David Muñoz-Rojas

Daniel Bellet

Hervé Roussel

Laetitia Rapenne

At this very ending of my PhD life, I am looking back to the three years of thesis and feeling so complicated. There were exciting moments and stressful period, but at this moment I feel so grateful for those people walked with me through it. Thus, I want to give my most sincere thanks to all of them.

First of all, I am so grateful for my family that they always supported me for what I want to do. They are the harbours of my mind which helps me keep peaceful mind. Even though they miss me a lot, they always encourage me to see the world. Then, I was

Ré sumé

Au cours des dernières années, l'électronique à base d'oxydes métalliques a attiré de plus en plus d'attention au sein de la recherche, principalement grâ ce à leur potentiel en termes de ré duction de coûts ainsi que la possibilité de dé velopper une é lectronique transparente. Il existe plusieurs applications potentielles concernant les oxydes mé talliques : le photovoltaï que, les transistors à couche mince et la photo-é lectrochimie. Il existe plusieurs oxydes mé talliques de type n avec d'excellentes proprié té s é lectroniques, telles que l'oxyde de zinc dopé à Al. Mais la mise au point de dispositifs entièrement à base d'oxydes mé talliques est largement entravé e par les mauvaises proprié té s é lectroniques des oxydes de type p jusqu'à pré sent é tudié s. Par consé quent, il est né cessaire de dé velopper des maté riaux semi-conducteurs d'oxyde mé tallique de type p pré sentant de meilleures proprié té s é lectriques.

Dans cette thè se, l'optimisation du dé pôt de films minces de Cu2O a é té effectué e par MOCVD assisté par aé rosol (AA-MOCVD). Par consé quent, des films de Cu2O homogè nes et de trè s forte cristallinité ont é té dé posé s à basse tempé rature (environ 335 °C) sans contamination dé tectable de carbone. De plus, grâ ce à l'incorporation de l'humidité durant les dé pôts, la taille des grains et l'orientation des films Cu2O peuvent ê tre modulé es, ainsi des films de Cu2O avec une texturation (111) et une taille de grains > 300 nm ont é té obtenus.

Pour les films Cu2O optimisé s, la mobilité peut atteindre un maximum de 15 cm 2 / V.s avec une concentration de porteur de l'ordre de 10 15 cm -3 . Enfin, un excellent comportement diode a é té observé en combinant les films de Cu2O optimisé s avec du ZnO, obtenant un rapport on/off supé rieur à 10 En raison de l'incompatibilité entre la fenê tre de stabilité thermique associé e à AgCuO2 et les tempé ratures né cessaires pour dé poser des composé s Ag et Cu par CVD avec les pré curseurs utilisé s, le dé pôt direct d'AgCuO2 n'a pas pu ê tre obtenu. Ainsi, des techniques de revê tement couche mince à base de solution ont é té adopté es pour le dé pôt de film AgCuO2. En particulier, la mé thode SILAR a permis le dépôt de films minces d'AgCuO2.

Grâce à une couche d'amorce sur substrat de verre appropriée, des couches d'AgCuO2 denses et continues ont été revêtues, avec une valeur RMS minimale de 8 nm. Les films d'AgCuO2 dé posé s avaient une phase presque pure. Les proprié té s optiques et de transport des films minces AgCuO2 ont donc é té analysé es pour la premiè re fois. Les mesures de transmittance ont confirmé la faible largeur de bande interdite prédite d'AgCuO2 (1.2 eV), tandis que grâ ce à l'utilisation de la formule de Tauc, nous avons constaté que ce maté riau est plus susceptible d'avoir une bande interdite directe, en accord avec les calculs DFT publié s. Grâ ce aux mesures de l'effet Hall, les films AgCuO2 dé posé s ont é té confirmé s comme é tant de type p.

La plus faible ré sistivité atteinte est de 0.2 Ω.cm. En outre, ces films avaient une densité de porteurs de charge de l'ordre de 10 17 cm -3 et la meilleure mobilité atteinte é tait de 24 cm 2 / V.s. En comparaison avec les composé s de type p de delafossite pré cé demment rapporté s (M, Al, Cr, Ga, etc.), ce maté riau pré sente la plus petite largeur de bande interdite (inté ressant notamment pour l'application photovoltaï que) et une conductivité assez é levé e.

La caracté ristique la plus inté ressante est que le problè me gé né ral de la faible mobilité des transporteurs dans ces composé s delafossite a é té ré solu dans cet AgCuO2, grâ ce à sa structure é lectronique mixte et à la dé localisation des charges. Ainsi, ces ré sultats de caracté risation sans pré cé dent ouvrent la voie à l'utilisation de films AgCuO2 dans des dispositifs fonctionnels.

Abstract

In recent years, metal oxide electronics has attracted more and more attention in research, mainly thanks to their potential lower cost and the possibility they offer to develop transparent electronics. There are several potential applications concerned with metal oxides including photovoltaics, thin film transistors and photo-electrochemistry. There are several n-type metal oxides with excellent electronic properties being well developed, such as Al doped zinc oxide. But the fabrication of devices fully made with metal oxides is largely impeded by the poor electronic properties of the p-type oxides so far studied. Therefore, there is the need for developing p-type metal oxide semiconducting materials with better electrical properties.

In this thesis, the optimization of pure Cu2O thin film deposition was conducted using Aerosol Assisted MOCVD (AA-MOCVD). As a result, homogenous Cu2O films were deposited at low temperature (about 335 °C) without detectable amount of carbon contamination with high crystallinity. In addition, by incorporation of humidity during the deposition, particle size and the orientation of the Cu2O films could be tuned, thus Cu2O films with (111) textured large grain sizes (> 300 nm) were achieved. For optimized Cu2O films, the mobility can reach a maximum of 15 cm 2 /V.s with carrier concentration in the order of 10 15 cm -3 . Lastly, an excellent diode behavior was observed by combining the optimized Cu2O films with ZnO, obtaining an on-off ratio exceeding 10 4 .

Besides the Cu2O optimization, the deposition of AgCuO2 by MOCVD was also tackled. In order to do so, the deposition of silver and silver oxide thin films was previously optimized.

For that, two new silver precursors, namely, Ag(hfac)phenanthroline and Ag(hfac)triglyme were synthesized and fully characterized. High quality Ag coatings could be obtained with both precursors. Silver oxide films were obtained through electrochemical oxidation and oxygen plasma treatment of pre-deposited Ag coatings.

Due to the incompatibility between the thermal AgCuO2 stability window and the temperatures needed to deposit Ag and Cu compounds by CVD with the precursors used, the direct deposition of AgCuO2 could not be obtained. Thus, solution based thin film coating techniques were adopted for AgCuO2 film deposition. In particular, Successive Ionic Layer Adsorption and Reaction (SILAR) allowed the deposition of AgCuO2 thin films.

Using a proper seed layer on glass, dense and continuous AgCuO2 films were coated, with minimum RMS value of 8 nm. The deposited AgCuO2 films had almost pure phase. The optical and transport properties of AgCuO2 thin films were thus carried out for the first time.

Transmittance measurements confirmed the predicted low bandgap of AgCuO2 (1.2 eV), while by using the Tauc formula, we found that this material is more likely to have a direct bandgap, in agreement with published DFT calculations. Thanks to Hall Effect measurements, the deposited AgCuO2 films were confirmed to be p-type. The lowest resistivity achieved was 0.2 Ω.cm. In addition, those films had carrier density in the order of 10 17 cm -3 and the best mobility achieved was 24 cm 2 /V.s. Comparing with the previously reported CuMO2 (M= Al, Cr, Ga etc) delafossite p-type compounds, this material has shown the lowest bandgap (appropriate for photovoltaic application) and rather high conductivity.

The most interesting characteristic is that the general problem of low carrier mobility in those delafossite compounds has been solved in this AgCuO2, thanks to its mixed-valence electronic structure and charges delocalization. Thus, those unprecedented characterization results pave the way for using AgCuO2 films in functional devices. In Chapter IV, the work involving the deposition of copper oxides is presented. Given that the deposition of CuO has been already studied in LMGP, we focussed on Cu2O. During this chapter, the optimization of the AA-MOCVD system is presented. With the optimized configuration, high crystalline and carbon free Cu2O films were deposited using Cu(II) trifluoro acetylacetonate. A study of the effect of i) the oxygen ratio, ii) the presence Ag nanoparticles or ii) the use humid carrier was also performed and is detailed. The effects of those factors during deposition were elaborated from the point view of morphology, crystal orientation and electronic properties. Lastly, through a classical thermodynamic model, we propose a rationalisation of the Cu2O deposition process.

In Chapter V, the work about AgCuO2 thin film deposition will be introduced. Firstly the results of co-depositions using both silver and copper precursor solutions in AA-MOCVD is be introduced. Due to the incompatibility between the precursor decomposition temperature and the AgCuO2 phase, no AgCuO2 was obtained. Then, to deposit AgCuO2 films, alternative low-cost and low temperature deposition techniques were adopted, namely, solution based Successive Ionic Layer Adsorption and Reaction (SILAR). With this technique, 100 nm thick continuous and dense AgCuO2 was coated on Cu2O seed layers.

Preface vii

The optimization of such deposition process is detailed. Then, the UV-Vis characterization of AgCuO2 films are reported for the first time, implying a direct 1.2 eV bandgap.

Furthermore, Hall Effect measurements were also conducted on those samples. Comparing with the previously reported CuMO2 ternary compounds, the AgCuO2 films presented a low resistivity down to 0.2 Ω.cm, higher carrier density in the order of 10 17 cm -3 ~ 10 18 cm -3 and high carrier mobility as high as 24 cm 2 /V.s.

Lastly, in Chapter VI, through the collaboration with Prof. Mari-Ann Einarsrud from NTNU, Norway, the in-situ XRD characterizations of the hydrothermal synthesis of AgCuO2 was conducted in ESRF synchrotron. The results of the characterizations of different reactions with different pH solution, reaction temperatures are presented and discussed. As a result, the hydrothermal reaction mechanism of AgCuO2 formation in this case will be revealed, and an activation energy of 67 kJ/mol was obtained.

