IV-74 b), similar considerations can be inferred for this system, since the first peak also increases in height, and the second peak splits into two. However, it is important to refer the differences in spectrum related to Cu2O:Mg annealed at 450°C. In this case, there is an intermediate state for both peaks, Cu-O 1 st and Cu-Cu 1 st , not detected before for the intrinsic Cu2O system.

In Figure IV-75, the spectra of the Cu2O and Cu2O:Mg group of samples were fitted using the combination of Cu2O and CuO pure EXAFS spectra provided by the boundaries samples of undoped Cu2O films and using the Athena software. This allowed to quantify the amount of each oxide depending on the annealing temperature. Until the annealing treatment at 350°C, both films present similar quantities of Cu2O, more than 85%; however the annealing stage at 450°C, shows drastic differences between the intrinsic and doped Cu2O. While in the Cu2O film, more than 95% of the film is already CuO, in the Mg -doped Cu2O a majority of 60% of the film is still Cu2O. This indicates a delayed or partially inhibited transition from Cu2O to CuO in the presence of the dopant. In the case of annealing at 500°C both films are mostly CuO, however with some residue of Cu2O in the Cu2O:Mg case.

The preliminary fitting of the Cu2O and CuO pure EXAFS spectra using Arthemis software is presented in Appendix B.

Electric characterization under temperature

We use the monitoring of the electrical properties as an additional way to follow the transformation of Cu2O into CuO in situ during the annealing step. As a matter of fact, the annealing at 250°C has been used previously to decrease the sheet resistance of our materials, but a complete evaluation of the effect during the annealing was still then lacking. We specifically devote this section to the task. For this study, the same asdeposited sample was cut in several pieces for the different annealing temperatures and the experiments were conducted with in-situ sheet resistance measurements using a 2-probe system. This setup allows us to observe the sheet resistance variation during the heating and cooling process, as well as the effect during the annealing step at a given temperature. In the Cu2O samples and considering only the final sheet resistance, the most efficient annealing approach is the one performed at 250°C, since we obtain a lowering of the resistance during the annealing step, as well as a lower resistance value after the treatments are finished. The total reduction of the resistance with the annealing is 13%, leading to final sheet resistance of 5.5 MΩ/sq. These results corroborates the selection of the temperature for the optimization of temperature. Annealing at 350°C provokes an increase of resistance during the thermal stage, leading to a final resistance higher than the original film (57%). Finally, the annealing at 450°C leads to an even larger increase of resistance during the treatment, potentially attributed to the formation of CuO. The sample annealed at 450°C presents the worst electrical response.

The following graphs, presented in

We describe now the behaviour during the annealing step at constant temperatures, where the sheet resistance changes in different ways, picture in Figure IV-76 a). On the one hand, at 250°C we observe a constant decrease of resistance, while on the other hand, at 350°C and 450°C, the resistance initially decreases during the heating ramp, but immediately increases from the beginning of the thermal step. The decreases in resistance can be explained by a mechanism associated to the formation of copper vacancies at lower temperature, increasing the density free carriers and reducing the resistivity of the film. The increase of the resistance is related to the phase transformation from Cu2O to CuO at higher temperature, which leads to a general increase of resistivity of the film.

Concerning the Mg doped films, presented in Figure IV-76b), we observe a different overall trend that contrasts with the intrinsic Cu2O samples. The most efficient annealing treatment concerning the decrease of sheet resistance is the one performed at 450°C, with a reduction of 84% on sheet resistance, resulting in a film with 0.7 MΩ/sq. For the three selected temperatures, we obtain a more conducting film than the original one after the thermal treatment. Additionally, if we analyse now the resistance behaviour during each step, it is observed that there is a decrease of resistivity for both 250°C and 350°C. This phenomenon can be attributed to copper vacancies generation, as previously stated for the intrinsic Cu2O. In the 450°C annealing curve, we have a slight increase of resistance during the thermal stage at the fixed temperature, probably related to the phase transformation from Cu2O to CuO, but nonetheless, this film present the lowest resistivity value after the annealing.

Using the variation of the resistivity during the heating and cooling down stages of the annealing treatment is possible to extract the activation energy in the two Cu2O types of film. The Arrhenius plots of resistivity in logarithmic scale for the intrinsic and doped The values of sheet resistance were converted into resistivity to correlate this results to Srdoped Cu2O previously studied 3 . In a general analysis of all the annealing treatments, the slopes of the intrinsic Cu2O films are larger than the Cu2O:Mg films, during both heating and cooling stages, represented by the guided triangles in Figure IV-77 b). In the case of intrinsic Cu2O samples and for the films annealed at 250°C and 350°C the slopes during heating and cooling are similar; however, for the annealing at 450°C, the slope increases in the cooling process, especially at high temperatures (Figure IV-77c)). This phenomenon can be attributed to the higher quantity of CuO, formed at high temperature, which changes the resistivity behaviour under temperature variation. The samples containing magnesium present a different characteristic than that of undoped ones, in particular on the heating part.

There is clearly a change of slope at around 170°C -in the graph at 2.25 (1000/K) -for all 3 annealing curves. By taking this into consideration, we separate the heating ramp of the annealing into different sections, one from room temperature to 170°C, and other from Similar to previous studies on Cu2O resistivity variation with temperature 25 , the linearization of these curves allow us to analyse through the use of an Arrhenius equation, needed to extract the activation energies.

𝜌 = 𝜌 0 𝑒 𝐸 𝑘 𝑏 𝑇 (IV.1)
where ρ is the resistivity, ρ0 a constant, E is the activation energy and kb is the Boltzmann constant. In order to obtain this linearization, we depreciate the impact of the mobility dependence with temperature, which allows us to attribute the activation energy to the defects present in the film. The energies obtained from this fittings are plotted in the Figure IV-77d). We have calculated the energy obtained from curves under heating and cooling, but also taking into account the second slope visible in the Cu2O:Mg samples for temperatures higher than 170°C. For the variation of resistivity of undoped Cu2O films under heating, the increase of slope with the temperature of annealing leads to values varying from 220 and 240 mV, which are normally attributed to the formation energy of simple cooper vacancies, VCu -, in theoretical 1 and experimental 26 studies. This effect is more evident in the activation energy obtained from the cooling step, reaching a value of 273 mV for the film annealed at 450°C. As described before, in the Mg-doped Cu2O system, two different regimes are visible, with a temperature transition around 170°C. In the low temperature regime, the activation energy varies between 162mV and 176mV, while at higher temperature, the activation energy increases with the temperature, from 209mV to 250mV. This value is represented in Figure 27d) as Cu2O:Mg heating curve. The variation for temperatures higher than 170°C and labelled as Cu2O:Mg heating (2 nd energy) leads to energies comparable with the heating activation energy for the intrinsic case (~170 mV). For the energy obtained from the cooling curves, we observe lower activation energies, below 160mV, which decrease down to 117mV, after the annealing treatment at 450 °C. This value is lower than that of the simple copper vacancy and it is comparable with the secondary defect observed in Cu2O:Sr 3 with an ionization energy of 133mV, as well as acceptor levels detected in Nitrogen doped Cu2O at an energy of 121 meV 27 .

To complete the characterization of the electrical properties of the samples and to evaluate the effect of the annealing on the mobility and free charge carrier concentrations, the samples were analysed by ex situ Hall Effect. All samples showed p-type behaviour. As previously observed in the resistance measurements, the resistivity of the intrinsic Cu2O shows a minimum in the film annealed at 250°C, followed by a drastic increase up to 1000 Ω.cm. Concerning the 18% Mg doped film, the lowest point is visible at 450°C, with a value in the 10 Ω.cm range. due to high resistivity values. Concerning the mobility of Mg-doped Cu2O films an increase from 1 to 4 cm 2 .V -1 .s -1 with annealing temperature is observed; while in the undoped case, the mobility is maintained roughly constant around 9 cm 2 .V -1 .s -1 . The free charge carrier density, associated mainly to holes, varied with the two group of samples. On the one hand, the Cu2O showed a maximum of charge-carriers density of 6x10 16 cm -3 at 250°C, followed by a decrease when annealed at higher temperature. While on the other hand, the Mg-doped Cu2O films show a general increase of charge carries, presenting the higher value for the film annealed at 450°C, close to 2x10 17 cm -3 .

Discussion

The effect of magnesium incorporation on the stability of the Cu2O phase under different annealing treatments was studied in order to clarify the mechanism of CuO formation. Apart from the surface oxidation detected by XPS on samples as-deposited or annealed at 250°C, the temperature of the formation of CuO is similar, starting below 350 ºC, in both intrinsic and Mg doped films. This was confirmed by XRD, Raman, EXAFS and indirectly by resistance measurements. However, the formation of CuO occurs in a different way between the two sets of films. In intrinsic Cu2O films at high temperature, the formation of grains attributed to CuO is generalized at the surface of the film, while in the Mg-doped films the grains are only formed on a fraction of the surface. This leads to a difference of resistivity behaviour at 350 ºC; a significant reduction of resistivity is caused in the doped samples, while the intrinsic films become more resistive when the CuO phase is present. When the annealing temperature increases to 450°C, the undoped Cu2O films show a dominant content of CuO, increasing drastically its resistivity. In the opposite way, the magnesium-doped Cu2O films maintained a higher fraction of Cu2O combined with low resistivity values in the 10 Ω.cm range.

The different behaviours between the two sets of samples led us to deduce two distinct mechanisms for the CuO formation. On the one hand, the intrinsic Cu2O starts to form CuO by oxidation at the surface and interior of the film. At the surface the Cu + cations are easily oxidized into Cu 2+ in contact with oxygen forming new grains of CuO phase.

Though in the film, the formation of split copper vacancies induces the formation of Cu 2+ centres, which creates CuO nucleation centres inside the grains. In a first stage, the split copper vacancies leads to an inefficient generation of free holes due to the highly localized character of this defect, which later contributes for a large increase of resistivity by the formation of the parasitic phase throughout the film. The electric transport measurements confirm this hypothesis, due to the increase of resistivity during the annealing stage at 350°C.

On the other hand, in the presence of Mg, the transformation seems to start at the surface of the film, while in the interior of the film, simple copper vacancies are the major defects to be created, facilitated by the presence of the dopant. The large creation of simple copper vacancies increases the number of free charge carriers and consequently reduces the resistivity. As suggested by Isseroff and Carter, doping with Mg can prevent the formation of split vacancies for a single cation vacancy, since the divalent cation would assume a position similar to a split copper vacancy in the crystallographic structure 13 . This fact seems to be confirmed by the low resistivity of the films, even when the surface is covered by the CuO phase. Additionally, the XRD shift detected after the annealing at 450°C can suggest an activation of this mechanism by the changes in crystallographic structure. The inhibition to create split vacancies would reduce the parasitic centres inside the film, being the main mechanism for conduction loss during the annealing stages. Furthermore, the first activation energy associated to Cu2O:Mg of 165mV can be compared to the energy value of a large size impurity-vacancy complex found previously in the Sr-doped system 133mV 3 . One type of defect that can explain this energy is a clustering of a second simple vacancy with the dopant, [MgCu -2•VCu] -,which has been suggested by Isseroff and Carter 13 . The use of EXAFS by synchrotron radiation at ESRF in Grenoble was additionally fundamental to clarify the crystallographic changes with the presence of the dopant as well as the mechanism for phase transformation between Cu2O and CuO. Nevertheless the creation of a crystallographic model to fit the data is still ongoing. The combination of these stability results with the changes in conductivity previously discuss can provide a way to better apply this material both in solar cells and electronics, or even new copper oxidebased devices.

Cu 2 O:Mg/ZnO heterojunctions

Introduction

The lack of a stable n-type semiconductor Cu2O thin film, diminished drastically the possibility of solar cells production based on copper oxide homojunctions. This fact can be overcome by the use of a different n-type semiconductor that present a band alignment compatible with Cu2O to form a heterojunction. ZnO, as an n-type semiconductor oxide, can be combined with Cu2O, which in this case work as p-type absorber material, due to the lower band-gap. The maximum theoretical efficiency of a ZnO/Cu2O solar cells is as high as 20% 28 , however the maximum value achieved by functional devices is only 4.12% 29 for an intrinsic ZnO/Cu2O junction [2], which was recently increase to 8.1% in a MgF2/ZnO:Al/Zn0.38Ge0.62O/Cu2O:Na heterojunction 30 . Besides, this record of efficiency was achieved by thermal oxidized Cu2O, a process that creates large grains in the tens of microns range, reducing the fraction of grain boundaries and increasing the mobility.

However, this technique is not suitable for large scale production, since it requires very high temperatures, around 1000 ºC 31 . Therefore, the use of chemical deposition techniques for the deposition of Cu2O combined with the possibility of depositing ZnO at low temperature pave the way for future implementation of a cost-efficient oxide based solar cells.

The Mg-doped Cu2O study, previously discussed, provided two essential results that can have a significant impact in the electrical behaviour of pn junctions. On the one hand, the presence of the cation contributed to a higher stability of the Cu2O phase, reducing the appearance of CuO parasitic phase. This phase is reported to reduce the performance of an oxide-based solar cell, due to the creation of highly conductive paths in the grains boundaries that short-circuit the device 32 .

On the other hand, the Cu2O:Mg thin films showed a high charge-carrier density, up to 8x10 17 cm -3 , due to a doping mechanism based on copper vacancy assisted cation incorporation. The larger density of charge carriers (holes) can have an important contribution to an increase of efficiency on a Cu2O based solar cell, by reducing the size of the depletion region 33 . This enables a higher electric field and charge collection throughout the cell, by reducing the number of charges lost due to recombination and higher reflection in the back electrode 33 .