Chapter I Introduction

As it is known, Si based semiconductor materials are one of the most important foundations for the modern electronic industry. However, due to a high melting point of Si/SiO2, high purity crystalline Si wafer production is always a heavily energy demanding process, which is one of the major reasons for its high cost. 1 Also, Si based photovoltaic technologies are the most widespread nowadays, but due to the indirect bandgap of crystalline Si, thus much thicker absorber layers are required in solar cells. This imposes a higher degree of purity and crystallinity of the Si used and thus high cost of material fabrication is one of the hurdles for photovoltaic energy to compete with fossil fuels. With the development of new Si based materials such as amorphous-Si (a-Si: H), the production cost is largely reduced with much lower processing temperature. However, such a-Si layers suffer from low carrier mobility (~0.5 -1 cm 2 /V.s), which limits its potential for wider type of applications, such as transistors. 2,3 In addition, crystalline Si has very poor mechanical resistance. And even when using a-Si thin layers, applying Si crystalline/amorphous layers with low-cost coating techniques for electronics on flexible substrates is still rather challenging.

To find an alternative solution, researchers have turned their interests in metal oxide semiconductors. Comparing with tradition Si based semiconductor materials, metal oxide semiconductors have several merits. First of all, many metal oxides are earth abundant and can be processed through lower cost chemical approaches at low temperature. Thus, a much cheaper cost of the semiconductor material fabrication can be achieved. The application of metal oxides layers in solar cells can thus increase the competence of the solar energy in the market. Furthermore, with a lower thermal budget and better mechanical resistance in some cases, some metal oxide materials still showed excellent carrier mobility (e.g. mobility exceeding 100 cm 2 /V.s was reported for Cu2O films deposited at 200 °C by radio frequency Magnetron Sputtering (rf-MS) method 4 ). Thus metal oxide semiconductors have great potential in flexible electronics as well, such as flexible circuits for organic light emitting devices (OLED). [5][6][7] Besides the promising applications with metal oxide semiconductors, since there are several metal oxide semiconductors that have a wide bandgap (such as ZnO, over 3 eV), thus the combination of both transparency or semitransparency with conductivity is possible with metal oxides. Those type of metal oxide materials are called transparent conductive oxides (TCOs), and typical ones are Al doped ZnO (AZO), Indium doped Tin oxide (ITO) and Fluorine doped Tin Oxide (FTO). Combining all characteristics of low cost, transparency and conductivity, metal oxides have even brought new possibilities in several fancy applications such as smart windows and transparent electronic devices. [8][9][10][11] However, the current situation for metal oxide electronics is still facing lots of challenges, one of which is the lack of p type oxides with identical excellent electronic properties as the n type counterparts. 12,13 As a result, the performance of electronic devices requiring p and n metal oxides are dragged down. Only unipolar n type metal oxide based electronic devices have shown a promising future for commercialization. Therefore, developing new p type metal oxides with better electrical properties is urgently needed.

Meanwhile, from the point view of electronic structure configurations, there are several difficulties to achieve a p type metal oxide with similar performance to n type oxides. As stated from the work of Z.Wang et al., 12 in n type metal oxides, the electrons are created by oxygen vacancies and the pathway conduction band minimum (CBM) is composed of widely dispersed metal s orbitals. Thus, there is sufficient hybridization and delocalized CBM in n type oxide, which allows the electrons to travel through the material with high mobility.

However, in p type metal oxides, since the origin of the hole is from metal ion vacancies and the valence band maximum (VBM) is made up with the anisotropic and localized O 2p orbitals, thus low carrier mobility is rather commonly observed in p type metal oxides. 14,15 In addition to the low mobility, due to the high formation energy of metal ion vacancies and low formation energy of oxygen vacancies, the native generation of hole carriers is limited.

Combining both aspects, finding new p type metal oxide material with proper electrical structure for hole conduction is essential to achieve a good performance of metal oxide based electronic devices.

I.1 Electrical properties of existing p type metal oxides

So far several types of p type metal oxides have been reported. With a rough classification, there are SnO, NiO, ZnM2O4 (M= Ir, Rh and Co) spinel type oxide, copper oxide and it's ternary oxide compounds. 12 Table I-1 presents the optical and electronic properties of common p type metal oxides (we have also include n type semiconductors for the sake of comparison). Among the binary compounds, NiO has rather high bandgap more than 3.6 eV and the lowest resistivity, merely about 0.4 Ω.cm. Meanwhile, Cu2O has a smaller bandgap transition in those TCOs, between 2 and 2.7 eV. The resistivity of Cu2O films are also generally high, normally more than 100 Ω.cm. [16][17][18] Even though the optical and electronic properties of Cu2O films can be tuned by doping with other metal ions or anion doping (N mainly) [19][20][21][22][23] , they normally suffer from reduced carrier mobility. The optical bandgap of other binary oxides, such as SnO are between those two values and the reported SnO has relatively low mobility as well. Thus, NiO seems to be promising for transparent electronic with its high bandgap and low resisitivity.

Concerning the ternary metal oxide compounds, recently there are reports about ZnM2O4

(M= Ir, Rh and Co) spinel type compounds that have presented p type conductivity combined with a large bandgap reaching almost 3 eV as well. Thus, those type of metal oxide compounds are also attracting more and more interests for their promising applications in transparent electronics. On the other hand, due to a cheaper cost and low toxicity with copper oxide, there are several works focusing on developing copper based ternary metal oxide compounds. In fact, the largest family of p type ternary metal oxides are copper based, usually with delafossite structure CuMO2 (M= Al, Ga, Cr, etc) or the spinel compound SrCu2O2, as shown in Table I-1. They generally have shown enlarged optical bandgap, up to 3.75 eV with CuAlO2, and reduced resistivity in the range of 0.1~20 Ω.cm. However, due to the presence of high VBM tail states in the electronic structure, most of the reported copper ternary oxide compounds have improper high carrier concentration and low mobility [29][30][31][32][33][34][35] . Those characteristics largely limit their applications, such as in transistors. In addition, by reviewing the listed p type metal oxides in Table I-1, p type metal oxides with small bandgap are rarely reported, which adds up to the difficulty of achieving full metal oxide composed solar cells with good performance. To use the p type metal oxide as absorber layer in solar cell, smaller bandgap around 1.5 eV is required.

I.2 New emerging potential Ag-Cu oxide family for photovoltaic application: AgCuO2 and Ag2Cu2O3.

In addition to our interests on optimization of pure Cu2O phase, the recent family of silver copper mixed oxides family attracts our attention. 36,42,43 The very origin of the birth of Ag-Cu oxides family is from the attempts so seek alternatives for toxic element Hg in superconductive materials. 36 Until now, two types of ternary Ag-Cu mixed oxides have been reported, namely, Ag2Cu2O3 and AgCuO2. With more recent study of their crystal structure and electronic and optical properties, Ag-Cu oxides compound have shown fascinating features for photovoltaic application. In this section, we introduce the crystal structures of Ag2Cu2O3 and AgCuO2, the electronic and optical properties and the synthesis techniques used so far to obtain them.

I.2.1 Crystal and electronic structures of Ag-Cu oxides family

The origin of the silver-copper mixed oxide family was actually from the research for copper based superconductors. Among those, the cuprates containing mercury presented the highest Tc points, however Hg is severely toxic and non-environmentally friendly. 44 Thus, there was a lot of work attempted to replace this element in the structure. The electronic structure of Ag 1+ and Hg 2+ are similar, therefore there was lots of efforts involved and attempted to substitute the Hg element in the superconductor, and this was the initial motivation of synthesizing Ag-Cu oxides. The first Ag-Cu oxide compound reported was Ag2Cu2O3 from the work of P.Gómez-Romero et al. in 1999 by using acid bath chemistry or molten hydroxide fluxes to precipitate the precursor and the crystal structure was proposed. 45 In another work from this group, the Ag2Cu2O3 phase was obtained by reacting Cu(NO3)2 and

AgNO3 in 3 M NaOH solution for 6 hours at room temperature. From that, more detailed studies on this material were investigated including thermal, magnetic and electronic properties. 46 As shown in Figure I-2. Schematic illustration of Ag2Cu2O4 crystal structure with octahedral Ag-O6 (2+4), in which red atom represents oxygen, blue for copper and grey for silver atom (inside octahedral) 47 Besides the Ag2Cu2O3 phase, in a work from D.Muñoz-Rojas et al., a new Ag-Cu oxide phase AgCuO2, was synthesized based on electrochemical oxidation of Ag2Cu2O3. 48 The crystal structure is schematically illustrated in Figure I-2, 47 in which the copper atoms (blue) are surrounded by 4 oxygen atoms (red) each, while silver atoms (grey) are inside a distorted octahedral structure formed by six oxygen atoms. The crystal structure of Ag2Cu2O4 resembles a lot that of AgO (Ag(I)Ag(III)O2) 49 , therefore there is a tendency to view the oxidation state of Ag and Cu in the silver copper oxide as 1+ and 3+.

To explore the charges thus the electronic structure of the AgCuO2, XPS characterization of such compound was carried out and the full spectra is shown in For Ag2Cu2O3 XPS spectra, the positions of the two peaks, locating at 367.9 eV and 373.9

eV, corresponding to Ag 1+ oxidation state in this compound. 47 While in the case of electrochemically deposited Ag2Cu2O4, each originally single peak at Ag 3d3/2 and Ag 3d5/2 is spitted into two peaks. By fitting, they were ascribed to a coexistence of Ag 1+ and Ag 3+ charge states, a mixed valence state classified as type II (fast hopping between Ag 1+ and Ag 3+ ) or type III, totally delocalized. 47,50 Similar mix valence charging states can be also found in Figure I-3 c), implying that Cu ions in this oxide compound could carry both 2+ and unusual 3+ oxidation state. Lastly, the peaks corresponding to O 1s are presented in Figure I-3 d). The fitted lines showed three components corresponding to three binding energies, 528.9, 530.6 and 532.3 eV. Binding energies at 528.9 and 530.6 eV refer to oxide species (O 2-) and binding energy 523.3 eV would correspond to peroxide species (O2 2-). This implies that the extra charges from Ag and Cu are delocalized onto oxygen atoms.

After the discovery of the Ag-Cu oxide tertiary family, other derivatives from this family were also reported. Such as the discovery of Ag2MnCuO4, by substitution of one Cu atom in the structure with Mn atom. 51 The SEM image of such phase is shown in

I.2.2 Transport properties of new Ag-Cu oxides

Since the initially synthesized AgCuO2 was in powder form, to investigate the transport properties of the new phase, low density pellets were pressed, as reported in the work by F.

Sauvage et al. 43 In this work, both AgCuO2 and AgCu0.5Mn0.5O2 pellets were prepared and Considering the size and shape of the particles, a conductivity of 17 S.cm -1 was estimated (one order of magnitude higher than for the low density pellets reported previously).

I.2.3 Thin film deposition of Ag-Cu mixed oxides

The Ag-Cu oxides reported above were all initially synthesized as powder. For Ag2Cu2O3 compound, since it has less limitation to be synthesized in less oxidizing environment, there are several works already reported for its deposition as thin films. Most of the Ag2Cu2O3 thin films were deposited by sputtering or PLD [52][53][54] . In the work from J.F. Pierson et al., 55 the possibility of tuning the Ag-Cu ratio in the compound films by Reactive Sputtering technique was reported. In the forward scan, the current rose for potential higher than 0.58

V. In the backward scan, the current decreased down to even lower potential than 0.58 V.

This implies reduced slight irreversibility of the electrochemical oxidation reaction. Thus, among all the Ag-Cu oxides family, currently only Ag2Cu2O3 phase can be sputtered as a thin film, thus its electronic and optical properties have been evaluated and confirmed

that it has direct optical bandgap of 2.2 eV. The mobility can reach maximum 10 cm 2 /V.s and carrier density can get maximum 10 19 cm -3 . In the case of AgCuO2, due to a more critical oxidation environment required during its formation, currently no dense films are reported except nanoplates electrodeposited on ITO, therefore it requires more optimization or new thin film coating techniques in order to evaluate its electronic and optical properties.

I.3 Cu2O structures and electronic properties

Among the p type metal oxides previously introduced, Cu2O has drawn our attention, this is so since copper element is earth abundant, non-toxic and it can be processed through low cost chemical approaches. Cu2O has a direct bandgap around 2 eV, depending on the synthetic technique. Thus, it has a wide range of applications, especially for application in photovoltaics, the calculated theoretical efficiency for Cu2O based solar cells can reach an astonishing 20%. 56

23,57

As it is known, pure Cu2O without any doping is an intrinsic p type semiconductor oxide.

To further understand the origin of the p type conductivity, the work of H. Raebiger et al. 58 gave an explanation from the point view of formation energy. They pointed out that such intrinsic p type conductivity arises from the non-stoichiometry of the Cu2O lattice, by removing copper atoms and leaving an ionized copper vacancy VCu 0 , as described in Equation 1.

𝐶𝑢 𝑙𝑎𝑡𝑡𝑖𝑐𝑒 (𝐼) → 𝑉 𝐶𝑢 0 + ℎ + (1)
To confirm the origin of p type conduction in Cu2O, the formation enthalpies of different types of defects and vacancies were calculated and compared in copper rich/oxygen poor and copper poor/oxygen rich conditions, as presented in Lastly, one significant characteristic of Cu2O films is that the electronic properties, including even the bandgap, heavily rely on the synthesis technique used. As shown in Table I-2, Cu2O films can be fabricated through almost all major thin film coating techniques. Meanwhile the electronic properties obtained with the different techniques varies a lot. Among those different synthesis techniques, the best quality Cu2O films were produced through thermal oxidation of copper sheets and they presented high mobility over 100 cm 2 /V.s, thanks to the enormous grains size obtained (tens of microns). However, the daunting high temperature (above 1000 °C) for such oxidation process is very energy demanding. As well, with such high processing temperature, it's inconvenient and expensive for scalable production, and not well suited to device fabrication. On the other hand, Cu2O film with good electronic properties can be also synthesized through several physical deposition techniques, like Pulsed Laser Deposition (PLD) or radio frequency Magnetron Sputtering (rf-MS). 4,16,59 But, long deposition duration and critical high vacuum are usually required with these physical deposition techniques, making them still rather expensive and difficult to achieve scalable production.

Comparing with those techniques mentioned, chemical thin film coating techniques have shown their advantages in terms of low cost and scalability, especially with chemical solution based techniques. But, unfortunately, , due to the existence of higher density of defects in the grains boundaries, Cu2O films synthesized through solution approaches commonly have low electrical performance. For example, electrodeposited Cu2O films generally only reach carrier concentration about 10 14 cm -3 and very low mobility around 1 cm 2 /(V.s) . 60 Applying those types of Cu2O films in devices can severely damage the final performance. Considering both cost, scalability and thin film quality, Metal Organic Chemical Vapor Deposition (MOCVD) and related techniques seem like an optimal option.

We will discuss more in details about MOCVD synthesis and optimization of Cu2O in Aerosol Assisted MOCVD in Chapter IV. This difference of electronic properties can be partly attributed to the synthesis temperature.

In the work from E.Fortunato et al., [START_REF] Kim | p-Channel oxide thin film transistors using solution-processed copper oxide[END_REF] There are other reports about fabricating Cu2O p type channel TFTs through other fabrication techniques like spin coating [START_REF] Mahalingam | Galvanostatic deposition and characterization of cuprous oxide thin films[END_REF] , spray pyrolysis [START_REF] Pattanasattayavong | P-channel thin-film transistors based on spray-coated Cu2O films[END_REF] , etc. But due to the undermined electronic property of synthesized Cu2O, on off ratio only reach 10 2 and field effect mobility just reaches 0.01 cm 2 / V.s. Therefore, due to a higher mobility of Cu2O produced through several physical deposition techniques, including rf-MS and PLD, the final devices generally have better performance. However, those physical deposition process are rather expensive and time demanding, thus lower cost and more rapid chemical deposition approaches should be optimized.

I.3.2 Application Cu2O in water splitting

Besides applying Cu2O layers for transistors, Cu2O is also very interesting for water splitting via photocatalysis, also called photo electrochemical In addition to using Cu2O as photoelectrode and solving its stability issues, in the work of Y.F. Lim et al., [START_REF] Lim | Sol-gel deposited Cu2O and CuO thin films for photocatalytic water splitting[END_REF] CuO was also under study for PECs cells with comparison of Cu2O. Both phases were achieved through sol gel spin coating and followed with annealing in different atmosphere under 500 °C. Although CuO has an indirect bandgap about 1.2 eV, it has shown a higher photocurrent of -0.35 mA cm -2 (Cu2O -0.28 mA cm -2 ) and better stability in liquid. In addition to the works on Cu2O stability though using composite layers, in the work of Q.B.

But

Ma et al., [START_REF] Ma | Cu2O photoelectrodes for solar water splitting: Tuning photoelectrochemical performance by controlled faceting[END_REF] the facet preferential orientation of Cu2O could be tuned between (200) plane and (111) plane during the process of electrochemical deposition thus obtaining a different performance of Cu2O as photo electrode. Cu2O films with (111) texture exhibits a better stability without corrosion in PEC test. Besides using Cu2O films as PEC cell electrode, using Cu2O nanowires as electrodes was also reported. [START_REF] Luo | Cu2O Nanowire Photocathodes for Efficient and Durable Solar Water Splitting[END_REF][START_REF] Shi | Carbon coated Cu2O nanowires for photo-electrochemical water splitting with enhanced activity[END_REF] In those works, unprecedentedly high photon current of 10 mA cm -2 was reported. At the same time with an excellent stable performance over 50 h test. Thus, even though Cu2O has shown great potential in photo electrochemical cells for hydrogen generation, stability of such electrodes need to be carefully tackled.

I.3.3 Application of Cu2O in photovoltaics

In addition to the water splitting, by forming heterojunction normally with ZnO or AZO (Al doped ZnO), Cu2O is also applied as the absorber layer in ultra low-cost all oxide solar cells.

With AZO (or ZO)/ Cu2O structure, the highest efficiency reached so far is merely 3.83% using thermally oxidized Cu2O by T. Minami et al. 18 Other work from the same group, by using thermally oxidized Cu2O film coupled with ZnO/Ga2O3 as the n layer, the highest efficiency achieved was 5.38%, with structure ZnO/Ga2O3/Cu2O. There are several reasons leading to the poor performance of the Cu2O based photovoltaic devices (as said before, maximum theoretical efficiency of Cu2O based cells is 20% [START_REF] Ievskaya | Fabrication of ZnO/Cu2O heterojunctions in atmospheric conditions: Improved interface quality and solar cell performance[END_REF][START_REF] Lee | Hall mobility of cuprous oxide thin films deposited by reactive direct-current magnetron sputtering[END_REF][START_REF] Chen | Fabrication and optical properties of Cu2O--ZnO composite opal[END_REF] . To understand the actual problems with those devices, we will elaborate those issues from the point view of Cu2O materials itself.