As proof of concept, we first simulated a pn junction of intrinsic and Mg-doped Cu2O on n-type ZnO in a semi-classical approach using numerical simulations based on a finite-difference scheme. For the experimental part, we fabricated these structures by combining the Mg-doped Cu2O thin films with ZnO, deposited by pulsed injection metal organic chemical vapour deposition (PI-MOCVD) and by spatial atomic layer deposition (SALD), in order to study the rectifying response of these structures. More specifically, two different structures were studied, pn and np, in order to understand the electric and optical impact on the device. The study is completed by photovoltaic tests to evaluate the possible application of these structures into solar cells.

Simulation of a pn junction formed by Cu 2 O/ZnO

Formerly to the fabrication of pn junctions experimentally, we decided to simulate the band structure and consequent diode response of these two materials. The simulations of this semiconductor device were based on a mathematical model, using the so-called basic semiconductor equations. The semiconductors must be non-degenerated, with a general behaviour, without any particular assumptions, derived from Maxwell-Boltzmann statistical relations and Maxwell's equations. A numerical solution of the system of basic semiconductor equations is then obtained through a finite-differences method. The simulation was a semi-classical approach to the problem, where we took into account the conservation of charge and mass 34 . The equations are solved in the steady-state regime where the applied voltage is constant, which additionally allows to solve the system in reverse and forwards bias situations.

The group of parameters selected for the two semiconductors were based other works 35 , 36 , properties of samples fabricated and measured in our labs and general assumptions for pn junctions. We assumed a thickness between 200 nm and 1 m for each film, with values for band-gap energy, mobility, permittivity and effective mass referred in the literature 35 , 36 . The concentration of defects, both in the interface and trap states was assumed as minimum at a level of 10 10 cm -3 , to avoid significant contributions to the simulation, while the capture cross section and radiative recombination coefficient were based on pn junctions formed by group IV semiconductors 37 . Therefore, in Table IV-10, all the different fixed parameters for the simulation are represented. Defects traps (cm -3 ) 1x10 10 1x10 10

Mobility electrons (cm 2 .V -1 .s -1 ) 1 10

Mobility holes (cm 2 .V -1 .s -1 ) 1 10

Relative permittivity 18,1 10,8

Electrons effective mass 0,99 0,3

Holes effective mass 0,58 0,5

Capture cross section (cm 2 ) 1x10 -12 1x10 12 Radiative recombination coefficient (cm 3 .s -1 ) 1x10 -11 1x10 -11 Here, we discuss two main simulations where the varied parameters are the acceptors concentrations in the p-type semiconductors to simulate intrinsic and doped conditions for Cu2O. Considering the previous deposited Cu2O:Mg films, we defined the density of acceptors to be from 5x10 15 cm -3 to 8x10 18 cm -3 in intrinsic and doped cases, respectively. For ZnO, we used well establishing values from materials previously obtained in Grenoble LMGP 38 , around 10 19 cm -3 38 . Additionally, the size of the simulated device was adapted to the width of the depletion region between the two cases. These parameters are represented in Table IV

(IV.2)
where J is the current density, J0 the reverse current density, q the electron charge, V the applied voltage, A the contact area, RS the series resistance, n the non-ideality factor, kb the Boltzmann's constant and T the absolute temperature.

Even if the slope are the same and represent a non-ideality factor of 1, it is important to refer the differences in series resistance. In the Cu2O case, we observe a change in slope at higher applied voltages, caused by the higher series resistance of the Cu2O layer. This fact is originated by the lower density of charge carriers in the oxide and the large width of this junction, which results in an increase of the total resistance. (2017) 38 . The use of different thin films staking order was constructed to understand the impact of ZnO as absorbing layer. Since the n-type film presents a higher band gap of 3.4

eV in the UV range, the change of order under solar illumination could increase the light absorption, 1 st structure, which can result in higher efficiency of the cell.

The use of two separate technique is then justified by the temperatures required to deposit each film. As ZnO by MOCVD was to be deposited at 400°C, it is not possible to process it after Cu2O due to formation of CuO, previously discussed. Additionally, ZnO is stable at 350°C, temperature used for Cu2O deposition, so we would preserve the general properties of the film. Nevertheless, it is important to report the possible slight changes in charges carries concentrations or mobility of ZnO during heating before the Cu2O deposition. The ZnO deposited by SALD allow the use of lower temperatures, 200°C, where Cu2O is still stable, therefore the inverse structure could be fabricated. The annealing was performed at 250°C for 2 hours, in order to avoid any CuO formation.

In Figure IV-90, the diode response between -1V and 1V is clearly visible. In terms of nonideality factor, there were no significant variations in any diodes. Nevertheless after annealing there was an increase of current in forward and reverse sense. This fact may be attributed to the high concentration of holes in Cu2O, which leads to a lower sheet resistance.

lines). ). Current density presented in absolute value

The diode parameters were extracted by fitting the experimental curve with the typical diode equation, considering the series and shunt resistances as follows:

𝐽 = 𝐽 0 (𝑒 𝑞(𝑉-𝐽𝑅 𝑠 ) 𝑛𝑘 𝐵 𝑇 -1) + 𝑉 -𝐽𝑅 𝑆 𝑅 𝑆ℎ (IV.3)
where J is the current density, J0 is the reverse current density, q is the electron charge, V is the applied voltage, RS is the series resistance, n is the non-ideality factor, kB is the Boltzmann's constant, T is the absolute temperature, and RSh is the shunt resistance. In the case of the first structure ZnO/Cu2O, RS, RSh, and J0 could not be extracted by the fitting to the implicit equation. Only n was determined by the ideal diode equation. The fitted parameters of the different experimental diodes are shown in Table IV-13. The comparison of all diodes led us to conclude that the ones where SALD ZnO is deposited on top of Cu2O, show better parameters in general. The non-ideality factor obtained is as low as 2.0, with a rectifying ratio (RR±1V) above 100 for both intrinsic and Mg-doped Cu2O. Nevertheless, even if the series resistance is low compared to the ZnO MOCVD cases, being favourable for a diode, the shunt resistance is also lower, which increases the possibility of shortcircuit and reduces the global performance of the device. These experimental results are difficultly comparable to the simulated junctions, especially due to diverse rectifying behaviour, and the current density values obtained.

Nevertheless, the effect of a more conductive Cu2O is the similar in both cases, leading to a reduction of the sheet resistance. This fact is particularly visible in the ZnOMOCVD/Cu2O:Mg heterojunctions, since a reduction of sheet resistance is clear after the annealing treatment at 250°C.

Cu2O/ZnO diodes from other studies show similar results in terms of non-ideality factor and rectifying ration. Heterojunction of Cu2O and ZnO deposited by RF sputtering showed in dark IV curves with an non-ideality factor larger than 2.0, reaching up to 3.8 with annealing treatments 39 . These results are similar to the ones here reported, mainly the1 st structure. In other work between these two oxides deposed by RF magnetron sputtering 40 , the ratio between the current at -1V and +1V is marginally higher than 100, comparable to the best diodes reported in our work.

Photovoltaic performance of the pn junctions

The goal of these structures is the implementation into solar cells. Therefore photovoltaic performances were analysed under AM 1.5G illumination conditions. In In the SALD cases, the current density values increased around 3 orders of magnitude for both p-type layers, when compared to the previous case. The Cu2O cell shows a current density of 0.79 mA.cm -2 , the highest measured in these experimental results. In the case of Mg doped Cu2O, the VOC is the highest found, 379 mV, with a reduced JSC of 0.16 mA.cm -2 . Overall, even with the presence of photovoltaic effect, the efficiency of the cells produced are unable to pass the 0.1% mark. In the two cells prepared by SALD, we can extract the highest values of JSC and VOC for us obtained, 0.8 mA.cm-2 and 380mV, respectively. Considering the physical and chemical deposition techniques of thin films, the values of photovoltaic performance are in the range of other studies on Cu2O. In a study by H Akinaga et al. 39 , the deposition by magneton sputtering of Cu2O and ZnO thin films led to the formation of a solar cell with 0.4mA.cm -2 and 300 mV of JSC and VOC, respectively. The thickness of the oxide layers is also in the hundreds of nanometres, as in our study, however, the metals used for contacts are different, based on different layers of Pt, Al and Au. Other study by S. Jeong et al. 41 , Cu2O thin films (280nm) were deposited by MOCVD on ZnO thin films (130nm) in this case deposited by RF sputtering, presenting a comprehensive analysis on the temperature dependence of the PV parameters. The direct comparison of the PV response at room temperature shows VOC and JSC values of 300mV and 1 mA.cm -2 . They attributed these meagre PV results to an interface recombination mechanism that is the dominant carrier transport mechanism in these solar cells. 41 .

The use of spray-pyrolysis to deposit Cu2O was also found in the literature, however in this case combined with TiO2, as n-type oxide layer 42 . The open circuit voltage was reported to be up to 350 mV, while the short circuit current density reached a maximum of 0.4 mA.cm -2 . From the same group was also reported in 2016 a study with the incorporation of Mg in Cu2O, deposited also by spray-pyrolysis 43 . In this case, the current density value reached up to 0.9 mA.cm -2 , with a VOC in the same 350 mV range 43 . The JSC improvement on the solar cells was attributed to a higher photo-conductivity of cuprous oxide with the Mg incorporation. ZnO:Al presented a high leakage current. Moreover, the junctions could only be used as photodiodes as no electrical power could be harnessed upon solar illumination.

These previous reports using sputtering or chemical vapour techniques are drastically supplanted by two other techniques for the growth of Cu2O: electrodeposition and thermal oxidation. Both techniques enable the growth of Cu2O with mobility values higher than 10 cm 2 .V -1 .s -1 . In the case of electrodeposition of Cu2O in combination with ZnO, several studies report short-current density values higher than 1 mA cm -2 44, 45, 46 . A study by T. Buonassisi and R.G. Gordon 47 reports efficiency values as high as 2.85%, with VOC and JSC of 622 mV and 7.25 mA cm -2 , respectively. Finally, the use of thermal oxidation of copper sheets is at this moment the best technique to obtain Cu2O base solar cells. Various studies by T. Minami 30,31,[48][49][50][51][52] show the highest values in terms of photovoltaic parameters, with short-current density of 10 mA.cm -2 49 , open-circuit voltage of 1.1V 30 and efficiency of 8.1% 30 , using sodium-doped Cu2O.

Future studies on these pn junction could also be performed to access the applicability of Cu2O/ZnO structures for photocatalysis in water splitting processes.

Summary

The integration of Cu2O, intrinsic and Mg-doped, into pn junctions was successfully achieved using only chemical vapour deposition approaches. Rectifying behaviour was observed in the junctions created, even with significantly more modest performance when compared to numerically-simulated devices. The Mg doping effect was especially visible in films deposited on MOCVD-grown ZnO, reducing the leakage current on the device.

When ZnO was the last oxide to be deposited by SALD, both intrinsic and doped films showed similar electrical responses.

The idea of this approach was to enable fast, scalable and low-cost technology aiming for full oxide-based solar cells. Nevertheless, photovoltaic performances presented here are obviously far for the best literature devices. The use of AA-MOCVD to produce Cu2O limits strongly the mobility of these thin films, more than one order of magnitude below the thermally-oxidized Cu2O films. The impact of this low mobility is strongly visible on the current density values and as consequence the fill factor and efficiency.

Additionally, an increase of current density in the diodes annealed at 250 °C, did not lead to improvements in PV performance, as proposed.

Conclusions

A comprehensive study on the impact of magnesium incorporation in cuprous oxide thin films was developed in this chapter. The presence of the dopant was confirmed up to 17% concentration using aerosol-assisted metal-organic chemical vapour deposition under atmospheric pressure condition at 350ºC. As main influences, the films showed a higher roughness, a lower electrical resistivity (6.6 .cm), concentration of holes up to 8.1x10 17 cm -3 and a higher stability of the Cu2O phase under oxidizing environment. The magnesium presence seems to change the mechanism of copper vacancies generation, blocking the formation of split copper vacancies. This leads to a reduce transformation of Cu2O into CuO.

5 Chapter V: CuCrO 2 delafossite thin films

Deposition and characterization of CuCrO2 films

Introduction

The lack of improvements in transparency of Mg-doped Cu2O thin films lead us to test an alternative approach, the p-type delafossite phase CuCrO2. The intrinsic higher band-gap of 2.8eV 1 and the lower resistivity of 0.1 Ω.cm 2 , when compared to Cu2O, can enable the application of this material into transparent semiconductor devices. A detailed presentation of the state of art of this material can be found in section I.1.4. In the study here presented, we optimize the deposition of CuCrO2 thin films by AA-MOCVCVD. The deposition conditions were based on a particular study by Farrell et al. 3 , which reports the growth of undoped CuCrO2 thin films by spray pyrolysis. This work showed the nanocrystallization of the delafossite phase at temperatures below 400 °C, without any postannealing treatments. Interestingly, the metal organic precursors and solvent used by Farrell et al. 3 are compatible with aerosol system in LMGP, based on a piezoelectric transducer.

Experimental

Using the AA-CVD reactor in configuration 2 previously described in Chapter II, the experimental parameters of the solution and deposition are presented in Table V-15.The precursors used were copper acetylacetonate, Cu(acac)2, and chromium acetylacetonate, Cr(acac)3, with ratio between Cu and the total concentration of cations (Cu/Cu+Cr) varied from 40% to 60% for a total concentration of precursors of 0.02 mol.l -1 in ethanol. For this study, the time was increased to 90 minutes and the temperature to 400 °C in order to reproduce literature study 4 . Argon with a flow rate of 6 l.min -1 was used as carrier gas, while O2 with a flow rate of 1.5 l.min -1 was used as reaction gas as in Cu2O case. These flow-rate values allowed us to maintain a solution consumption rate of 2 ml.min -1 , resulting in a total solution consumption of 180 ml.

Characterization of CuCrO 2 films

The CuCrO2 films present an optical appearance completely different to Cu2O films, as presented in the amorphous or nano-crystalline phase. For the films with a higher Cu content, 46% and 56%, the XRD spectra is marked by the presence of the (012) and (110), diffraction peaks with various intensity at 36.4° and 62.4°, respectively. These diffraction peaks are attributed to the rhombohedral delafossite phase according to the 04-010-3330 file of the International Center for Diffraction Data (ICDD). The relatively small intensity and broad width of these peaks may reveal a low degree of crystallinity of the CuCrO2 thin film. This can be attributed to the deposition temperature of 400 °C, which is typically close to the lowest crystallization temperature of the delafossite phase 7 . The crystallite size is estimated to be between 18 and 20nm in the two samples, as deduced from the (012) peak at 36.4° using the Scherrer's equation 8 . and 3b, respectively, the oxygen atoms in special position 6c. According to the method of nuclear site group analysis 10 and after subtraction of the three acoustic modes (A2u + Eu) of the total representation of the crystal, the irreducible representations associated to the optical modes can be expressed as:

Γ 𝑜𝑝𝑡 = 𝐴 1𝑔 + 𝐸 𝑔 + 2 𝐴 2𝑢 + 2 𝐸 𝑢 (V. 1)
where A1g and Eg are Raman-active modes and A2u and Eu modes are infrared (IR)-active modes. It should be noted that vibrational modes coming from A and B atoms are IR-active only and have no contribution to the Raman spectrum, which means that only oxygen atoms move in the Raman modes. The Raman spectra obtained with a wavelength of 488 nm are presented in Figure V-99. In coherence with the XRD results, the spectrum obtained from the film with the lowest Cu content does not present well-defined Raman peaks. In the two spectra registered from the films with higher Cu content, two Raman modes at 457 and 707 cm -1 are detected that can be assigned to the modes Eg and A1g 11 of the CuCrO2 phase, respectively. The first mode, weakly observed on the decreasing tail of the Rayleigh line at 107 cm -1 , corresponds to an infrared-active mode Eu of the CuCrO2 phase 12 . Other additional modes of weak intensity particularly a in the 520-650 cm -1 , also occur in the CuCrO2 thin film spectrum. Such additional modes, can be assigned to defect-induced modes, have already been reported at 536 and 586 cm - 1 11 , and 538, 557, 623 and 668 cm - This fact is common of CuCrO2 films, due to high reflectance 16 . Comparing with Cu2O films, the transmittance is slightly higher, especially for the 46% Cu content film. In terms of absorption edge, we calculated the direct band-gap for the different films, represented in The origin of this lower resistivity was then investigated by Hall Effect measurements using a Van der Pauw configuration. Unfortunately, the hall coefficient sign was unable to be determined. As reported by the literature 6 , the mobility values of CuCrO2 are below 1cm 2 .V -1 .s -1 , which are under the limit of the Hall Effect system used.

Considering the mobility values in a 0.1cm 2 .V -1 .s -1 range, the density of free holes would vary from 10 18 cm -3 in the 24% Cu film and up to 10 21 cm -3 in the richest Cu content film.

Conclusion on CuCrO 2 thin films

The growth of CuCrO2 thin films by AA-MOCVD was achieved with noteworthy results in terms of transparency and resistivity, comparable with the literature 4,5 . The films prepare with an equimolar solution of both precursor showed a stoichiometry of 46% of Cu content, close to the ideal ratio for CuCrO2. They present a good compromise of optical and electrical properties, with a transmittance of 60% and resistivity values close to 10 Ω.cm.

This film ratio was then selected to be tested for devices application. The combination of CurCrO2 with ZnO has allowed to develop UV-sensor also discussed in section I.2.7. Nevertheless, nanostructured heterojunctions consisting in core-shell heterostructure based on ZnO nanowires (NW) have been ever reported so far. The use of nanowire arrays can improve the performance of a photo-electrical device, due to an increase of surface-to-volume ratio 17 , enhanced light trapping processes [18][19][20] , efficient charge carrier separation and collection over the nanowire, caused by the high mobility electron pathways in the monocrystalline NW 18,21 . In this work, we present the first fabrication of a ZnO/CuCrO2 core-shell NW heterojunction using surface scalable chemical deposition techniques, and its integration into a self-powered UV photodetector.

ZnO/CuCrO

Experimental

The growth of ZnO NW with controlled length and diameter has been developed and studied at LMGP in previous works 22 . ZnO NW arrays were grown by chemical bath deposition (CBD) on commercial 130 nm-thick ITO /glass substrates (Delta Technologies).

The ITO/glass substrates were initially cleaned in an ultrasonic bath of acetone and then in a beaker containing isopropanol for 15 min each to remove any contaminants. Then, a ZnO seed layer was deposited onto the ITO thin layer by dip coating based on sol-gel reactions. °C. The solution consumption rate of 2 ml.min -1 , resulting in a total solution consumption of 120 ml. Circular silver contacts with a diameter of 0.5 mm defining the size of the selfpowered UV photodetector were deposited by thermal evaporation on top of the ZnO / CuCrO2 core-shell NW heterostructures.

Structural characterization of core-shell structures

The observe a similar spectrum to the previous section (V.1), with Raman, IR-active and defect induced modes visible in the spectrum. In the ZnO NW case, the correlated ZnO Raman lines occur at 99 cm -1 (E2 low ), 334 cm -1 (E2 high -E2 low ), 439 cm -1 (E2 high ), 587 cm -1 (E1(LO)

) low ), 1106 cm -1 (2LO), and 1160 cm -1 (2A1(LO), E1(LO); 2LO) 26 . The remaining Raman peaks as identified by the dark open circles are attributed to the ITO thin layer and its correlated phonon modes at 591 and 1102 cm -1 also occur in the corresponding Raman spectrum. On the core-shell NW structure, the Raman lines at 458 cm -1 (Eg), 709 cm -1 (A1g), and the additional defect-induced modes, particularly those at ~106, 207 and 754 cm -1 , are direct evidence of the CuCrO2 shell crystallized in the delafossite phase 3,12,11 . Moreover, The rectifying ratio between the forward current density to the reverse current density at ±1 V is as high as 5500, which is considerably larger than the ratio of 70 at ±1.5

Electrical characterization of core-shell heterostructures

V for the diode made of ZnO/CuCrO2 thin films as reported by Tonooka et al. 30 . The JV tests also performed on planar films showed a lower diode response, pictured in Figure V-110. In this case, the diode behavior was detected for higher applied voltages, larger than 2V, with significantly lower rectifying ratio, around 20 at ±5 V. where J is the current density, J0 is the reverse current density, q is the electron charge, V is the applied voltage, A is the contact area, RS is the series resistance, n is the non-ideality factor, kB is the Boltzmann's constant, T is the absolute temperature, and RSh is the shunt resistance. The non-ideality factor n of the diode is found to be 2.2, indicating a good diode behaviour. The relatively large shunt resistance of 6.5 kΩ is correlated with the small reverse current density at -1 V, which can be attributed to the conformal CuCrO2 shell and to the high mobility in the ZnO NW. Moreover, the series resistance of 0.5 Ω originates from the quality of the metallic contact and from the high conductivity of the CuCrO2 shell 32 . Overall, the fitted values of J0, n, RS and RSh as well as the rectification ratio of the ZnO/CuCrO2 core-shell NW heterostructures reveal the high quality of the pn An absorption higher than 85 % below 400 nm is obtained in the UV region. This is due to the wide band gap energies of both oxides, which are higher than 2.8 eV, and to the antireflective character of the nanostructured NW array 33 . The responsivity of the selfpowered UV photodetector (i.e., the photo-generated current per unit incident light power)

shows only one peak centred around 374 nm, which very well correlates with the 3.37 eV band gap energy of ZnO and confirms that the photo-generation of charge carriers which demonstrate that this UV photodetector offers self-powering capabilities. The higher specific surface area of the heterojunction is expected to enhance light trapping processes in the ZnO NW arrays, as well as to exhibit a larger total volume of depletion region and longer carrier lifetime 36 . However, the VOC of 16 mV is fairly low, especially when compared to other delafossite phase material like CuAlO2, which presents a VOC of 250 mV when combined with ZnO NWs 37 . The use of a CuAlO2 layer filling the ZnO NW arrays and thus planarizing the electrical contact may account for the improvement of the PV performances. The reduced PV performances of our device may also be explained by the CuCrO2 properties themselves, especially the anisotropy of its electronic properties 38 , and its typical low hole mobility. The present findings reveal that an all-oxide core-shell NW heterotructure grown by chemical deposition techniques appear as a promising candidate as self-powered UV photodetectors.

Conclusions on the core-shell heterostructures

An original core-shell heterostructure made of ZnO NWs covered with a conformal CuCrO2 shell was successfully achieved through a low-cost surface scalable chemical fabrication route. The structural morphology of the CuCrO2 grains shows a columnar grain growth at the top of the ZnO NWs, while it consists of smaller nano-grains on their vertical sidewalls. Although no epitaxial relationships occur as revealed by ASTAR maps, the different grain morphology obtained on the different ZnO faces implies distinct growth mechanisms taking place in each case. This heterostructure successfully works as a diode, with a rectification ratio as high as 5500 at ±1 V, which is much better than similar planar diodes. It further exhibits a high absorption above 85% in the UV region. Moreover, the device presents a maximum responsivity of 187 µA/W under zero bias at 374 nm and a high UV-to-visible rejection ratio (374-550 nm) of 68 together with a PV effect under standard AM1.5G illumination. The combination of the high rectification ratio, of the high absorption in the UV region, and of the significant responsivity offers promising possibilities for the use of this heterostructure as a self-powered UV photodetector.

Eventually, the use of CBD and AA-CVD at temperatures below 500 °C is favorable for the implementation of all oxide-based self-powered UV photodetectors with reduced costs and large scale production. The combined finding about the material stability and the higher conductivity of As the increase of transparency of Cu2O could not be reached in the previously described work, other copper-based oxide systems with a larger band-gap were considered.

Therefore, the p-type semiconductor alloy CuCrO 2 was developed and thin films were successfully deposited by aerosol-assisted metal-organic chemical vapour deposition.

These materials exhibit optical and electrical properties similar to state-of-the-art reports, with a direct band-gap of 3.2eV and resistivity values of 0.1 Ω.cm.

As a consequence, this delafossite material was implemented into an original coreshell heterostructure in combination with ZnO towards application as a UV photodetector. This heterostructure functionally operates as a diode, with a rectification ratio as high as 5500 at ±1 V, denoting larger improvements compared to diodes with similar junctions under planar architecture, which present rectification ratio below 100. It further exhibits optical absorption above 85% in the UV region.

RésuméL

  'absence d'oxydes semi-conducteurs de type p de haute performance retarde le développement de d'électronique transparente et du photovoltaïque à base d'oxydes. Dans le groupe des composés semi-conducteurs, les oxydes à base de cuivre présentent des caractéristiques électriques, optiques et de fabrication prometteuses qui établissent cette famille de matériaux comme bien adaptés aux applications semi-conductrices de type p. Dans ce travail, nous nous concentrons sur la croissance de films minces d'une part de Cu2O dopée par des cations et d'autre part de CuCrO2, visant à améliorer leurs propriétés optiques et électriques. De plus, nous avons mis en oeuvre ces films d'oxyde dans des dispositifs de jonction pn tels que des cellules solaires et des photodétecteurs UV.

  Preface ................................................................................................................

1. 1

 1 Cuprous oxide (Cu2O) ................................................................................ 1.1.1 Conductivity in Cu2O ........................................................................... 1.1.2 Cu2O deposition ................................................................................... 1.1.3 Dopants............................................................................................... 1.1.4 CuCrO2 -Delafossite .......................................................................... 1.2 Devices based on oxide semiconductors .................................................. 1.2.1 Transparent electronics ...................................................................... 1.2.2 Photovoltaics ...................................................................................... 1.3 Conclusion ................................................................................................ 1.4 References ................................................................................................

  represented in the Figure I-1 a).

Figure I- 1

 1 Figure I-1 Crystallographic structure of Cu 2 O with oxygen atoms as the origin of the bcc cell. Oxygen represented in red and Copper in blue. (b) Cu 2 O atomic structure of 2x2x2 cells. The two interpenetrating Cu-O-Cu networks (one light coloured and the other dark coloured). Adapted from Nolan 2

  . The following Figure I-2 represents the bands structured of Cu2O, where Γ 6 + and Γ 7 + are the conduction and valence band, respectively.

Figure I- 2

 2 Figure I-2 Band structure of Cu 2 O. Adapted from Brochen et al 6

Figure I- 3

 3 Figure I-3 Band schematic of Cu or Ag based oxides. a) chemical bond between an oxide ion and a cation with closed metal d shell and b) Schematic of where the energy levels of M d 10 and O 2p 6 are assumed

Figure I- 4

 4 Figure I-4 Structure of a single Cu vacancy in Cu 2 O. Complete structure, simple vacancy and split vacancy. Copper atoms pictured as blue circles and oxygen as red ones.

Figure I- 5

 5 Figure I-5 Number of publications on "Cu 2 O", "doped Cu 2 O", "copper oxide", "cuprous oxide" and "copper (I) oxide" are plotted from 2000 to 2017, as obtained from Scopus.

Figure I- 6

 6 Figure I-6 Temperature pressure diagram from bulk Cu-Cu 2 O-CuO crystals. Adapted from Schmidt-Whitley 29

Figure I- 7 .