First of all, the effects from the electronic properties of Cu2O on the final solar cell performance are discussed. As has been stated above, the properties of Cu2O are very dependent on the deposition method and conditions. With electrodeposited Cu2O layers in solar cells, for example, the incompatibility between charge collection length and depletion zone is reported by K.Mussleman et al. [START_REF] Musselman | Incompatible Length Scales in Nanostructured Cu2O Solar Cells[END_REF] As shown in

𝑉 𝑏𝑖(𝑛) 𝑉 𝑏𝑖(𝑝) = 𝑁 𝐴 𝜀 𝐶𝑢2𝑂 𝑁 𝐷 𝜀 𝑍𝑛𝑂 (3)
In which, NA ad ND represent the acceptor and donor concentration in p and n type semiconductors , respectively. ε with different subscript are the absolute permittivity with values 6.2 for ZnO and 8 for Cu2O . Due to the fact that most electrodeposited Cu2O has poor carrier mobility (~ 1 cm 2 /V.s), effective charge collection length in such material is less than 1 μm. [START_REF] Musselman | Incompatible Length Scales in Nanostructured Cu2O Solar Cells[END_REF] On the other hand, such Cu2O has relatively lower carrier concentration (in the range of 10 13 -10 14 cm -3 ).By comparing the general ZnO carrier density in the range of 10 18

-10 20 cm -3 , the relation between build in potential in this type of Cu2O and full depletion length can be expressed in Equation 4:

𝑉 𝑏𝑖(𝑝) = 𝑞𝑁 𝐴 𝑥 𝑝 2 2𝜀 𝐶𝑢2𝑂 (4)
Where q is the unit electron charge and xp is the full depletion length in Cu2O. Therefore, according to Equation 4, in order to obtain a build-in potential in the range of 0.4 to 0.7 V, the full depletion length is estimated to be about 2. ), the shot circuit current with bilayer structure showed a peak. The increase of current could be attributed to more absorbed light thanks to a thicker absorber layer. On the other hand, the decrease of the current can be explained by that Cu2O layer was thicker than the effective charge collection length. Thus, even though more excitons were generated in the active layer, there was less possibility for the excitons to be separated and collected.

Meanwhile, in figure c) the evolution of the solar cell efficiency presented a similar trend. To solve this incompatibility between charge collection and full depletion length, better thin film coating techniques could be helpful. From the work of A.T.Marin et al., 61 Comparing with electrodeposited Cu2O, Cu2O used in this work has higher carrier concentration (2 orders higher than electrodeposited). Thus, to achieve 0.7 V build-in potential, by calculation using Equation 3 and 4, the full depletion length in this case was estimated about merely 250~ 300 nm. In Besides the work on Cu2O material its self, other works are focusing on the ZnO or AZO layer. Such as in the work of T.S.Gershon et al., [START_REF] Gershon | Chapter 1.Introduction II.4 References[END_REF] the fill factor of Cu2O/ZnO bilayer solar cell was improved up to 54% through optimizing ZnO surface morphology with larger grains.

With application of new deposition techniques such as atomic layer deposition, electronic property of ZnO/AZO can be also largely improved 82,37 . Therefore, Cu2O thin films have

shown a wide range of potential applications. For either application of TFT, water splitting or solar cells, with a better quality Cu2O film, the device performance can be always improved by tuning the deposition of Cu2O. Thus, concerning both cost and deposition quality, a proper thin film deposition technique for Cu2O should be chosen carefully.

Although the great potential presented with Cu2O, the material has its own limitations.

Especially for photovoltaic applications, the bandgap of Cu2O is not ideal enough as an absorber. For the sun light spectrum, absorber layers with a lower bandgap (about 1.5 eV) are more efficient to convert the majority of the photon energy into electricity. From this point of view, new emerging silver copper mixed oxides have much lower bandgaps, thus they might be more suitable for photovoltaic applications.

I.4 Objective of PhD thesis

With the rising demands for p type metal oxides with better performance, the main objective of this paper work is using low temperature, chemical approaches to deposit p type metal oxides with improved electronic properties. The new AgCuO2 phase has shown p type tranport nature with low resistivity. As well, from the DFT calculation, low bandgap about 1 eV was predicted. Given the promising characteristics of AgCuO2, this work will focus on the synthesis of AgCuO2 thin films. To achieve that, several approaches, namely MOCVD, AgCu alloy film oxidation and solution coating method, will be evaluated.

For the deposition of AgCuO2 by MOCVD, first the deposition of silver and silver oxide films, on the one hand, and of copper oxides, on the other hand, will have to be optimized.

Then, the co-deposition of mixed oxide phases will be parallel, given the relevance of Cu2O, the deposition conditions of Cu2O thin films by MOCVD will also be optimized.

For the oxidation of AgCu foam films, electrochemical and ozone oxidation wil be adopted.

Finally, the Successive Ionic Layer Adsorption and Reaction (SILAR) method will be evaluated for the direct deposition of AgCuO2 thin films. The governing condition will be optimized to achieve continuous and dense AgCuO2 coatings. The synthetic AgCuO2 thin films, will allow to determine the electronic properties and optical bandgap of AgCuO2. ). In the following section, a brief introduction to AA-MOCVD and PI-MOCVD principles, systems setups and the deposition process will be introduced.

II.1.1 Pulsed Injection MOCVD

PI-MOCVD is based on the flash evaporation of precursors or precursor solutions. This is carried out through the pulsed injection of small droplets of the precursor solution in a chamber under vacuum. In this way, the deposition rate is not dependent on the volatility of the precursor itself, and larger range of precursors can be available when compared with other MOCVD technqiues. 2,3 Due to the low quantity of precursor required for an efficient deposition, PI-MOCVD can allow a rapid precursor screening with small precursor consumption (about 25 ml of 0.01 M precursor solution for each test). Therefore, for the sake of convenience, depositions with more expensive or less volatile precursors were conducted via PI-MOCVD in this work. andc). Two injector lines, 1 and 2, are present and Ar is used as carrier gas. Oxygen is fed via another line through the top of the reactor. The vacuum chamber can reach a minimum pressure of 0.1 mbar. The heating plate is a Si wafer with diameter about 8 cm that is illuminated by an ultra violet lamp placed below (the temperature of the lamp is controlled by adjusting the lamp power).

Fast heating rates can be achieved and the temperature is monitored with an infrared detector on the right-top corner of the deposition chamber and controlled by adjusting the lamp power.

As shown in Figure c) right top corner, uniform depositions over a 5 cm diameter circumference can be achieved with our setup.

II.1.2 Aerosol Assisted MOCVD

AA-MOCVD is well known for the ability to use less volatile metal organic precursors during CVD deposition at atmospheric pressure (no vacuum). 4 The deposition process is shown in Figure II-2, in which the precursor solution is atomized into small droplets (diameter about several microns) by a piezoelectric ceramic and the droplets are transported by the carrier gas into deposition chamber. Before the droplets reach hot substrate, the solvent is firstly evaporated due to the temperature gradient close to the heating source and the precursor molecules are further evaporated when the flow is closer to the hot substrate.

Lastly the vaporized precursors are decomposed on the hot substrate thus dense and continuous thin films can be deposited. In this work, AA-MOCVD has been mostly used for the deposition of Cu2O thin films, although Ag deposition has also been tested. Ag coating was indeed also successfully carried out by this approach but the large precursor consumption makes it less convenient than PI-MOCVD, as stated above, when precursors are expensive. The AA-MOCVD system used during this work is presented in its initial configuration in Figure II-3. The precursor solution is subjected to ultrasounds generated by a piezoelectric ceramic and atomized droplets are formed. The mist is then carried by N2 carrier gas (CG#1) from the bottom. For the sake of security, oxygen or air are only introduced from the middle part, accelerating the precursor onto the top heating plate, where the reaction takes place. The decomposed precursor byproducts and evaporated solvent are ventilated from the top, going through a solvent trap cooled with liquid nitrogen.

It has to be noticed that, in our AA-MOCVD systems, in order to ensure the same amount of precursor consumption throughout the deposition process, the precursor solution level in the chamber has to be kept stable. In our case, this was realized by connecting the solution chamber with a specially designed small reservoir as shown in Figure II-3. Due to the balance of the pressure, the solution level in the mist chamber was the same as the level in the small reservoir. By shifting the position of the reservoir up and down, the level of precursor solution in the mist chamber was adjusted and maintained stable.

Lastly, an atmospheric plasma system was connected with the deposition chamber in the AA-MOCVD, which was used for providing highly oxidizing species in this work. The working principle of the atmospheric plasma system is introduced in section II.2.2. The AA-CVD system went through several modifications in order to improve the Cu2O thin film deposition quality, which will be discussed in details in Chapter IV. The schematic illustration of the final configuration of our AA-MOCVD system is shown in Figure II-4 a).

In this configuration, the hot plate is placed below the precursor flow. this thesis) 6 . The precursor was carefully grinded in a ceramic mortar before dissolving it in alcohol solvent. 0.01 M precursor solutions were prepared by adding 2.45 g (0.005 mol) of precursor AgP in 500 ml ethanol solvent with magnetic bar stirring at a speed of 300 cycle/min for 60 min. Again, to speed up the process, dispersants as mentioned above can be added.

II.1.3.2 Substrate cleaning

Glass or Si substrates were cut with a diamond pen into squares with size of about 2.5 x 5 cm 2 . The substrates were washed in a soap solution with ultrasounds for 10 mins, rinsed with deionized water and dried with a N2 flow. They were then washed in an ultrasonic bath again with acetone and isopropanol solvents in sequence, each for 10 mins at 40 °C. Lastly, they were dried with a N2 flow and kept in a clean plastic box.

II.1.3.3 Preparation of the MOCVD systems PI-MOCVD:

First of all, the injector was cleaned with the same solvent to be used, in this case ethanol or butanol-01. The injector was filled with 50 ml of precursor solution and fixed onto the precursor line. The temperature of the vacuum precursor line was adjusted to prevent the condensation of precursor before reaching the deposition chamber. The particular deposition conditions for each sample are detailed and discussed in the following chapters.

AA-MOCVD:

In MOCVD, the distribution of temperature on the heating plate has a huge impact on thin film homogeneity. To ensure a homogeneous heat distribution on the substrate, a Si wafer with a diameter of 10 cm was placed between the substrate and heating plate. The water cooling on the atmospheric plasma head needs to be always on during deposition, even when not using it, to avoid overheating. Meanwhile, before starting the deposition, the precursor solution in the atomization chamber is pre-warmed by switching on the piezoelectric ceramic and applying a power of 6 W, with frequency 5 Hz. Since the mist formation is dependent on temperature, this ensures that the mist is the same during the whole reaction.