 7 The film surface changed from smooth to granular with the increase of Sr concentration. On the electrical properties, Cu2O thin films containing 5-6 %at of Sr showed resistivity values of about 10 Ω.cm, visible in Figure I-7I, which were then decrease to 1 Ω.cm with thermal annealing treatments at 200 °C in air. This films show hall carrier density and mobility values of 2.8 × 10 17 cm -3 and 15 cm 2 .V -1 .s -1 , respectively, as reported in a follow-up study by Brochen et al 6 . Nevertheless, no significant effect on the band gap or optical properties was visible, as shown in Figure I-7II.

Figure I- 7 I

 7 Figure I-7 I) Morphology, resistivity and II) Transmittance of Cu 2 O:Sr thin films from Bergerot et

  The crystallographic structure of the CuCrO2 delafossite phase is represented in Figure I-8a), while the number of publications of CuCrO2 in Scopus since 2000 to 2017 is plotted in Figure I-8b).

Figure I- 8

 8 Figure I-8 a) CuCrO 2 crystalline structure. and b) number of publications of CuCrO 2 in Scopus since 2000 to 2017

  source, drain and gate, which are related to a junction transistor as emitter, collector and base, respectively. This type of transistor differs from a MOSFET in the stacking organization of the n-type or p-type component as visible in the Figure I-9. In the bottom gate approach here explained, the metal gate in deposited on top of the substrate, forming the first layer of the device. Then, a dielectric material covers the metal gate, working as isolating layer. The semiconductor material is deposited on top of this isolating layer, being referred as body. The source and gate composed of metal material are then deposited on two opposed edges of the semiconductor body. This configuration, visible in FigureI-9, represents a bottom gate field-effect transistor. Other configurations are also possible as top-gate with staggered or coplanar semiconductor bodies.

Figure I- 9

 9 Figure I-9 a) Field-effect transistor FET scheme; b) Characteristic curve of a TFT with current at the drain in function of voltage applied between source and drain; c) Image of transparent TFT. Adapted from E.Fortunato73

Figure

  Figure I-10c, where the grains show a homogeneous distribution with a grain size around 100 nm. The integration of the films in a TFT transistor, with Ni contacts, show the possibility of creation a transistor with an on/off current's ration of 1.1× 10 4 and a V threshold of -7.5V (Figure I-10d) 81 .

Figure I- 10

 10 Figure I-10 Transfer curves of different TFT based on Cu 2 O produced by different groups a) Matsuzaki et al. 8 b) Xiao Zou et al. 80 c) and d) Joonsung Sohn et al. 81 where c) SEM micrograph of a Cu 2 O film annealed at 500 °C

  . Each layer had a different effect on the TFT, since the Cu2O film worked as a capping layer, while the SnO on the bottom controlled the upper layer stoichiometry. The final 25 nm bilayered structure presented in Figure I-11b), showed a p-type channel performance with a final onto off current ratio of 1.5x10 2 and a threshold voltage of -5.2V.

Figure I- 11

 11 Figure I-11 Transfer curves of different TFT based on Cu 2 O produced by Fortunato et al. 83 and b) Bilayer FTF of Cu 2 O and SnO produced by Hala A. Al-Jawhari et al . 85

Figure I- 12 .

 12 Figure I-12. As illustrated below, W0 depends on the carrier concentration of the semiconductors and thus the n and p-type materials can contribute differently to the depletion region. In Figure I-12 for instance, the p-type semiconductor contributes more to W0 since it has a lower carrier concentration than the n-type semiconductor.

Figure I- 12

 12 Figure I-12 Band structure of an inorganic pn junction with zero bias.

Figure I- 13

 13 Figure I-13 a) Forward bias decreases depletion zone and built-in potential; b) Reverse bias increases depletion zone and built-in potential.

  Figure I-14, the open and close circuit schemes of the pn junction are represented.

Figure I- 14

 14 Figure I-14 a) Carrier movement in pn junction under illumination with no load; b) Carrier movement in p-n junction with load.

  the recombination current, opposite to the electric current generated in the cell. When the current is annulled, the recombination and photocurrent are equal, and the voltage applied is called the open circuit voltage VOC. This voltage depends on the p-and n-type properties of the materials and the contacts between the electrodes and the semiconductors on the cell.

Figure I-

  Figure I-15a) represents the variation of the produced power with the bias voltage, while the Figure I-15b) shows a typical IV curve.

Figure I- 15

 15 Figure I-15 -a) Power variation with bias voltage b) IV curve of a solar cell

  semiconductors based on earth abundant elements, mainly Si, Fe, Cu, O, S, and their theoretical capacity of producing enough energy to cover the demand 93 . In Figure I-16 the annual electricity potential of each material in shown, regarding the materials extraction costs. The graphs shows that Cu2O has higher potential and lower extraction costs than silicon, therefore being proposed as an interesting energy material.

Figure I- 16

 16 Figure I-16 Annual electricity potential vs. materials extraction cost for several semiconductors.

m, as opposed to the standard 3

 3 m for fully electrodeposited devices. Apart from enhanced electric field, the use of a thinner Cu2O layer and the nanoparticulated AP-SALD Cu2O layer also contributed to a more effective collection of charges generated by photons of wavelengths above 450 nm. The obtained VOC was lower than expected, due to the nanometric nature of the AP-SALD Cu2O films, from 50 nm to ~250 nm, in Figure I-17, which is believed to cause charge recombination.

Figure I- 17

 17 Figure I-17 a) Cross section SEM images of an SALD-enhanced cell; b) EQE spectra of a thin (∼75 nm) and thick (∼200 nm) SALD enhanced cell (ZnO/Cu 2 OED/Cu 2 OSALD) compared to a 1 μm-Cu 2 O and 3 μm-Cu 2 O fully-ED (ZnO/Cu 2 O) cell 103

Figure II- 18

 18 Figure II-18 Droplets in a mist generated vibration 8

  Figure II-20 shows the process of nucleation and growth of the film.

Figure II- 20

 20 Figure II-20 Aerosol process at the substrate representing the nucleation and growth of the film.

  acetylacetonate anion (CH3COCHCOCH3 -) which act as a bidendate ligand, sharing 2 chemical bonding for one shared electronic charge. Normally it is abbreviated as acac, with the copper case being commonly referred as Cu(acac)2. As visible in Figure II-21, the copper atom is connected to four oxygen atoms, which form 2 six-membered chelate ring.

Figure II- 21

 21 Figure II-21 Schematic representation of Cu(acac) 2 with copper atom represented in blue, oxygen in red, carbon in grey and hydrogen in white 9 .

Figure II- 22

 22 Figure II-22 a) and b) shows the appearance of the precursors used, Cu(acac)2 and Mg(acac)2 for the doped case. Cr(acac)2 was a similar appearance to Cu(acac)2 with a dark blue colour, however with larger sized particles. In Figure II-22 c) is shown the solution with ethanol as solvent and before the addition of ethylendiamine, while Figure II-22 d) depicts the stirring step, already with ethylendiamine as additive.

Figure II- 22

 22 Figure II-22 Precursor used in the solution a) Cu(acac) 2 and b) Mg(acac) 2 ; solution after the addition of the solvent; d) stirring step in the magnetic plate.

Figure II- 23

 23 Figure II-23 Metallic holder with Corning glass, Silicon wafer and glass micro-slide as substrates in a 2.5cm x 7.5cm holder

  visible in Figure II-24b) and the temperature starts to increase, controlled by a thermocouple. The scheme of deposition set-up is also visible in Figure II-24a).

Figure II- 24

 24 Figure II-24 a) Scheme of deposition set-up and b) heating plate with substrates supported by a metallic holder

Figure II- 25

 25 Figure II-25 a) solution vessel with the level of solution visible; b) piezoelectric transducer and created mist.

Figure II- 26 a

 26 ). The mixture consisting of the mist and the gases is transported to a home-made reactor, beginning the deposition on the substrates, placed in a vertical position, FigureII-26 c). In the configuration presented here, the mist enters preferably into the chamber from the side, which leads to a horizontal flow, however, a vertical flow is also possible. The pressure inside the chamber and consequent exhaustion of the mist is controlled by an external pump connected to a liquid nitrogen trap to collect the residual solution. The barometer that controls the exhaustion is visible in FigureII-26b).

Figure II- 26

 26 Figure II-26 a) Transport of mist, with the inlet of Argon and Oxygen highlighted; b) barometer for control of exhaustion; c) window of chamber with visible inlet tube, mist and heating plate.

Figure II- 27

 27 Figure II-27 Vertical configuration of the reaction; a) scheme with all the components and b) photograph

  Figure II-28 show both setups used.

Figure II- 28

 28 Figure II-28 a) Box furnace used for air annealing treatments; b) 2-probe stage used for annealing treatments with in situ resistance measurements

  evaporator and the crucible are visible in Figure II-29.

Figure II- 29

 29 Figure II-29 Thermal-evaporator for metals evaporation with b) molybdenum crucible

Figure II- 30

 30 Figure II-30 a) Colour chart for refracted light on Si 3 N 4 , comparable to Cu 2 O thin films using Brigham Young University Silicon Nitride colour simulator 15 ; b) Optical microscope setup -Leica DM/LM

  . The incident electron beam is focused on the surface of the sample by a system of electronic lenses, as visible in Figure II-31 a). The interaction between the beam and the surface is mainly dependent on the energy of the beam and the atomic number of the atoms in the samples, resulting in different types of electrons emission or photons radiation by the probed sample. The emitted entities can be X-rays, Auger electrons, primary backscattered electrons or secondary electrons, as represented in Figure II-31 b). The detection of secondary scattered electrons allows the creation of a scanned image with different brightness levels, reflecting the roughness differences of the specimen, determined by the amount of electrons detected. In direct comparison with optical microscope, SEM enables spatial resolution of the order of nanometres, while the visible-light based microscopy only allows resolution higher than 500nm, limited by the diffraction barrier.

Figure II- 31

 31 Figure II-31 a) Scheme of SEM function mode; b) Emitted particles from electron beam interaction (from Radiological and Environmental Management of Purdue University 17 ); c) FEI Quanta 250 equipment present in LMGP

  scheme of the TEM components is visible in Figure II-32 a).

Figure II- 32

 32 Figure II-32 a) Scheme of components of a TEM Drawing by Graham Colm, courtesy of Wikimedia Commons; b) JEOL JEM 2010 microscope in LMGP

  Figure II-33 a) and b) Description of AFM modes of operation( from University of California Santa Barbara 20 ) c) Veeco D3100 AFM in LMGP

Figure II- 34

 34 Figure II-34 a) X-Ray diffraction on a crystal, where the crystallographic planes and X-Ray wavevectors are represented 22 ; b) Bruker D8 Advance in LMGP

Figure

  Figure II-35 a) Scheme with Rayleigh and Raman scattering representation 25 ; b) Jobin Yvon/Horiba LabRam spectrometer in LMGP

(

  scheme of this technique is visible in Figure II-36 a). In the analysis of Cu2O films, this technique proves to be interesting for the detection of organic species in the material. In case of incomplete degradation of the precursors, residual organic substances can indeed be present in the film, which can lead to meagre physical properties. Additionally, metaloxygen bonds can be detected, confirming the presence of dopants in the films.

Figure

  Figure II-36 -a) Scheme of FTIR function mode 27 ; b) Bruker Vertex 70V spectrometer present in

Figure

  Figure II-38 a) Scheme of light pass in the UV VIS NIR spectrometer; b) Lambda 950 spectrophotometer from Perkin Elmer present in LMGP

  -4 probe provided by Lucas Lab, Figure II-39 b), with a distance between tips of 1.066 mm and connected to a Keithley multimeter, Figure II-39 a).

Figure II- 39

 39 Figure II-39 a) Keithley multimeter and b) 4 probe tips

Figure II- 40

 40 Figure II-40 a) Representation of van der Pauw measurements 30 b) Hall Effect potential measuring set-up 31 .

  The solar simulator is represented in Figure II-41 a). The photovoltaic performance of the solar cells was conducted by using a homedesigned holder, represented in Figure II-41 b) and c). This holder is composed of 11 gold pogo pins. Each pin contacts a different electrode on the sample, which is upside down.The light enters from the top mask that has circular holes with the same area as the metallic contacts.

Figure II- 41

 41 Figure II-41 a) Solar simulator; PV sample holder b) without mask, c) with mask

  appearance of the films observed directly after the deposition gives basic information related to the optical and structural properties. For the lowest temperatures, 300 °C, 325 °C and 350 °C, the films showed good homogeneity with a yellow color, as visible in FigureIII-42 a). The deposition at 375 °C, however, led to the formation of loose powder on top of the substrate. The powder formation is expected at high temperature in this process due to the starting reaction in homogenous phase, rather than in the heterogeneous phase at the substrate surface. Consequently, powder is formed before the arrival of the precursor at the substrate surface. It subsequently settles on the substrate without the required adhesion.The presence of loose powder is not adapted for the type of application that we are aiming for. The SEM micrograph of top view and cross-section corresponding to the film deposited at 350 °C are visible in FigureIII-42 b) and c). The film presents a homogeneous surface, with no individual grains resolved at this magnification level; however, a small roughness is observed, probably due to a nano-crystalline film. The cross-section confirms the homogeneity of the film thickness, with an average thickness in the 100 nm range.

Figure III- 42

 42 Figure III-42 Results obtained from a Cu 2 O thin films deposited by AA-MOCVD using a butanolbased solution and Cu(acac) 2 as precursor. a) visual appearance on Si (left part) and glass (right part) substrate; b) and c) SEM image of top view and of cross-section, respectively, of film deposited at 350 °C;

Figure III- 43

 43 Figure III-43 Cu 2 O thin films deposited by AA-MOCVD: a) XRD spectra on Si substrate with the crystallite size dependence with temperature and b) sheet resistance of films with different temperatures of deposition on glass substrate. Reference spectrum of Cu 2 O: JCPDS n° 04-007-9767 presented in Appendix B

  Figure III-44 Effect of annealing treatment on intrinsic Cu 2 O thin films a) Raman spectra for the film annealed at 300 °C, inset shows a zoom of the 200 -500 cm-1 region; b) Optical direct transmittance of films and c) Sheet resistance dependence with temperature of annealing.

  Figure III-45 shows the plot of dopant amount in the film in relation to total cations ( 𝑋 𝑋+𝐶𝑢 ), as well as the sheet resistance values, after annealing. The thickness of the films presented similar values, all around 100nm.

Figure III- 45

 45 Figure III-45 Atomic concentration of dopant atoms in films (left) and sheet resistance (right) in relation to the dopant atomic number for each tested element. Dashed blue line represents the optimal sheet resistance value of intrinsic Cu 2 O thin film. Dashed red line represents the limit of detection by EDS. Green and blue dopants represent smaller or bigger radius, respectively.

Figure III- 46 ,

 46 the surfaces of the films deposited in the presence of precursors doping with Li, Ca, Co, Ni, Zn, Sr and Mg, as well as intrinsic Cu2O for comparison are represented. Here are the main observations.

Figure III- 46

 46 Figure III-46 SEM microstructure images of Cu 2 O films: intrinsic and deposited with precursors containing Li, Ca, Co, Ni, Zn, Sr and Mg.

Figure III- 47

 47 Figure III-47 XRD spectra of Cu 2 O films: intrinsic and deposited with precursors containing Li, Ca, Co, Ni, Zn, Sr. Reference spectrum of Cu 2 O: JCPDS n° 04-007-97679767 presented in Appendix B

  Figure III-48a) and b), leading to a large transparency as shown in Figure III-48 c), and consequently and increase of bandgap to form 2.2 to 2.8 eV. In this case, we assume that an increase in tin content could lead to the formation of SnO or SnO2. Therefore, we have to discard the possibility of creating a mixed amorphous phase between the two oxides. However this film showed a larger sheet resistance, larger than 100 /sq, which limited the Hall Effect measurements and the determination of the semiconductor type. In order to determine whether we could combine high transparency with lower resistivity, we increased the Sn content furthermore, until both cations had the same concentration in solution. This lead to a large increase of transparency as shown in Figure III-48 c), with a band-gap of 3.0 eV and at the same time a drastic reduction of the sheet resistance to values below /sq as shown in and Figure III-49 a). In terms of cation incorporation, the Sn-doped Cu2O films presented a similar Sn atomic concentration of cations, Sn/Sn+Cu, as the cationic ratio of the prepared solution as shown in Figure III-49b), where measurements were performed by WDS.

Figure III- 48

 48 Figure III-48 Visual appearance Cu 2 O thin films a) intrinsic Cu 2 O and b) Cu 2 O:Sn; c) Optical Transmittance of Cu 2 O thin films doped with Sn with variation of band-gap.

Figure III- 49

 49 Figure III-49 a) Sheet resistance and b) EDS quantification of Sn doped Cu 2 O thin films

Figure III- 50

 50 Figure III-50 a) XRD and b) Raman quantification of Sn doped Cu 2 O thin films Reference spectrum of Cu 2 O: JCPDS n° 04-007-9767 presented in Appendix B

Figure III- 51

 51 Figure III-51 SEM images of Sn doped Cu 2 O thin films: a) 10% Cu 2 O:Sn; b) 33% Cu 2 O:Sn; c) and d) 50% Cu 2 O:Sn top-view and cross-section respectively

  Magnesium has been highlighted in the frame of the elements' screening as a promising dopant for Cu2O from the first electrical resistivity results in Chapter III. The following chapter is focused on an in-depth study of the deposition and characterization of magnesium-doped cuprous oxide (Cu2O:Mg) thin films. The first part of the chapter describes in detail the effect of magnesium on cuprous oxide by varying the concentration of the dopant. The main content of this study was published as a peer-review article: Resende, J., Jiménez, C., Nguyen, N. D. and Deschanvres, J.-L. (2016), "Magnesiumdoped cuprous oxide (Cu2O:Mg) thin films as a transparent p-type semiconductor" Phys. Status Solidi A, 213: 2296-2302. Additionally, it is important to report on the work about strontium-doped cuprous oxide and the comparison to the magnesium case developed in this chapter, which finally led to another publication: Brochen, S., Bergerot, L., Favre, W., Resende, J., Jiménez, C., Deschanvres, J. L., & Consonni, V. (2016). "Effect of Strontium Incorporation on the p-Type Conductivity of Cu2O Thin Films Deposited by Metal-Organic Chemical Vapor Deposition". Journal of Physical Chemistry C, 120(31), 17261-17267. On the second part of the chapter, we focus on the study of the stability of the magnesium-doped cuprous oxide thin films under different oxidizing thermal treatments. In the last third section of this chapter, we study the feasibility and the performance of an oxide-based pn junction integrating Mg-doped Cu2O films. For the purpose of that investigation, the selected n-type semiconductor is ZnO, which is stacked in combination with intrinsic and Mg-doped Cu2O thin films. The pn junction study is divided in two different parts: numerical simulations and device measurements of experimentally deposited films. The ZnO oxide thin films were obtained by chemical methods in collaboration with two other PhD students from LMGP: Thomas Cossuet for the PI-MOCVD and Viet Nguyen in the SALD case. Some results of this study led to a peerreviewed publication: Nguyen, V. H., Resende, J., Jiménez, C., Deschanvres, J.-L., Carroy, P., Muñoz, D., Bellet, D. and Muñoz-Rojas, D. (2017) "Deposition of ZnO based thin films by atmospheric pressure spatial atomic layer deposition for application in solar cells." Journal of Renewable and Sustainable Energy 9, no. 2: 021203.

  Cu2O film (5% in solution) is shown in Figure IV-52. On glass, the coating presents a yellow colour, transparent and homogenous throughout the whole sample (area of 4 cm 2 ), while on silicon the colour depends on the film thickness. In the case of the sample shown in Figure IV-52, the thickness was 77 nm and the surface appears mostly dark blue. In both cases, the films adhere to the substrate, resist to water and Scotch-Tape tests. Similar appearance is visible in all Cu2O:Mg films.

Figure IV- 52

 52 Figure IV-52 Photographs of the Mg doped Cu 2 O thin film on glass (left) and silicon (right) with a thickness of 77nm using a solution prepared with 5% atom of Mg.

Figure IV- 53

 53 Figure IV-53 EDS spectra obtained at an acceleration voltage of 5 keV of Cu 2 O:Mg thin films on glass. The Mg ratio indicated was quantified by EDS analysis (Table IV-6). The different elements X-Ray

Figure IV- 54 .

 54 Figure IV-54. This suggests a more efficient deposition of copper than magnesium under the selected deposition conditions. Higher temperature of deposition might increase the presence of magnesium as reported previously4 , since MgO polycrystalline films are normally deposited at temperatures between 400°C and 500°C. Though, these high temperatures must induce the presence of CuO phase5 or powder deposition on the substrates, as reported in Chapter III.

Figure IV- 54

 54 Figure IV-54 Relative atomic Mg content in the deposited Cu 2 O:Mg films on glass as obtained from EDS data compared to the amount in solution. Black dotted line represents the 1:1 relation.

Figure IV- 55

 55 Figure IV-55 SEM images of Cu 2 O:Mg thin films with different Mg content deposited on glass at 350°C. a) Cu 2 O intrinsic, b) 1%, c) 3%, d) 9% and e) 17% Cu 2 O:Mg with respective cross-section. Same magnification was used for all the top-view pictures. Top right sample corresponds to a cross-section view of intrinsic Cu2O. Grain size enlarges with the increase of Mg content

Figure IV- 56

 56 Figure IV-56 HRTEM image of the 17% Mg doped Cu 2 O thin film: a) cross section view on the silicon interface region, b) corresponding diffraction pattern identified as Cu 2 O structure; c) and d) detail of the film top part removed by scratching

Figure IV- 57

 57 Figure IV-57 AFM images 2µmx2µm of the Cu 2 O:Mg thin filmsa) Cu 2 O intrinsic, b) 1%, c) 3%, d) 9% and e) 17% Cu 2 O:Mg. The Z scale is different for each sample.

Figure IV- 58

 58 Figure IV-58 Spectra of the Cu 2 O:Mg thin films on silicon obtained in a Bragg-Brentano configuration. Reference spectrum of Cu 2 O: JCPDS n° 04-007-9767 presented in Appendix B

Figure IV- 59

 59 Figure IV-59 a) Raman spectra of the Cu2O:Mg thin films on glass; b) FWHM analysis of 149 cm-1 peak with the variation of Mg content

Figure IV- 60

 60 Figure IV-60 a) FTIR spectra of the Cu 2 O:Mg thin films on silicon. b) FWHM and position analysis of 436cm - 1 peak with the variation of Mg content

Figure IV- 61

 61 Figure IV-61 Transmittance spectra from 250 nm to 1500 nm of the Cu2O and Cu2O:Mg (1% and 17%) thin films on glass. The bare glass substrate transmittance is also represented as reference

Figure IV- 62

 62 Figure IV-62 Electrical resistivity dependence on the magnesium content of Cu 2 O:Mg thin films after thermal annealing. The dotted line is represented as a guide for the eye.

Figure IV- 63

 63 Figure IV-63 Mobility (left vertical axis) and charge-carrier density (right vertical axis) as function of the magnesium content in Cu 2 O:Mg thin films

Figure IV- 64

 64 Figure IV-64 Incorporation of Mg in a tetrahedral position of Cu 2 O crystalline structure. Copper atoms pictured as blue circles, oxygen as red ones and magnesium as black ones.

  change from Cu2O to CuO is accompanied by a change in colour and morphology of the films. As the CuO phase presents a darker brown colour, the visual analysis of the films is a clear evidence when the films change their phase. In the Figure IV-65, the 8 samples studied are represented.

Figure IV- 65

 65 Figure IV-65 Optical appearance of the copper oxide thin films. Cu 2 O: as deposited (a), 250ºC (b), 350ºC (c), 450ºC (d); Cu 2 O:Mg: as deposited (e), 250ºC (f), 350ºC (g), 450ºC (h).

  present in lower proportion. When analysing results obtained from undoped films presented in Figure IV-66, only the diffraction peaks corresponding to the Cu2O phase are detected in the as-deposited film and on the film annealed at 250ºC. When the annealing temperature increases to 350ºC, diffraction peaks at 35.5 ° and 38.7° are detected, which can be attributed to the (11-1) and (111), reflection of the CuO phase. For the sample annealed at 450ºC, the majority of the diffraction peaks are associated to the CuO phase, while a small residue corresponding to the Cu2O phase is still visible, mainly diffraction peaks at 36.4°, 42.3° and 61.3° angles corresponding to the (111) (200) and (220) reflections.

Figure IV- 66

 66 Figure IV-66 GI-XRD spectra of undoped Cu 2 O thin films on glass corresponding to the asdeposited and annealed samples. Reference spectrum of Cu 2 O: JCPDS n° 04-007-9767 and CuO: ICDD n° 00-048-1548 presented in Appendix B and C

  intense. It is worth noting that the Cu2O diffraction peak also shows a shift of 0.42º degrees from the original position of 36.42º to 36.84º. This can be an evidence of the change in lattice parameter from 4.270 Å to 4.226 Å, decreasing the size of the unit cell.

Figure IV- 67

 67 Figure IV-67 GI-XRD spectra of the Mg doped Cu 2 O thin films on glass corresponding to the asdeposited and annealed samples. Reference spectrum of Cu 2 O: JCPDS n° 04-007-9767 and CuO: ICDD n° 00-048-1548 presented in Appendix C and D

Figure IV- 68

 68 Figure IV-68 SEM micrographs of Cu 2 O thin films: as deposited (a), 250ºC (b), 350ºC (c), 450ºC (d); Cross-section image visible on image a)

Figure IV- 69

 69 Figure IV-69 SEM micrographs of Mg doped Cu 2 O thin films: as deposited (a), 250ºC (b), 350ºC (c), 450ºC (d); Cross-section image visible on image a)

  STEM to identify the local distribution of Mg. Mappings for each element obtained from the Cu2O:Mg film annealed at 350°C are presented in Figure IV-70. This analysis confirms the presence of Mg through all film without significant segregation in the grain boundaries, not easily distinguished. However, the dopant (Figure IV-70c)) appears to present a gradient with the thickness, being preferably located at the interface with the substrate, in the first 100 nm of the film.

Figure IV- 70

 70 Figure IV-70 a) Bright-field STEM image of Cu 2 O:Mg thin film annealed at 350°C b-e) Corresponding EDS-STEM elemental mapping of the b) Si, c) Mg, d) Cu, and e) O elements, respectively. f) Superimposition of the corresponding EDS-STEM maps of the Si, Mg, Cu and O elements. The 250 nm scale bar is valid for all images.

  Raman spectroscopy was used to clarify the early stage of phase transformation in these sets of samples. This technique is complementary to XRD analysis, but allows for a more local analysis. Raman spectra of as-deposited and annealed Cu2O thin films are represented in Figure IV-71 a) and b), for the undoped and doped samples respectively.

Figure IV- 71

 71 Figure IV-71 Raman spectra of the Cu 2 O thin films on glass: a) intrinsic Cu 2 O; b) Mg doped Cu 2 O.

  Figure IV-72 shows the different spectra, with the bonds related to Cu2O, CuO and MgO phases indicated. As already described in the previous section, Mg-O are detected in the Mg doped films. The main difference in this group of samples is the detection of a bond (478 cm -1 ) related to CuO phase on the films annealed at 350°C which confirm the trend visible previously in GIXRD and Raman.