The ventilation of the system is monitored by the pressure gauge, as shown in Figure II-5, connecting with the deposition chamber on one side and the other exposed to atmospheric pressure. By changing the value of the gauge between 0 and 300 (maximum), the ventilation can be adjusted. Without ventilation the pressure gauge is at zero and the higher the ventilation power is, the higher the value on the pressure gauge. The ventilation used in this work was adjusted by setting the pressure gauge to about 3 cm with 10 L/min of carrier gas entering the deposition chamber. At least during 5 minutes before the deposition actual start, the gases with the correct ratio for the deposition need to be pre-introduced in the deposition chamber.

II.2 Oxidation techniques: electrochemical oxidation and oxidation with atmospheric oxygen plasma II.2.1 Electrochemical deposition and oxidation

In the majority of cases, as illustrated in Figure II-6, an electrochemistry system is made up of a potentiostat, and a cell containing the electrolyte and three electrodes, namely, working electrode (WE), counter electrode (CE) and reference electrode (RE). Among those electrodes, WE is where the expected reaction happens and RE is the electrode to fix the reference potential for the WE independent from the change of the solutions. 7 During the electrochemical reaction process, voltage or current can be applied on the WE with the potentiostat. If the potential on the WE is positive, then WE acts as the anode, where electrons are removed and oxidation takes place. Conversely, the CE is the cathode. Where electrons are supplied and thus reduction takes place. the forward curve represents an oxidation process at bias Ep a . For reversible reactions, the backward curve shows a reduction reaction at potential Ep c . The degree of reversibility of the electrochemical process is described by the difference between Ep a and EP c , the bigger the difference the less reversible the reaction is. In addition, the ratio of intensity ip a and ip c should be close to 1:1 for a reversible reaction. In the case of oxide deposition or oxidation, after identifying reaction potentials with LSV or CV, a particular oxidation state of the metal ions can be achieved by applying the necessary voltage (i.e. by performing a CA, were current is monitored). Reactions can be carries out in CP mode, where the current is set to a certain value as required and the potential evolution is monitored.

II.2.2 Oxygen plasma oxidation

Another oxidation technique adopted in this work has been atmospheric oxygen plasma oxidation. The working principle of atmospheric plasma is illustrated in Figure II-9, in which the gas flow passes through the two electrodes with high voltage applied in between. The ultra-high potential between the two electrodes ionizes the gas molecules passing through.

Thus, the emission plasma flow contains electrons, ions and radicals. The plasma generator is set at the bottom of the deposition chamber besides the carrier gas flow tube.

As it is shown, Ag/Cu metallic films were placed on the heating plate and subjected to the plasma with different powers (up to 500 W, with an O2 flow of 15 L/min). During the oxidation process, H2O mist was introduced at the same time to reduce corrosive radicals emitted from the plasma head.

The effect of temperature was also studied during plasma oxidation. The results are detailed in Chapter III and Chapter V.

II.2.3 Deposition of AgCuO2 films by the SILAR method

Successive

𝑚𝑀 𝑛+ (𝑎𝑞) + 𝑛𝐴 𝑚-(𝑎𝑞) → 𝑀 𝑚 𝐴 𝑛 (𝑠) (5)
Thus, theoretically SILAR deposition can be conducted in a layer-by-layer manner, yielding dense and high quality films. -the elastic scattering, known as the Rayleigh scattering, where the incident and scattered photons have the same energy, meaning that there is no energy transfer during the process.

-the inelastic scattering where the scattered photon has an energy different from that of the incident photon. Raman scattering is one of the most important inelastic scattering processes;

it involves transitions between the vibrational/rotational levels due to the combination of an absorption of the incident photon through a transition to a higher energy level, represented as an intermediate virtual state, followed by a spontaneous emission of the scattered photon.

If the scattered photon has less energy than the incident one, the molecule is excited to a higher energy level and the process is called the Stokes Raman process. It corresponds to the excitation of the molecule from the vibrational ground state (v=0) to the first excited state The main advantage of Raman spectroscopy is that it is a local probe; moreover, it is non destructive, cost-effective and doesn't need sample preparation. It gives complementary information to Infrared (IR) spectroscopy which also probes the vibrational modes through an absorption process where the molecule is excited to a higher vibrational level within a given electronic state, usually the ground state.

Raman spectroscopy has mainly been used in this work for the identification of crystal phases and detection of carbon species in the thin films. In addition, Raman intensities also gave us information about the crystallinity of the materials.

Raman spectra were recorded from the film surface using a Jobin Yvon/Horiba Labram spectrometer equipped with a charge-coupled device (CCD) detector cooled with liquid nitrogen. The experiments were conducted in the micro-Raman mode at room temperature.

Among the three exciting wavelengths available in LMGP (red, blue and green), we used the blue (at 488 nm) and green (514.5 nm) ones Ar 1+ laser with a power as low as ranging from 60 W to 6 W (filter 1/100 to 1/1000) at the sample surface to avoid sample heating.

The laser was focused to a spot size close to 1 m on the sample with using a × 100 or × 50 objective. Before each spectrum acquisition, the Raman line of a Si substrate was measured and used as reference for calibration (theoretical position at 520.7 cm -1 ).

II.3.1.2 X-Ray diffraction (XRD): Bragg-Brentano and Grazing Incidence.

X-Ray Diffraction is a widely used characterization technique for studying crystalline matter, providing information including phase identification, texture and grain size, etc. The basic working principle of this technique relies on the Bragg diffraction. When X-Ray radiation (having wave length similar to the distances between atoms in crystals) reaches a crystalline material, it is scattered by the electrons around each atom. Depending on the angle between the incident radiation and the atomic plane, collective constructive scattering can take place. This is expressed by Bragg's equation (Shown in Equation II 6):

2d hkl × sinθ = nλ ( 6
)
Where dhkl is the distance between two crystal planes with Miller index (hkl), θ is the incident angle between the X-Ray and the (hkl) crystal plane, n is the diffraction order and λ is the X-Ray wavelength.

The X-Ray Bragg diffraction is schematically presented in Figure II-13. In the figure, the scattering vector is defined as K = keki as the difference between the wave vector ke of the scattered wave, and the wave vector ki of the incident X-ray. In Bragg-Brentano geometry, during XRD acquisition the X-Ray source is fixed and both sample plane and detector turn along the same rotation axis. θ is defined as the angle between incident X-Ray beam and the sample surface while 2θ is the resulting angle between incident beam and scattered beam (the detector is moved so that it keeps the same θ angle with respect to the sample, thus 2θ with respect to the incident beam). In this work, a Bruker D8 Advance diffractometer was used for acquiring XRD patterns with Bragg-Brentano Geometry, with wavelength 1.5406 Å from Cu Kα1 radiation. A 1D detector with resolution about 0.05° was used for signal collection. 2.5° and 0.3° Soller slits were used for divergent slit, while anti-scatter slits were 6.8 mm width.

For very thin films (< 10 nm) Grazing Incidence geometry (shown in Figure II-16) can be applied. 12 Grazing incidence (GI) geometry is very sensitive to phase composition even for samples with less radiation volume, like thin films. With the Bragg-Brentano geometry, the incident X-ray can easily penetrate through the thin film material, the resulting collected signal being too weak to be distinguished from the background noise, especially with an amorphous substrate. In GI geometry, the incident X-ray beam is aligned to the sample surface with a very small incident angle, typically ranging from 0.25° -0.5°. 15 The positions of X-ray source and sample stage are fixed, only the detector scans in 2θ angle to collect the diffracted signal. With such a close to parallel beam, larger amount of thin film material can be radiated thus stronger diffraction signal can be acquired. In this way, the interference signal from the substrate can be minimized. But it has to be noticed that since the X-Ray In addition to phase identification, the XRD patterns acquired under Bragg-Brentano geometry can also give information about grain orientation in the thin film (i.e. texture).

Such a preference or texture can be evaluated using a texture coefficient. 16 The texture coefficient Tc can be calculated with Equation II 7.

T c (h 𝑖 k 𝑖 l 𝑖 ) = I(h 𝑖 k 𝑖 l 𝑖 ) I 0 (h 𝑖 k 𝑖 l 𝑖 ) ⁄ ( 1 n ⁄ )[∑ I(h 𝑗 k 𝑗 l 𝑗 ) I 0 (h 𝑗 k 𝑗 l 𝑗 ) 𝑛 𝑗=1 ] (7)
Where Tc(hikili)is the texture coefficient of the specific (hikili) plane, I is the measured diffraction intensity (under Bragg Brentano geometry) for a particular Miller index, I0 is the intensity for the same index in the reference pattern and n is the number of reflections taken into account. Indexes with Tc values bigger than 1 indicate that there is a preferred orientation along the respective direction Besides texture, by measuring the Full Width at Half Maximum (FWHW) of the reflections, the crystallite size on each particular direction can be estimated through the Scherrer Equation, 17,18 as shown in Equation II 8. 19

β L = Kλ L cosθ ( 8
)
Where βL is the peak width (obtained by subtracting the width associated to the instrument),

K is a numerical constant to account for instrumental widening (0.9 in our case), λ is the wavelength of the X-Ray source, L is the average crystallite size and θ is the incident angle.

II.3.2 Morphological characterization: SEM: EDS, EBSD, AFM II.3.2.1 Scanning Electron Microscopy (SEM)

Scanning Electron Microscopy is a widely used characterization technique for acquiring topographical information on the sample surface. It is also a versatile analysis technique providing qualitative and semi-quantitative information about sample composition. Finally, it can also provide information about crystal structure and texture. In SEM, a primary electron beam is allocated towards the sample surface and causes a series of phenomena (see When the probe tip scans across the surface, any change in the sample morphology will cause the deflection of the cantilever thus a change in the reflection of the laser spot on the detector. To maintain the laser spot at the center of the detector, a feedback signal will be applied on the tip cantilever through a piezoelectric piece in the Z axis. Thus, the morphology change of the sample surface is converted into electrical signal and recorded to recreate the surface 2/3 D image afterwards.