Figure

  Figure IV-72 FTIR spectra of the Cu 2 O thin films deposited on Si. Set of samples attributed to intrinsic and Mg doped Cu 2 O, both as deposited and annealed at 350°C.

Figure

  Figure IV-73 a) Cu 2p XPS spectra obtained for the Cu 2 O thin films deposited on glass corresponding to intrinsic and Mg doped Cu 2 O, both as deposited and annealed at 250°C. b) Reference spectra of metallic Cu, Cu 2 O and CuO adapted from "www.xpsfitting.com"

Figure IV- 74

 74 Figure IV-74 EXAFS raw data from Athenaa) Cu 2 O and b) Cu 2 O:Mg

  Figure IV-76 a) and b), show this variation for the three annealing temperatures, in the case of undoped and doped films, respectively. To compare the results we use the resistance at any time normalized by the initial value, given for each set of samples at the initial point of the curve. The sheet resistance values given for each curve correspond to the values obtained after annealing.

Figure IV- 76

 76 Figure IV-76 Normalized sheet resistance of the Cu 2 O thin films during the 3 different annealing steps: 250°C, 350°C and 450°C. a) Intrinsic Cu 2 O films; b) Mg doped Cu 2 O

  Cu2O films are shown in Figure IV-77 a), b) and c) for the three annealing temperatures.

Figure IV- 77

 77 Figure IV-77 Dependence of the resistivity of the Cu 2 O thin films with temperature during the 3 different annealing treatments: a) 250°C, b) 350°C and c) 450°C. Grey and red triangles visible in b) to help the visualization of different slopes; d) activation energies for each sample during heating and cooling.

Figure IV- 78

 78 Figure IV-78 shows the resistivity of the two group of films, measured at room temperature.

Figure IV- 78

 78 Figure IV-78 Resistivity dependence on the temperature of the annealing treatments for Cu 2 O and Cu2O:Mg thin films.

Figure IV- 79

 79 Figure IV-79 a) and b) shows the effect of the annealing temperature on mobility and charge carrier density, respectively, of the intrinsic and Mg-doped Cu2O films. It was not possible to perform the Hall Effect characterization on the Cu2O film annealed at 450°C

Figure IV- 79

 79 Figure IV-79 a) Mobility and b) charge carriers density dependence on the temperature of the annealing treatments for Cu 2 O and Cu 2 O:Mg thin films.

  The band diagram and charge concentrations for simulation A are plotted in Figure IV-80. In the Cu2O case, the band diagram shows the formation of a depletion region mostly on the Cu2O side, as the concentration of acceptors is lower, compared to the donors on ZnO. Consistently to what was suggested in the literature 33 , this region is larger than 1 m in intrinsic Cu2O reaching up to 3m in weakly, non-intentionally, doped Cu2O, 10 14 cm - 3 . The concentration of electrons and holes in Figure IV-80b) represents the expected profile of majority carriers in both semiconductors, holes in Cu2O and electrons in ZnO, with an abrupt variation especially in the n-type side case.

Figure IV- 80

 80 Figure IV-80 a) Band diagram and b) defects concentration of simulation A Cu 2 O thin film with low density of acceptors. Depletion region is represented at yellow in figure a)

Figure IV- 81

 81 Figure IV-81 a) Band diagram and b) defects concentration of simulation B -Cu 2 O:Mg thin film with high density of acceptors on ZnO

Figure IV- 82

 82 Figure IV-82 a) Band diagram and b) Current Density throughout the device of simulation B Cu 2 O:Mg thin film with high density of acceptors on ZnO thin film under an applied bias of +2.0V

Figure IV- 83

 83 Figure IV-83 JV curve of both simulated diodes Cu 2 O/ZnO and Cu 2 O:Mg/ZnO between 0 and 2V

Figure IV- 84

 84 Figure IV-84 pn structures created for Cu 2 O based pn junction study: a) 1 stglass/ ITO / ZnO (MOCVD) / Cu 2 O / metal/; b)2 ndglass/ ITO / Cu 2 O / ZnO (SALD) metal. Direction of illuminated light represented by brown arrows

4. 3 . 3 . 1 1

 331 st Structure: ZnO (PI-MOCVD) and Cu 2 O:Mg (AA-MOCVD) A photograph of the first structure is shown in Figure IV-85, where the ITO uncoated layer is visible in the edge, used for bottom contact. The ZnO film covers the rest of the transparent part of the device, while the central square of Cu2O was masked during deposition. Finally, the structure is completed by a circular Au contact of with 3mm diameter, thermally evaporated.

Figure

  Figure IV-85 Photograph of the first structure: glass / ITO / ZnO (MOCVD) / Cu 2 O / Au.

Figure

  Figure IV-86 SEM micrograph of the top and cross section of a heterostructure of ZnO and Cu 2 O (a) and b)) and Cu 2 O:Mg (c) and d))

Figure

  Figure IV-87 SEM micrograph of the cross section and mappings obtained by EDS of the Cu 2 O:Mg/ZnO junction

Figure

  Figure IV-88 a) XRD spectra of Cu 2 O:Mg and Cu 2 O:Mg/ZnO structure b) Transmittance spectra of Corning glass, ITO covered glass, ZnO/ITO film and the whole structure: Cu 2 O:Mg/ZnO/ITO

Figure

  Figure IV-89 JV characteristic of pn junction ZnO/Cu 2 O (intrinsic and Mg doped). Current density presented in absolute value

Figure IV- 90

 90 Figure IV-90 JV characteristic of pn junction between ZnO and Cu 2 O (intrinsic and Mg doped) with and without the anneal step at 260°C. ). Current density presented in absolute value

Figure IV- 91

 91 Figure IV-91 Photograph of the second structure: glass / ITO / Cu 2 O / ZnO (SALD) / Au. Size of 2.5cm by 1.5cm

Figure

  Figure IV-92SEM micrograph of the ZnO/Cu 2 O junction; a) intrinsic Cu 2 O and b)Cu 2 O:Mg

Figure

  Figure IV-93 JV characteristic of pn junction ZnO/Cu 2 O (intrinsic and Mg doped) with ZnO obtained by SALD (dot curves) and MOCVD (lines). ). Current density presented in absolute value

Figure

  Figure IV-94 are plotted the JV curves under illumination. There were six different pn junctions tested: Cu2Oas-dep/ZnOMOCVD, Cu2O250°C/ZnOMOCVD, Cu2O:Mgas-dep/ZnOMOCVD, Cu2O:Mg250°C/ZnOMOCVD, CuO2/ZnOSALD and Cu2O:Mg/ZnOSALD. All of them showed measurable photovoltaic responses except the Cu2O/ZnOMOCVD pn junctions, neither asdeposited nor annealed. The relevant photovoltaic parameters of the four working solar cells as open-circuit voltage (VOC) and short-circuit current density (JSC), fill-factor (FF) and efficiency (η) are represented in Table IV-14. In the MOCVD-grown case, the structure with the as-deposited film showed a VOC of 221mV that was drastically decreased after annealing, down to 21.4mV. Even if the annealing temperature used was below the formation of CuO, the thermal treatment seem to short-circuit part of the cell, reducing the overall VOC. Additionally, the current density values are quite low, below 1 A/cm 2 , resulting in efficiency values far below 0.01%.

Figure

  Figure IV-94 JV curves under AM 1.5G illumination of different pn junctions a) Cu 2 O:Mg/ZnO (MOCVD) as deposited and annealed at 250°C; b) Cu 2 O/ZnO (SALD) with Cu 2 O and Cu 2 O:Mg as-deposited.

  In a recent work(2017), developed in the SPIN lab at the University de Liege by Y. Malier under a Master Thesis internship, undoped and Mg-doped Cu2O thin films were prepared by radio-frequency magnetron sputtering. Pn junctions were then arranged in combination with ZnO:Al. Similar results were reported in terms of diode behaviour, nevertheless, the current-voltage characteristics of pn junctions created together with

4

 4 Photodetectors 5.2.1 Introduction 5.2.2 Experimental 5.2.3 Structural characterization of core-shell structures 5.2.4 Electrical characterization of core-shell heterostructures 5.2.5 Conclusions on the core-shell heterostructures 5.3 Conclusions 5.4 References

  Figure IV-52. The visual inspection of the film shows clearly the difference of phase, compared to Cu2O (Figure IV.1), since the gap energy is higher, so the film is more transparent in the visible range. Nevertheless, due to the absorption in the farend of the visible range, a light grey color is observed when deposited on glass. As seen before, a good homogeneity throughout the whole sample (area of 4 cm 2 in each substrate) is obtained by this AA-MOCVD method. The film deposited on Silicon shows a thickness dependent color, in this case is close to100 nm, creating a light blue color film.

Figure V- 95

 95 Figure V-95 Photographs of the CuCrO 2 thin film on glass (left) and silicon (right) with a thickness

Figure V- 97

 97 Figure V-97 SEM images of CuCrO 2 thin films deposited on glass at 400°C. Same magnification was used for all the top-view pictures (50000x) with the same scale for all the images. Cu/Cu+Cr ratio: a) 24%, b) 46% with the respective cross-section and c) 56%

Figure

  Figure V-98 XRD spectra in the Bragg-Berentano configuration of the CuCrO 2 thin films on glass. Reference spectrum of CuCrO 2 : ICDD n° 04-010-3330

1 13 .

 13 The comparison of our spectrum with IR spectra of the CuGaO2 and CuAlO2 delafossite compounds from Pellicer-Porres et al.14,15 indicates that the CuCrO2 additional modes in the 520-650 cm -1 range very close to those of the three Eu and A2u CuGaO2 and CuAlO2 modes, as it is already the case for the first Eu mode at 107 cm -1 (122 and 144/149 cm -1 in CuGaO2 and CuAlO2, respectively). Owing to the very weak Raman signal of the additional modes, we cannot accurately symmetry assign them, but they are very likely IR modes. The presence of these modes, normally Raman-inactive in the CuCrO2 Raman spectrum, indicates relaxed selection rules generated by a local symmetry breaking probably due to the presence of defects in the structure, such as Cu vacancies or oxygen interstitials. Moreover, the Raman spectra of both the CuCrO2 thin films also emphasize the presence of carbon-based residual species, related to the by-products of the organic ligand of the AA-CVD precursors through the observation of the D-and G-bands characteristic of C-C bonds pointed at 1356 and 1600 cm -1 .

Figure V- 99

 99 Figure V-99 Raman spectra of the CuCrO 2 thin films on glass

Figure V- 100

 100 Figure V-100 b). The two samples with higher Cu shows a value close to 3.2eV, while the 24% film has a value higher, which is attributed to the amorphous phase discussed previously.

Figure V- 100

 100 Figure V-100 a) Transmittance spectra from 250 nm to 2000 nm of the CuCrO 2 thin films on glass. The bare glass substrate transmittance is also represented as reference. b) Direct band-gap obtained from the transmittance spectra

Figure V- 101

 101 Figure V-101 Resistivity values obtained for the CuCrO 2 thin films as a function of copper content. Resistivity values for Cr 2 O 3 and Cu 2 O films are also represented

  morphologies of the ZnO/CuCrO2 core-shell NW heterostructures at different elaboration step are shown in Figure V-102 by top-view and cross-sectional FESEM images. At left, SEM images show the ZnO NWs with a mean length around 800 nm and a mean diameter in the range of 35 to 65 nm. They grow vertically on the 40 nm-thick polycrystalline ZnO seed layer to form a dense array of about 53 NW/µm 2 . They typically nucleate on the c-axis oriented grains of the seed layer, and subsequently elongate along the polar c-axis 23 . Following the deposition of the CuCrO2 shell, the smooth surface and hexagonal shape of ZnO NW is transformed into a rougher surface and rounded morphology over the whole NW array, as seen in Figure V-102 c). Consequently, the ZnO/CuCrO2 core-shell NW heterostructures have a broader apparent diameter in the range of 50 to 200 nm and are separated each other by a decreased mean distance. The CuCrO2 shell follows the elongated shape of the ZnO NWs with high uniformity and conformity, from their bottom to their top, as shown in Figure V-102 d). Top-view and cross-sectional FESEM images of the final whole self-powered UV photodetectors made of ZnO/CuCrO2 core-shell NW heterostructures, including the thermo-evaporated Ag contact, are presented in Figure V-102 e) and f). The Ag contact forms a rough thin film and appears to cap the ZnO/CuCrO2 core-shell NW heterostructures, by mostly covering their top and slightly their vertical sidewalls to an approximate depth of about 300 nm.

Figure V- 102

 102 Figure V-102 FESEM images at different steps during the fabrication of the core-shell heterostructure: a) Top-view and b) cross-sectional of ZnO NW arrays grown on ITO/glass substrate, c) Top-view and d) cross-sectional of CuCrO 2 / ZnO NW heterostructures, and e) Top-view and f) crosssectional of CuCrO 2 / ZnO NW heterostructures covered with the Ag contact. The 500 nm scale bar is valid for images (a-d) and 1µm for f) only.

Figure V- 103

 103 Figure V-103 a) GIXRD patterns of the CuCrO 2 thin film deposited on glass substrate, of ZnO NW arrays grown on ITO/glass substrate, and of ZnO/CuCrO 2 core-shell NW heterostructures grown on ITO/glass substrate.

  Raman spectra of both the CuCrO2 thin film and the ZnO/CuCrO2 core-shell NW heterostructures also emphasize the presence of carbon-based residual species, related to the by-products of the organic ligand of the deposition technique, through the observation of the D-and G-bands characteristic of C-C bonds pointed at 1356 and 1600 cm -1 .