There are mainly three types of AFM scanning modes, namely, contact mode, tapping mode and non-contact mode (or vibration mode). In contact mode, the AFM tip is dragged along the surface of the sample, the repulsive force between the atoms from the tip and sample causing a deflection of the cantilever. Such imaging mode shows better resolution than the other two imaging modes but some fragile or weakly attached samples are not resistant to such dragging force. Thus the other semi/non-contact modes will be more appropriate. The other imaging mode is called tapping mode or semi-contact mode, in which the cantilever is driven by the piezoelectric with a certain frequency and amplitude. With closer interaction with the surface, the amplitude of the oscillation will be affected thus recorded as the morphological information. Lastly, by taking advantage of the Van Der Waals force interaction between the tip and sample atoms, both the oscillation and amplitude modulation are the characterization approaches in non-contact mode. Since there is no physical contact between the tip and sample, non-contact mode is preferable when dealing with soft samples.

In this work, the AFM system used was a Digital Instruments D3100 Nano scope, working in tapping mode. Si and Ti coated Si probes were used.

II.3.2.3 Transmission Electron Microscopy (TEM)

Transmission Electron Microscopy allows both sample imaging and in situ crystallographic analysis. Its basic working principle is built on the interactions between the electrons in the material and the e-beam crossing the sample (thickness has to be less than 100 nm) to form images or diffraction patterns. This requires preparation of samples thin enough for characterization. Therefore TEM is a destructive characterization technique, but it can allow observing the fine details of the sample on atomic scale. In some cases, powder samples can also be studied and in this case very little amount of sample is needed.

For thin film characterization in this work, there were two ways to prepare the samples, mainly by scratching the surface to get small clusters of the thin film material or by cutting and mechanical milling followed by polishing the cross-section with an Ar + beam. The scratching methods allows rapid preparation but it may take a lot of time during the characterization process to find the parts of interest. Preparation of cross-sections can take much longer but it is much easier to find the area of interests. In a typical TEM

II.3.3 Ultraviolet-Visible (UV-Vis) spectroscopy

UV-Vis spectroscopy is mainly used for optical characterization. In this work, a Perkin Elmer Lamda 950 spectrometer was used, equipped with a deuterium and a tungsten lamps to emit light with different wavelengths ranging from 250 nm to 2500 nm. For signal collection, two photon detectors work in sequence switching at wavelength 850 nm (a photomultiplier for UV and visible light; An InGaAs sensor for infrared). The spectrometer is equipped with an integrating sphere and an automated reflectance transmittance analyser (ARTA). Therefore, transmittance, reflectance and absorbance spectra can be acquired. (image from: https://www.perkinelmer.com)

Both total transmittance/reflectance or direct transmittance/reflectance can be measured.

𝑇 𝑎 = 1 -T 𝑡 -T 𝑟 (9)
Based on the optical measurements, the absorption coefficient can be approximately calculated following Equation II 10 and 11, 24 𝑇 ≈ (1 -𝑅) 2 𝑒 -𝛼𝑑 (10)

𝛼 = - 1 𝑑 ln ( 𝑇 (1 -𝑅) 2 ) (11)
Thus, optical bandgap of the thin film semiconductor material can be identified using the Tauc plot as shown in Equation II 12, 24 (𝛼ℎ𝜈) 1/𝑛 = 𝐴(ℎ𝜈 -𝐸 𝑔 ) (12) in which α is the absorption coefficient, h the plank constant, ν is the photon frequency and Eg is the bandgap. The value of n depends on the nature of the bandgap 25,26 : 3) 27 . Thus, the resistance can be calculated as:

direct
R = V I = ρ L Wt = ρ t L W (13)
in which R is the resistivity, V is the voltage, I is the current, ρ is the conductivity, W is the width, L is the length and t is the thickness. The sheet resistance Rs (unit Ω/□) for thin films can be obtained from ρ and the film thickness (t), as shown in Equation II 10.

R S = ρ t (14)
Therefore,

V I = R = R s L W (15)
R s = 𝑊 𝐿 • 𝑉 𝐼 = k • V I ( 16
)
Where k is a correction factor based on sample size and space between each probe, which is taken as π/ln2 in this work for thin film measurements. 28,29 Thus, Rs can be rewritten as in Equation II 13.

R s = π ln2 × V I = 4.53 V I (17)

II.3.4.2 Hall Effect measurements

In addition to the four point probe measurement, the Hall Effect measurement can allow direct measurement of more detailed electronic properties besides resistivity, including identifying the type of semiconducting material, carrier concentration and mobility.

As it is shown, the working principle of Hall effect measurement is based on the Lorentz force generated when passing charged particles (electrons in this case) through an electric field while applying a perpendicular (Z axis) magnetic field: qvE × BZ, where qv is the carried charge, E is the electrical field and BZ is the magnetic field in the Z axis. In a semiconducting material, due to the unequal quantity of mobile charges (electrons and holes) present, when the current is applied along the X axis (IX), a potential difference appears in the Y axis, which is defined as the Hall voltage (VH), from which carrier concentration (n)

and mobility (μ) can be obtained from Equation II 18 and 19, provided the Rs value is known. 30,31

n = I X B Z eV H t (18)
μ = 1 eR s nt ( 19
)
where t is the thin film thickness and e is the elemental charge. to the high resistivity of Cu2O films, a current amplifier was applied for a better accuracy.

To ensure an ohmic contact between the measured Cu2O films and the Pt tips, gold contacts were evaporated on each corner on the 0.5 × 0.5 cm 2 square sample.

Chapter In this chapter, first of all, the deposition of metallic silver films using different precursors in both PI-MOCVD and AA-MOCVD will be introduced. With commercialized silver precursors, a cheap cost silver precursor, silver acetate, has been firstly reported for MOCVD deposition. In addition, two brand new silver metal organic precursors, Ag(hfac)phenanthroline and Ag(hfac)triglyme, will be introduced. The deposition results from those precursors are characterized. Secondly, the process to obtain silver oxides films are also included in this part of work, two oxidation techniques are adopted and the results will be presented.

Chapter III Thin film deposition of silver and silver oxides

Silver is a noble metal with unique physical properties including high conductivity and plasmonic, anti-bacterial and self-cleaning effects. [1][2][3] There are several techniques used for depositing Ag coatings, such as electroless plating, magnetron sputtering and sol gel, etc. 4, 5, the same method was used to optimize the deposition of Ag and silver oxides. As shown in the Chapter I, MOCVD is an industrially relevant technique for obtaining functional coatings, including metallic Ag coatings. Due to the noble nature of Ag, a high oxidation environment is required for further oxidation into silver oxide films. In this chapter, several types of silver precursors were tested including two newly designed silver metalorganic precursors. Ag depositions could be achieved with the commercially available and cost-efficient precursor silver acetate (AgAc), in specific MOCVD systems, and with the new metal organic precursors. Better quality and uniformity were achieved in both AA-MOCVD and PI-MOCVD when using the two new precursors. Finally, different approaches to oxidize silver films deposited by MOCVD are also detailed, including electrochemical oxidation and oxygen plasma treatment.

III.1 Ag Precursor design, synthesis and characterization

As indicated in Chapter II, in MOCVD the process can be largely altered by the way in which precursors are designed and delivered. Thus, depending on how precursors are fed into the deposition chamber, MOCVD is classified into several branches: HW-MOCVD, CW-MOCVD, PI-MOCVD, AA-MOCVD, etc. 7 This big family of MOCVD systems also brings new challenges for the design of new silver precursors. In particular, AA-MOCVD is an interesting CVD approach since it operates at ambient pressure. For AA-MOCVD, alcohol solvents are preferred due to the concerns related to the use of other more hazardous solvents, thus precursors being soluble in alcohol solvents are desirable. 8,9 In the case of the currently known silver metalorganic precursors, solubility in alcohols is low since silver complex molecules tend to polymerize when dissolved in the solvent. This limits the number of available silver precursors suitable for AA-MOCVD systems. AgNO3 has indeed been reported as precursor for Ag coatings using AA-MOCVD by Ponja et al. and also in a pulsedspray-evaporation CVD system by N. Bahlawane et al. 10,11 But given the nonorganometallic nature of AgNO3, the utilization of such precursor may be limited to some type of CVD reactors. In addition, the development of new organometallic Ag precursors is also interesting, for instance for application in ALD, where suitable Ag precursors are still needed, especially for thermal ALD 12 . As shown in Table III-1, the different silver metalorganic precursors reported to date were mostly used in conventional or PI-MOCVD systems. Few examples of ALD of Ag coatings have also been reported previously. [13][14][15][16] In the cases where AA-MOCVD was used, non-environmentally friendly solvents, such as THF and toluene, were used in combination with organometallic precursors. [17][18][19] Concerning the design of Ag metalorganic precursors, previous works have introduced different systems, including β-diketonate adducts Ag(hfac))(L) 20 or Ag(fod)(L), where L were tertiary phosphines (PPh3), 2,5,8,11,14-pentaoxapentadecane or bis(trimethylsilyl)-acetylene) or carboxylates (Ag(O2C-L)). 19,21,22 Apart from the low solubility of these precursors in alcohol solvents, other issues such as low volatility or the high toxicity of the ligands have been also reported. 23,16,20 In addition, some of the ligands are rather expensive, contributing to a higher price of the final Ag precursor. Thus, new silver precursors with less toxic ligands, good solubility in alcohol solvents and clean rapid decomposition, are still needed, ideally using less expensive ligands.

Synthesis of [Ag(hfac)(phen)] (AgP):

The synthesis was performed suspending 2.20 g of Ag2O (9.49 mmol) in 50 ml of dichloromethane. 2.54 g of Phenanthroline (14.14 mmol) were then added to the suspension.

After 10 min, 2.94 g of H-hfac (14.14 mmol) were added to the mixture, which was refluxed under stirring for 1 h as shown in Figure III-1. In order to ensure the full consumption of the ligand, excess silver oxide was added and the unreacted oxide was filtered off after the reaction. Lastly, yellow crystals were obtained after removal of the solvent upon evaporation.

A pure crystalline solid was obtained after purification with small quantities of n-pentane, in which the adduct is insoluble, and then vacuum dried.

Ag 2 O + 2 H -hfac + 2 phen → 2 Ag(hfac)phen + H 2 O ( 20
)
𝐴𝑔 2 𝑂 + 2 𝐻 -ℎ𝑓𝑎𝑐 + 2 𝑡𝑟𝑖𝑔𝑙𝑦𝑚𝑒 → [𝐴𝑔(𝑡𝑟𝑖𝑔𝑙𝑦𝑚𝑒) 2 ] + [𝐴𝑔(ℎ𝑓𝑎𝑐) 2 ] -+ 𝐻 2 𝑂 (21)
Precursor characterizations:

Elemental microanalysis were performed using a Carlo Erba 1106 elemental analyzer.

Fourier Transform Infrared Spectroscopy (FT-IR) was done using a Jasco FT/IR-430 spectrometer with nujol mulls between NaCl plates. 1 H NMR (400.13 MHz) and 13 C NMR (100.61 MHz) were acquired on a Bruker AvanceTM 400 spectrometer. Chemical shifts (δ) are expressed in parts per million (ppm). Spectra were referenced to the residual proton solvent peaks; coupling constant (J) values are given in Hz. Thermogravimetric analyses were made using a Mettler Toledo TGA/SDTA 851e. Dynamic thermal analyses were carried out under purified nitrogen flow (30 sccm) at atmospheric pressure with a 5 °C/min heating rate. Weights of the samples were between 10-15 mg. Melting points were taken on very small single crystals using a Kofler hot-stage microscope.

Ag coating characterizations:

The Ag coatings were analysed using a Scanning Electron Microscope (SEM-FEG Environmental FEI QUANTA 250). Particle diameter distribution was calculated from SEM images using the Image J software. 28 The Ag coating were scratched for further characterization with Transmission Electron Microscopy (TEM-JEM2010-JEOL). Carbon residues were analyzed with RAMAN Spectroscopy (Jobin Yvon/Horiba LabRam) equipped with a liquid nitrogen cooled charge coupled device detector. Structures were solved by using the SIR-2004 package 30 and subsequently refined on the F2 values by the full-matrix least-squares program SHELXL-2013. 31 Geometrical calculations were performed by PARST97 32 and molecular plots were produced by the program Mercury (v3.7). 33 The crystal data and refinement parameters of the new Ag adducts are reported in Table III- The solid state molecular structure of AgP consists of dimers made up of two silver complexes related by an inversion center (-x+1, -y, -z+2), shown in Figure III-2 a). In each complex the silver ion is bound to the oxygen and nitrogen atoms provided by the -hfac and -phen ligands, respectively. The coordination sphere is a slightly distorted square planar (as provided by the angle between the AgO2 and AgN2 planes, 7.2(2)°, and the small displacement of the silver ion with respect to the four donor atoms, 0.1110(5) Å). Even if quite uncommon 34 , this kind of geometry has been observed in similar Ag(I) complexes. 35 Bond distances and angles are within the expected range for this kind of compounds (Table III-3). 37 Additionally, as already found in other [Ag(hfac)-L] complexes, 36,38 the Ag-O bond distances are significantly different, Ag(1)-O(1) = 2.276(5) vs Ag(1)-O(2) = 2.439(3). The distance separating the two metal centres (3.0843(7) Å) is definitely longer than the silversilver one in the metal (2.888 Å). 39 As a consequence, the two metal ions do not interact significantly. 38,40 However, given the fact that there are no bridging ligands holding together the dimer, a weak stabilizing interaction between the metal ions might be postulated. 41 In addition, intra-dimer π stacking appeared between the -hfac and -phen rings (the distances separating the centroids (CT), centroid of the hfa ring as defined by O(1)C( 2 III-3), it is inside the Ag-O range as retrieved in the Cambridge Structural Database (CSD v. 5.37). 43 The two triglyme molecules are almost perpendicular to each other, the angle between the mean planes defined by the non-hydrogen atoms of the two molecules being 85.4(1)°. Finally the resulting coordination polyhedron could be described as a distorted dodecahedron.

III.1.2 XRD analysis and structure elucidation

Concerning the [Ag(hfac)2] -anion, the silver atom Ag( 1) is four-coordinated by the oxygen atoms of two -hfac anions in a resulting pseudo-tetrahedral arrangement. This coordination arrangement resembles well that already found in the similar silver complex

III.1.3 New Ag metalorganic precursor molecular structures and characterizations

In both cases, the synthesis was reproducible and yielded an anhydrous product. Adducts were soluble in dichloromethane and could be isolated as yellow crystalline powders by removal of the solvent. The silver precursors were characterized by FT-IR spectroscopy in the wave number range of 4000-500 cm -1 . The results for AgP and AgT are shown in Figure III-4 a) and b) respectively. In both cases, the FT-IR spectra do not show bands in the 3300-3600 cm -1 interval, which indicates the absence of water coordinated to the silver cation.

Peaks at 2923, 1461, and 1375 cm -1 are associated with the nujol used to prepare the mull.

In both silver complexes, the β-diketonate ligand can be evidenced by the presence of the C=O stretching at 1658 cm -1 , while the C=C stretching vibration is associated with the absorptions at 1500-1550 cm -1 .

In addition, the bands present in the interval 1050-1400 cm -1 can be attributed to typical C-F vibrations of the -hfac ligand present in both adducts. For AgP, the bands observed in the 700-1100 cm -1 range are considered fingerprints of the phenanthroline coordination to silver ion. In AgT, the broad band observed in the 1000-1300 cm -1 range might be associated with absorptions from polyether C-O bending and/or stretching overlapped with the C-F stretching of the β-diketonate. Furthermore, bands at 700-1100 cm -1 are related with glyme modes. The C-H glyme stretching modes, lying in the 2800-3000 cm -1 range, overlap with nujol features. Assignments are reported in the Table III-4 together with the notation used for phenanthroline.

The 1 H NMR spectrum of AgT (Table III-4) can be assigned using comparative arguments with data from the related lanthanum complex. 31 * The notation has been used for phenanthroline: † The notation has been used for triglyme: precursor commercially available and two newly synthesized as previously described. All the precursors were tested in both MOCVD systems.

III.2.1 Ag coatings via PI-MOCVD.

First of all, PI-MOCVD was used since less precursor solution is needed. Our PI-MOCVD reactor has been introduced in Chapter II. All the Ag precursors listed in Table III- Also, in all silver coatings deposited with PI-MOCVD, due to the geometry of the MOCVD system, the deposition center with diameter about 4 cm are always visible. These films were used to obtain silver oxides by oxidation, as explained in section III.

III.2.2 Ag coatings by Aerosol Assisted MOCVD.

Due to that the final objective of this work was to deposit Ag and Ag-Cu oxide films, thus different precursors were also tested with our AA-MOCVD system, since it incorporates an atmospheric oxygen plasma that can be used to create highly oxidizing conditions.

As with PI-MOCVD, AgAc was tested in AA-MOCVD as precursor for Ag coating as well.

At higher deposition temperature (more than 250 °C), no Ag could be deposited on the glass substrate. At lower temperatures (210 °C) a coating was obtained, mainly composed of Ag To obtain better quality Ag coatings with AA-MOCVD, the two new silver precursors were tested as well with deposition temperatures ranging from 250 °C to 350 °C. We could obtain silver coatings with both precursors at deposition temperatures as low as 270 °C. Due to the difficulty for characterizations on glass substrate, only the results of thicker coatings deposited at 320 °C will be presented here with the deposition duration shown in Thus, although AgP is able to yield Ag coatings when used as precursor in AA-MOCVD, these present carbon species which are expected to affect the properties of the final coating (i.e. conductivity) and thus can be not suitable for certain applications. At the same time, the direct deposition of Ag@C core-shell nanoparticles could be indeed interesting especially for plasmonic applications.

III.3 Oxidation of Ag coatings

Having optimized the deposition of Ag coatings via AA-MOCVD using the two new precursors, the deposition of silver oxides was tackled, as an optimization step for the final deposition of AgCuO2. Due to the noble nature of the Ag metal, and thus the low thermal stability of Ag oxides, obtaining a high quality silver oxide film is very challenging. In this section we describe three approaches that have been evaluated for the deposition of Ag oxides. Those are, direct deposition of oxides by using plasma assisted AA-MOCVD, and oxidation of pre-deposited Ag films by electrochemical oxidation and by oxygen plasma treatment.

Direct deposition of silver oxide in AA-MOCVD system was thus tested, but due to the high decomposition temperature of the silver precursor (above 270 °C for an efficient deposition), the silver oxides are unstable at this condition. In addition to that, the oxidation gas (mainly being ozone and radicals generated by the plasma) are highly corrosive and thus deposition rate needs to be higher than etching rate. This can be achieved by increasing the precursor concentration in the solution. Given the high cost of Ag precursors, and the rather high amounts of precursor needed in AA-MOCVD for each single deposition we decided to perform post treatment of the Ag metal coatings rather than further attempting to optimize the direct deposition of Ag oxides. The experiments performed and the results obtained are detailed next.

III.3.1 Electrochemical oxidation of Ag coatings

The setup of electrochemical oxidation has been introduced in Chapter II, Experimental part.

In this section, the study of the electrochemical oxidation of metallic Ag coatings has been studied (the study of the oxidation Ag/Cu alloys will be presented in Chapter IV). As usual, This is in accordance with the reference works 54,55 . 1 M NaOH solution was used as the electrolyte, since several works reported for AgCuO2 synthesis used such high pH (>13). 56,57 As it will be shown by works introduced in Chapter V, the high pH is crucial for AgCuO2 phase formation. At forward bias potential, two oxidation peaks at +0.34 V, +0.607 V are clearly visible, corresponding to Ag→Ag 1+ and Ag 1+ →Ag 2+ , respectively. When scanning back, two reducing peaks at +0.31 V and +0.05 V appeared, corresponding to the reduction of silver following the inverse reactions with respect to the oxidation sweep. The morphology of the oxidized Ag films obtained with various durations are shown in In addition, in order to fully oxidize the thin film, a conductive substrate is preferred.

III.3.2 Oxidation of Ag film by Oxygen Plasma

Apart from using an electrochemical approach, silver can also easily be oxidized with ozone.