Figure V- 104

 104 Figure V-104 Raman scattering spectra of the CuCrO 2 thin film deposited on glass substrate, of ZnO NW arrays grown on ITO/glass substrate, and of ZnO/CuCrO 2 core-shell NW heterostructures grown on ITO/glass substrate. The CuCrO 2 defect-induced modes are indicated by a green star.

Figure V- 105

 105 Figure V-105 a) TEM image of a ZnO/CuCrO 2 core-shell NW heterostructure. b) Corresponding SAED pattern identifying the CuCrO 2 rhombohedral delafossite phase along with a theoretical diffraction pattern of the present phase as computed with the software JEMS. c) Corresponding HRTEM image of the ZnO / CuCrO 2 interface recorded close to the top face of the ZnO NW.

Figure V- 106

 106 Figure V-106 a) Bright-field STEM image of a ZnO/CuCrO 2 core-shell NW heterostructure. b-e) Corresponding EDS-STEM elemental maps of the b) Zn, c) Cu, d) Cr, and e) O elements, respectively. f) Superimposition of the corresponding EDS-STEM maps of the Zn, Cr, and Cu elements. The 100 nm scale bar is valid for all images.

Figure V- 107

 107 Figure V-107 a) ASTAR map representing the crystalline phases of a ZnO/CuCrO 2 core-shell NW heterostructure, the ZnO and CuCrO 2 crystal phases being represented in red and green, respectively. The insets are the electron diffraction patterns collected at the points designated by A and B on the corresponding map, in which the open circles represent the theoretical electron diffraction patterns of wurtzite ZnO and rhombohedral CuCrO 2 . b) Corresponding ASTAR map representing the crystalline orientations of the ZnO/CuCrO 2 core-shell NW heterostructure, along the X-and Y-axes as indicated in the inset. The color scales for wurtzite ZnO and rhombohedral CuCrO 2 are given in the insets.

A

  Figure V-108 a) together with the corresponding energy band diagram in Figure V-108 b).

Figure V- 108

 108 Figure V-108 a) Schematic of the self-powered UV photodetector based on ZnO/CuCrO 2 core-shell NW heterostructures. b) Corresponding energy band diagram.

Figure V- 109

 109 Figure V-109 J-V measurements under darkness conditions of the self-powered UV photodetector made of ZnO/CuCrO 2 core-shell NW heterostructures. The value of the rectification ratio at ± 1 V is 5500.

Figure V- 110

 110 Figure V-110 J-V measurements under darkness conditions of ZnO/CuCrO 2 core-shell NW and planar heterostructure. SEM cross section of both devices are also represented.

  of the ZnO/CuCrO2 core-shell NW heterostructures as deduced from total transmittance and direct reflectance measurements together with the responsivity of the corresponding self-powered UV photodetector are shown in Figure V-111.

Figure V- 111

 111 Figure V-111 Responsivity and absorption measurements of the self-powered UV photodetector made of ZnO/CuCrO 2 core-shell NW heterostructures, as well as absorption measurements of the ZnO NW array deposited by CBD on ITO/glass and of the CuCrO 2 thin film deposited on glass.

6

  Chapter VI -Conclusions and future perspectives In the work here presented, we successfully prepared two different copper-based oxides thin films, aiming on the electrical and optical enhancements, and we integrated them into pn junctions for energy and light sensing devices.After a preliminary screening of different cations as possible p-type dopants for Cu2O, the incorporation of magnesium in cuprous oxide thin films was successfully achieved by aerosol-assisted metal-organic chemical vapour deposition under atmospheric pressure condition at 350 ºC. The magnesium content could be increased up to 17% concentration in cubic Cu2O phase without the appearance of any magnesium-related phase. The incorporation had a strong effect on the electrical properties, reducing the resistivity to 6.6 .cm, by increasing the charge carrier density up to 8.1x10 17 cm -3 . This result was attributed to the increase of simple copper vacancies, induced by the presence of magnesium in a tetrahedral position, leading to an increase of the hole concentration, as calculated by Nolan et al.Post-deposition annealing treatments in oxidizing conditions are commonly performed on Cu2O thin films to enhance electrical and optical properties; nevertheless, parasitic CuO phase usually starts to be formed in the films at 300 ºC. The dopant presence in the films had a substantial impact on the stability of the Cu 2 O phase under these oxidizing conditions. The Mg-doped Cu2O thin films exhibit a reduced amount of CuO phase than the intrinsic Cu2O films when compared on the same thermal treatments conditions. This fact is due to a different CuO formation mechanism between intrinsic and doped thin films. As suggested by Isseroff and Carter, doping with Mg can prevent the formation of split vacancies for a single cation vacancy, since the divalent cation would assume a position similar to a split copper vacancy in the crystallographic structure.Consequently, the reduction of split cooper vacancies leads to lower concentration of CuO nucleation centres, which inhibits the formation of the parasitic phase in the bulk of the film. Moreover, the reduce presence of CuO phase in the Mg-doped Cu2O has advantageous consequences for thin films properties, showing a lower resistivity and higher transparency than the intrinsic Cu2O.Concerning the doped Cu2O material, a comprehensive analysis of other dopants, especially Sn, Ni and Co, could be required to understand the impact of these divalent cations in a broader way. The use of other deposition techniques, using chemical or physical routes, would possibly provide a more efficient way to incorporate these elements, which posed substantial challenges for the constraints of the AA-MOCVD technique. A comprehensive study on the temperature-dependent Hall Effect on the Cu2O:Mg thin films could also clarify the energy level of different acceptors. Moreover, other strategies can be explored to improve the conduction properties of Cu2O thin films, combining it with other conductive materials.

Cu2O:

  Mg led to the application of these thin films to pn junctions formed with ZnO, an n-type transparent conducting oxide. The structures showed diode-like behaviour with rectification ratios in the 100 range and photovoltaic response with an open-circuit voltage values up to 300mV. Nevertheless, the solar cells showed low current density values, below 1mA.cm -2 due to intrinsic characteristics of the CVD-prepared films. Low carrier mobility and presence of carbon species hinder the path to junction improvement. These observations stress that the need of using supplementary techniques, as thermal oxidation of metallic copper or electrodeposition in the perspective of Cu2O:Mg-based solar cell design.

Table I -

 I 1 Intrinsic point defects on Cu 2 O with respective reaction equation, charge and formation

	enthalpy energy
	Defect point		Reaction equation	Defect	Formation Enthalpy
							Charge	(eV/defect)
	𝑽 𝑪𝒖 ′	1 2	𝑂 2 → 𝑂 𝑂 𝑋 + 2𝑉 𝐶𝑢 ′ + 2ℎ .	-1	0.41-1.66 13,14
	𝑶 𝒊 ′′		1 2	𝑂 2 → 𝑂 𝑖 ′′ + 2ℎ .		-2	1.94/1.87 14
	𝑽 𝑶 ..	𝑛𝑢𝑙𝑙 → 𝑉 𝑂 .. + 2𝑒 ′ +	1 2	𝑂 2	+2	9.9 15
	𝟐𝑽 𝑪𝒖 ′ + 𝑽 𝑶 ..		𝑛𝑢𝑙𝑙 → 2𝑉 𝐶𝑢 ′ + 𝑉 𝑂 ..

Table I -

 I 2 -Cu 2 O literature review with deposition technique temperature, transparency, band gap,

	conductivity and mobility
	Technique	Temperature (ºC)	Band-gap (eV)	Resistivity (Ω.cm)	Mobility (cm 2 .V -1 .s -1 )	Year Ref
	Electrodeposited	RT	2.0	-	-	2011 23
	ALD	225	2.52	125	5	2012 24
	Spray pyrolysis	350	2.2	104	0.2	2014 25
	Electrodeposited	RT	2.32	-	-	2014 26
	DC magnetron sputtering in	RT Annealing @ 280 in air	2.51	7.3	2.7	2014 20
	DC magnetron sputtering	RT -Annealing @ 475 in vacuum	2.4	149	51	2015 27
	RF magnetron sputtering	RT	-	65	-	2016 28
	RF sputtering	RT -Annealing @700 in vacuum	2.43	104	28	2016 22

  Metal-organic Precursors ................................................................... 52 2.1.2 Description of an AA-MOCVD experiment ...................................... 54 2.1.3 Modification of the reactor configuration .......................................... 61

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2.1.4 Annealing treatment ........................................................................... 61 2.1.5 Metals Thermal evaporation ............................................................... 62 2.2 Techniques of characterization 64 2.2.1 Morphological, chemical and structural characterization .................. 64 2.2.2 Spectroscopic and spectrometry techniques ....................................... 70 2.2.3 Electrical and functional properties of materials and devices ............ 76 2.3 Simulation 80 2.3.1 Basic equations ................................................................................... 80 2.3.2 Model's algorithm .............................................................................. 82 2.4 References 84

Table II -

 II 3 Deposition parameters and impact of the thin films deposition

	Parameter	Impact on	Result in maximal conditions	Variation or range
	Solvent	Dissolution of precursor	Precursor residue in case of low dissolution	Butanol or Ethanol
		Evaporation during deposition	Residue of carbon species detected in film
	Solution	Dissolution of precursor	Faster film growth with smaller grains.	From 0.01 mol.l -1 to 0.03
	Concentration	Amount of precursor in the vapour	Precursor residues in glass pieces	mol.l -1
	Frequency of	Size of the droplets	Smaller size of droplets	Around 800 kHz for
	vibration			resonance
	Piezoelectric	Amount of aerosol generated	Denser vapour, faster deposition, smaller grains, lower	35% to 50% of the
	Power		O2 content in the film	generator power
	Gas Flux
	Time	Thickness and saturation	Up to 600nm films	1h to 6h

Velocity and concentration of vapour Possible turbulent flow, faster deposition, smaller grains 5 l.min -1 to 10l.min -1 O 2 ratio Amount of oxygen in the film High O2 content in the film or presence CuO phase 20% to 33% Temperature Evaporation of vapour, decomposition of precursor and grain growth Higher film thickness and grain growth, reduction of organic species, possible formation of CuO phase 300 o C to 400 o C Extraction pressure Velocity and concentration of vapour Laminar flow, higher deposition speed Pressure of 4~5 mmH2O

  𝑅 𝑛 and 𝑅 𝑝 are the net recombination or generation rates for electrons and holes. The last 2 equations required for the simulations come from the definition of current density, in terms of carriers concentration, mobility and difusivity, within a semi-classical drift/diffusion model :𝑱 𝒏 = 𝑞 𝑛 𝜇 𝑛 𝑬 + 𝑞 𝐷 𝑛 ∇𝑛 𝑱 𝒑 = 𝑞 𝑝 𝜇 𝑝 𝑬 + 𝑞 𝐷 𝑝 ∇𝑝 (II.29) where 𝜇 𝑛 and 𝜇 𝑝 are respectively the electron and hole mobilities. The diffusion constants of both charges is defined as 𝐷 𝑛 and 𝐷 𝑝 . The physical consideration for the model is completed by using Boltzmann's approximations for carrier concentrations in the nondegenerate case and Fermi-Dirac occupied level concentration. The equations for the electrons, holes and electrons in trap states are expressed as:

	∇ • 𝐉 𝐩 + 𝑞	𝜕𝑝 𝜕𝑡	= -𝑞 𝑅 𝑝	(II. 26)
	-𝑞	𝜕𝑛 𝑡 * 𝜕𝑡	= -𝑞 (𝑅 𝑛 -𝑅 𝑝 )	(II. 27)
					(II. 28)
		𝑛 = 𝑁 𝑐 𝑒 -𝐸𝑐-𝐹𝑛 𝑘𝑇	(II. 30)
		𝑝 = 𝑁 𝑣 𝑒 -𝐹𝑝-𝐸𝑣 𝑘𝑇	(II. 31)
		𝑛 𝑡 * =	𝑁 𝑡 1 + 𝑔 𝑡 𝑒 -𝐸𝑡-𝐹 𝑡 𝑘𝑇	(II. 32)
					25)

where

  3.1 Intrinsic Cu2O thin films .......................................................................... 87 3.2 Cation doped Cu2O thin films: screening different elements ................... 92 3.3 Tin doped Cu2O thin films ....................................................................... 96 3.4 References ................................................................................................ 99

  The lower temperature value was 300°C, motivated by the temperature at which the Cu(acac)2 precursor decomposes, namely 284 °C. All other parameters, shown in TableIII-4, were kept constants as based on the previous work on Cu2O and the AA-MOCVD expertise achieved in LMGP. The general impact of each parameters was discussed in Table1of Chapter II. In the films here presented, the total gas flow, extraction pressure and time were fixed at, 7 l.min -1 , 4.5 mmH2O and 1 hour to obtain 100 nm thick films.

Table III -

 III 4 Parameters fixed for the deposition of Cu 2 O

	Parameter	Value
	Solvent	Butanol
	Solution Concentration	0.02 mol.l -1
	Frequency of vibration	800 kHz
	Gas Flux	7 l.min -1
	O 2 ratio	25%
	Extraction pressure	4.5 mmH 2 O
	Time	1h

Table III -

 III 5 List of elements tested with respective radius, coordination number precursor used and its molar mass from Sigma Aldrich supplier6 7 . The elements with smaller radius than Cu + 4-fold coordinated are represented in green while the ones with larger radius are represented in blue.

	Cation	Atomic	Radius (pm)	Precursor and molecular mass
		Number
	Li +	3	73(IV)	Li(acac) M= 106.05 g.mol -1
	Mg 2+	12	71(IV)	Mg(acac) 2 M= 258.55g.mol -1
	Ca 2+	20	100(IV)	Ca(acac) 2 M= 257.15g.mol -1
	Co 2+	27	72(IV)	Co(acac) 2 M= 356.26g.mol -1
	Ni 2+	28	69(IV)	Ni(acac) 2 M= 256.91 g.mol -1
	Cu +	29	60 (II) 74(IV)	Cu(acac) 2 M=261.76 g.mol -1
	Zn 2+	30	74(IV)	Zn(acac) 2 M= 263.61g.mol -1
	Sr 2+	38	132 (IV)	Sr(acac) 2 M= 285.84 g.mol -1
	Cd 2+	48	95	Cd(acac) 2 M= 310.63 g.mol -1
	Sn 2+ /Sn 4+	50	132(VI)/69(IV)	Dibutyltin diacetate M= 351.03 g.mol -1
	Ba 2+	56	149(VI)	Ba(TMHD) 3 M= 503.86 g.mol -1
	La 3+	57	177(VI)	La(TMHD) 3 M= 688.71g.mol -1
	Er 3+	68	103(VI)	Er(TMHD) 3 M= 717.06 g.mol -1

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(12) Isono, T.; Fukuda, T.; Nakagawa, K.; Usui, R.; Satoh, R.; Morinaga, E.; Mihara, Y. Highly Conductive SnO2 Thin Films for Flat-Panel Displays. J. Soc. Inf. Disp. 2007, 15 (2), 161-166.

Table IV -

 IV

6

. The magnesium content in the solid film increases linearly with the content in the solution, but the value is about 2.3 times lower than that in the solution, as shown in

Table IV -

 IV 6 Mg/(Mg+Cu) atomic ratio in solution and detected by EDS together with the thickness of each film

			Thickness (nm)
	0	-	86 ± 3
	5	1 ± 1	77 ± 2
	9	3 ± 1	99 ± 9
	17	9 ± 1	82 ± 4
	40	17 ± 1	77 ± 5

Mg/(Mg+Cu) solution (%) Mg/(Mg+Cu) EDS (%)

Table IV -

 IV 7 Crystallite size, roughness and thickness of the Cu 2 O:Mg thin films

	Mg/(Mg+Cu) (%)	RRMS (nm)	Crystallite size (nm)	Thickness (nm)
	0	2.6	30	86 ± 3
	1	8.4	32	77 ± 2
	3	11	30	99 ± 9
	9	7.2	29	82 ± 4
	17	16	41	77 ± 5

Table IV -

 IV 9 General properties of deposited Cu 2 O and Cu 2 O:Mg films for the stability study

	Sample	Mg/(Mg+Cu) in solution	Mg/Mg+Cu in film	Thickness	Sheet Resistance
		(%)	(%)	(nm)	(M/sq)
	Cu2O	0	0	244	6.3
	Cu2O:Mg	33	18%	164	4.7

Table IV -

 IV 10 List of constant unvaried during the simulations

	Material	Cu 2 O	ZnO
	Band Gap (eV)	2,17	3,4
	Defects interface (cm -3 )	1x10 10	1x10 10

  -11 and Table IV-12.

Table IV -

 IV 11 Simulation A -parameters simulating the intrinsic case: Cu 2 O and ZnO films

		Cu 2 O	ZnO
	Thickness (m)	2	2
	Acceptors density (cm -3 )	5x10 15	1x10 10
	Donors density (cm -3 )	1x10 10	1x10 19

Table IV -

 IV 12 Simulation B -parameters simulating the doped case: Cu 2 O:Mg and ZnO films

		Cu 2 O:Mg ZnO
	Thickness m)	0.2	0.2
	Acceptors density (cm -3 )	8x10 17	1x10 10
	Donors density (cm -3 )	1x10 10	1x10 19

Table IV -

 IV 13 Parameters obtain from fittings of pn junctions using equation 1

	Structure	Conditions	J (1V)	RR	n	Rs	RSh	J (0V)
			mA.cm -2	±1V		Ω.cm 2	kΩ.cm 2	A.cm -2
		As deposited	80	2.43	16.	-	-	-
	1 st
	ZnO MOCVD /Cu2O	Annealed	190	2.86	15	-	-	-
		250°C
		As deposited	33	90.0	2.6	8.0	8.0	0.4
	1 st
	ZnO MOCVD /Cu2O:Mg	Annealed	64	88.4	2.7	4.0	7.4	1.2
		250°C
	2 st	As deposited	130	170	2.0	3.5	1.9	2.7
	Cu2O/ZnO SALD
	2 st	As deposited	110	109	2.0	4.3	1.4	2.7
	Cu2O:Mg/ZnO SALD

Jsc mA.Cm -2 FF % η % 1 st MOCVD Cu 2 O:Mg

	Table IV-14 Parameters of the different solar cells created. Open-circuit voltage (V OC ), short-circuit
	current density (J SC ), fill factor (FF) and efficiency (η)
	Sample		V OC
			V
		As-deposited	0,221	-2,5 x10 -4	25,3	<0,01
		Annealed 250°C	0,021	-2,3 x10 -5	23	<0,0001
	2 nd -	Cu2O	0,194	-0,79	25,8	0,04
	SALD	Cu2O:Mg	0,380	-0,16	24	0,02

Table V -

 V 15 Parameters fixed for the deposition of CuCrO 2

	Parameter	Variation
	Solvent	Ethanol
	Solution Concentration	0.02 mol.l -1
	Frequency of vibration	800 kHz
	Temperature	400 °C
	Gas Flux	7.5 l.min -1
	O2 ratio	20%
	Extraction pressure	4,5 mmH2O
	Time	90 minutes

2 Core-Shell Nanowire Heterostructure for Self-Powered UV Photodetectors

	5.2.1 Introduction
	CuCrO2 presents a high interest as p-type semiconductor because of its high
	conductivity, and its wide direct band gap energy. A state of art of this material has been
	resumed in section I.1.4. Besides, it presents a suitable type-II band alignment with ZnO.

  solution of zinc nitrate hexahydrate [Zn(NO3)2•6H2O] and hexamethylenetetramine [C6H12N4] dissolved in deionized water. The substrates were placed face down into a sealed beaker filled with the equimolar solution and heated at 90 °C for 3 hours in a regular oven. After the growth, the substrates were rinsed with deionized water, and dried with nitrogen25 .The deposition of the CuCrO2 thin films in the form of a semiconducting shell covering the ZnO NW arrays was achieved by AA-MOCVD. The solution-based process was identical to the one explain in Section 1.2 in this chapter. The main parameters are resumed in Table V-15. Corning glass and ZnO/ITO covered glass were also used as substrates in order to obtain details on the film properties and compare the thin film with the nanowire configuration. The ZnO film used were deposited by Pulsed Liquid Injection Metal Organic Chemical vapor deposition (PI-MOCVD) at 400 °C into an Annealsys MC-200 reactor similar to the case presented in the Chapter IV.3. Both copper and chromium precursors were equimolar in concentration. The deposition was performed for 1 h at 400

	This ZnO layer is required to control the nucleation, growth and structural morphology of
	ZnO NWs 23 . A 0.375 mol.l -1 equimolar solution of zinc acetate (Zn(CH3COO)2•2H2O) and
	monoethanolamine (MEA) dissolved in absolute ethanol was used. More details can be
	found in Cossuet et al. 24 . A two-step heat treatment was subsequently performed at 300 °C

for 10 minutes and at 500 °C for 1 h to remove residual organic compounds and to crystallize the ZnO seed layer, respectively. The Chemical Bath Deposition (CBD) of ZnO NWs was achieved on the ZnO seed layer/ITO/glass substrate by using a 0.03 mol.l -1 equimolar

* * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * * *

Acknowledgments

Furthermore, I am grateful for the economic support provided by the International Doctoral School in Functional Materials under the project ERASMUS MUNDUS II 2009-2013. Last but not the least, I would like to thank my family: my parents, my sister and the little Tomás, for supporting me spiritually throughout This doctoral thesis is under the framework /). It is cosupervised by Dr. Jean-Luc Deschanvres from LMGP (CNRS, France) and Professor Dr. Duy Ngoc Nguyen from ); as well co-supervised by Dr. Carmen Jimenez (CNRS, France). The main work in this thesis has been conducted in two laboratories: LMGP (materials, and as well in the research group for Solid State Physics,

https://www.idsfunmat.u-bordeaux1.fr

Copper(II) oxide -Tenorite -CuO

complete transformation occurs at higher temperature, above 600 ºC 21 , due to the low rate speed of this phase transformation. Additionally, there is the report of the inverse phase transition from CuO to Cu2O by annealing treatments at 425 ºC under vacuum 22 .

In this study, we annealed both intrinsic and doped Cu2O thin film at different temperatures, from 250 ºC to 500 ºC in air and we analysed the effect of Mg incorporation on the stability of the Cu2O phase. We studied the changes on structural, morphological and electrical properties. Such a deeper investigation provide a comprehensive insight into the phase transformation and the electronic impact of the Mg incorporation.

Experimental

For the study on the stability of the Cu2O films under thermal annealing, two samples, with and without magnesium, were deposited. Once again, a butanol-based solution was used for the deposition by AA-MOCVD. The total concentration of the solution was fixed at 0.03 mol.l -1 . The first solution was composed of pure Cu(acac)2, while in the second one the Mg/Mg+Cu ratio in the solution was fixed at 33%, as it was considered the lowest resistivity sample in the previous Mg variation study (17% Mg in the films). In order to have a thicker film, the time of deposition was increased to 3 hours and the solution consumption rate decreased to 1.5 ml.min -1 . The lower solution consumption enabled a slow deposition rate, beneficial electrically for thicker films. The other deposition parameters were kept constant: 350 ºC for the temperature of the substrates (Corning 1737 and p-type silicon wafer); Argon flow of 6 l.min -1 ; O2 flow of 2.5 l.min -1 . The study of the stability films under thermal annealing was performed in a heating plate under air, coupled with in-situ resistance measurements preformed using a 2-probe system. Three stages of temperature were chosen, 250 ºC, 350 ºC and 450 ºC, all for 30 minutes, while the heating rate was kept constant at 10ºC/min. One additional annealing treatment was performed at 500 ºC for 30 minutes, only analysed by EXAFS. The temperature bounds were defined as follows: the lower limit is set by the preliminary annealing study in Chapter III for undoped Cu2O while the upper limit is determined by the major transformation to CuO phase.

Structural characterization

The as-deposited films were routinely characterized in order to obtain the Mg content in the film, thickness and sheet resistance. The preliminary results are visible in is plotted in Figure V-96, together with the film thickness obtained by WDS. The films are Cr-rich, suggesting a more efficient deposition of chromium than copper under the selected deposition conditions. Additionally, a similar trend was observed by J. Crepelliere et al. 5 using a pulsed-injection MOCVD system; where it was found a deficit in Cu content into films deposited from solution with atomic concentration of Cu/(Cu+Cr) from 40% to 60% . In terms of films thickness, the values vary from 73 to 170 nm with a large increase for a cationic concentration in solution of 60%. For practical reasons in the readability of the text, from this point of this chapter, we will use the Cu/(Cu+Cr) ratio values obtained by WDS to name the different films deposited, shown in Table V-16. predominantly occurs in the ZnO NWs. The maximum responsivity of 1.87 x 10 -4 A/W at 374 nm is achieved, followed by a sharp drop to 1.49x10 -5 A/W at 400 nm. Subsequently, it gradually decreases to reach values lower than 10 -6 A/W above 650 nm. Furthermore, the responsivity drop correlates well with the absorption drop observed for ZnO NWs on ITO/glass substrates without the CuCrO2 shell, as shown in The suitable photovoltaic effect with an open-circuit voltage (VOC) and short-circuit current density (JSC) of 16 mV and 37 µA/cm 2 , respectively, reveal that the ZnO/CuCrO2 heterojunction is able to properly separate and collect the photo-generated charge carriers, Moreover, the device presents a maximum responsivity of 187 µA/W under zero bias at 374 nm and a high UV-to-visible rejection ratio (374-550 nm) of 68 together with a PV response under standard AM1.5G illumination. The combination of the high rectification ratio, high absorption coefficient in the UV region, and of the remarkable responsivity offers promising possibilities for the application of such a heterostructure as a self-powered UV photodetector. Further optimization of this core-shell device can be achieved with the variation of nanowires size and diameter, CuCrO2 shell thickness or even the use of another delafossite material. Moreover, different applications as gas sensor or water splitting device would deserve some attention in the near future.

In conclusion, this research work was dedicated to the synthesis and characterization of copper-based oxides with the constant aim of ultimately incorporating them into opto-electric systems, fulfilling the initial intention, "from material to devices". Appendix 195 7 Appendix

A -EXAFS Data Fittings

In order to identify the EXFS spectra with the crystallographic structure of the material, a theoretical EXAFS spectra is created by Arthemis software, using the crystallographic parameters of the oxide and the variation of EXAFS parameters and then fitted by comparison to optimize the simulation. The list of variable parameters, presented in Table A-17, are related to the number of neighboring atoms (N), distance to the neighboring atom (R), change in half-path length (ΔR), energy shifts depending on the atom type (ΔE), mean squared displacement (σ 2 ) and passive electron reduction factor (S0 2 ).

In our case, we will fixed the N and S0 2 to theoretical values. Only the two first interaction paths will be used for Cu2O, corresponding to Cu-O 1 st and Cu-Cu 1 st first neighbors. N will be 2 for Cu-O and 12 for Cu-Cu. We will use three first interaction paths will be used for CuO, corresponding to Cu-O 1 st , Cu-Cu 1 st and Cu-Cu 2 nd first neighbors.

N will be 4 for all of them. The passive electron reduction factor (S0 2 ) was fixed to 0.813 that allows to fit the Cu-O 1 st peak for both phases and it is in agreement with the literature 3 . For this study we select the as-deposited Cu2O film and the film annealed at 450°C which