Silver oxidation by ozone has indeed been reported in the work of G.I.N. Waterhouse et al., who studied the reaction of ozone with a polycrystalline silver foil. 59 In addition, the interaction of silver powder with ozone have been also studied. 60 Even AgNO3 solution has been reported to react with ozone and form AgO or Ag7O8NO3 depending on the solution pH. 61 In addition, a solid state phase transition from Ag2Cu2O3 to AgCuO2 in ozone atmosphere has been reported in a work by D. Muñoz-Rojas et al. 62 In general, ozone has been proven to be a very efficient oxidizer gas to achieve silver oxides with high oxidation state. In this work, by taking advantage of the atmospheric plasma head coupled with the AA-MOCVD system, ozone can be generated thus to create highly oxidizing environment.

As stated above, while we tried to directly deposit silver oxide films with plasma, the precursors available to deliver silver coatings require deposition temperatures of at least 270 °C, which is too high for the existence of silver oxide phase. Thus, an alternative approach was taken by us, which is a two step process involving silver film deposition by PI-MOCVD and oxygen plasma treatment afterwards.

To choose a proper power and inlet gas for ozone generation, Air and O2 were tested and the plasma was analyzed by ICP-coupled Optical Emission Spectrometry (ICP-OES) (model Chapter 3. Thin film deposition of silver and silver oxide -99 -Thus, to make sure ozone is generated through atmospheric plasma, pure oxygen as gas source is preferred and a minimum power of 300 W is required. In addition, when oxidizing the Ag films, one thing to keep in mind is that radicals generated in the plasma can be very aggressive and thus corrosive for the Ag films. To minimize the erosion from radicals, the oxidation process was conducted with a humid environment (created by injecting water mist flow from a separate inlet into the chamber) and with minimum power (with 300 W).

III.4 Chapter summary

Silver coatings have been obtained through both PI-MOCVD and AA-MOCVD. During the process, 5 types of silver precursors have been tested, from which 3 precursors (including the cheap and commercially available AgAc and two new designed and synthesized precursors, AgT and AgP) allowed the deposition of Ag coatings. The silver coating obtained from AgAc is reported here for the first time and the effect of the dispersant molecules has been discussed, which gives the conclusion that with ethanolamine as dispersant more volatile silver acetate ammine complex can be formed, thus facilitating the precursor transport in PI-MOCVD vacuum condition and thus the high deposition rate. In addition, two new silver metal organic precursors have been synthesized Secondly, different approaches have been adopted to obtain silver oxide films to lay a foundation for further work of AgCuO2 phase synthesis. Direct deposition of silver oxide films coupling with the atmospheric plasma system were tested. Due to the relatively high deposition temperatures (at least 270 °C) for all the precursors available to us, silver oxide phase couldn't exist at such temperature. Thus, a two-step process to achieve silver oxide film was used, which was pre-deposition of Ag coatings by MOCVD and after-treatment with oxidation techniques (electrochemical oxidation and oxygen plasma). With both oxidation techniques, AgO phase was obtained, but there are still several issues to be solved.

For silver oxide films obtained from electrochemical oxidation, they presented a rough surface and conductive substrate is necessary for a full oxidation. Meanwhile for oxygen plasma treatment, even though it is a rapid oxidation technique and gives smoother and denser silver oxide films as compared with the previous technique, the corrosion from the radicals in plasma is difficult to control and minimize. In both cases, silver oxide films with pure phase composition are rather difficult to obtain.

III.5 Appendix

Chapter IV Deposition of Highly crystalline Cu 2 O films via AA-MOCVD

In Chapter III, the deposition of metallic silver and further oxidation to its oxides by MOCVD and post treatments have been detailed. To obtain the AgCuO2 phase by MOCVD, the parameters and setup optimization for the deposition of copper oxide thin films was also carried out. Given that oxidizing conditions are needed to obtain AgCuO2 and that plasma was only available in our AA-MOCVD system, we only studied the deposition of Cu oxides by AA-MOCVD. In this chapter of work, in order to overcome several issues, optimizations of our deposition system will be described. The effect of different parameters, such as oxygen ratio, presence of Ag and humidity, on the morphology of the Cu2O films will be Atomic Force Microscope (CAFM) and automatic crystal orientation and phase mapping technique (ASTAR), the correlation between conductivity and phases will be revealed.

Lastly, to understand the formation mechanism of larger size grains with the different deposition conditions, short duration depositions have been conducted and different thin film growth modes have been discovered. Furthermore, from the point view of thermodynamics, the formation mechanism has been explained by applying the classical nucleation theory.

IV.1 Optimization of deposition system and reaction parameters

IV.1.1 Deposition system configuration

The original AA-MOCVD system has been presented in Chapter II ( Comparing to the Cu2O films deposited with the same conditions using the later optimized case (figure b, see details below), the color gradient due to the uneven thickness is rather clear in figure a). Even worse, with the same precursor consumption rate in the original configuration, the Cu2O thin film deposition rate is much lower than in the optimized case, indicating a higher percentage of precursor waste. In general, under such configuration in Despite the improvement observed with the optimized system configuration, i.e. a homogeneous precursor flow was obtained with majority of the bigger droplets avoided, from time to time, several holes of small size (about 0.5 mm in diameter) were still visible.

Even after a careful cleaning of the glass substrate and the deposition chamber, the problem wasn't solved. Thus in order to collect the condensed liquid and further reduce big size droplets, minimizing the chances of having bigger droplets reaching the substrate, a simple

IV.1.2 Choice of precursors

Besides the MOCVD system optimization, the choice of precursor used in MOCVD also plays an important role in the final quality and properties of the deposited films For the deposition of copper oxides, there are several commercial metalorganic precursors available.

By taking into account both precursor cost and deposition quality, mainly Cu(AcAc)2

(Copper (II) acetylacetonate, STREM Chemicals 98%) and CuF (Copper (II) trifloruoacetylacetonate hydrate, STREM Chemicals 99%) were used in this work.

IV.1.2.1 Effects of oxygen ratio during deposition with Cu(AcAc)2 precursor

Cu(AcAc)2 was firstly tested for Cu2O thin film deposition with various oxygen/nitrogen ratios and deposition temperatures, as shown in Table IV-1. Precursor solutions with a 0.02 M concentration were prepared by mixing 2.62 g of precursor with 500 ml butanol and 1.5 ml of ethylenediamine. With magnetic bar stirring for 10 more min, a crystal clear blue solution was obtained. (More detailed information can be found in Chapter II). The samples were deposited on glass substrate at temperatures ranging from 350 to 375 °C for 1 h with a total carrier gas (CG) flow of 10 L/min (composed of nitrogen and oxygen). As shown in Figure IV-6 a), Cu2O samples deposited at 350 °C present several bands over 1200 cm -1 that correspond to carbon group contaminations in the films. With higher oxygen ratio used during the deposition process, the signal of carbon contamination decreased, which can be attributed to the higher oxygen ratio facilitating precursor decomposition. But even at the maximum oxygen ratio 80% used, the carbon group signal is still rather visible. At higher deposition temperature 375 °C, similar carbon contamination and trend with oxygen ratio can be found in Figure IV-6 b). In any case, even with the presence of carbon contamination, no peaks corresponding to CuO phase were observed.

IV.1.2.2 Deposition under oxidizing conditions using plasma with Cu(AcAc)2 precursor

Due to the high oxidation environment requirement to form AgCuO2 phase, we explored the use of the atmospheric plasma system coupled to the AA-MOCVD deposition system (see in Chapter II). In addition to providing highly oxidizing conditions, the plasma system was expect to facilitate precursor decomposition, to enhance film crystallinity and to reduce carbon contamination. 1,2,3 The emission characteristics of the plasma head have been reported in Chapter III, showing that a minimum power of 300 W is necessary to obtain ozone species when using pure O2 as gas source. Before using the plasma for copper oxide film deposition, two types of glass substrates (Frost microscope slide and Corning glass)

were tested under oxygen plasma with power of 500 W for 30 min. The glasses were then characterized by Raman spectroscopy. of copper oxide with plasma was tested on this type of glass substrates, no copper oxide films could be coated on top. Conversely, for corning glass, the Raman curves present no difference before and after the oxygen plasma treatment under the same conditions. In addition, copper oxide films could be obtained on corning glass when using the plasma (as detailed below). Thus, corning glass showed better stability than normal microscope glass slide under oxygen plasma and was used as the substrate for the optimization of the deposition of copper oxide films with plasma. Due to the highly reactive radicals emitted by the oxygen plasma, precursor molecules can react/decompose before reaching the substrate which leads to less reaction taking place on the substrate. As a consequence, higher precursor concentration is required (minimum 0.05 M). The copper oxide films were deposited under different oxygen ratio plasma (30% and 60%) using a plasma power of 350 W, and a deposition temperature of 300 °C, for 1 hour.

Since the purpose of using plasma during deposition was mostly focusing on the investigation of precursor decomposition at low temperature (so that the conditions remain compatible with the deposition of Ag oxides), thus tests at higher temperatures were not conducted. The films obtained were characterized by Raman spectroscopy. As shown in Figure IV-8, in both deposition conditions, and despite the use of plasma, the deposited copper oxide films exhibited strong signal of carbon contamination as previously observed, even though the signal slightly decreased with higher oxygen ratio. As shown by the optical images in the figure, the coatings showed a dark brown color. No peaks from CuO or Cu2O phase could be observed in the Raman spectra, possibly due to a low crystallinity at this low deposition temperature. Therefore, the deposition of carbon free Cu2O films using Cu(AcAc)2 as precursors were rather difficult in our AA-MOCVD system, better copper precursor with clean decomposition was needed.

IV.1.3.2 Optimization for oxygen ratio

It is well known that oxygen plays an important role during the MOCVD process, especially for metal oxide thin film deposition. [4][5][6] In this work, experiments were carried out with various oxygen ratios during deposition, as shown in Table IV

Sample names D1 D2 D3

Oxygen ratio (%) 12. intensities of the peaks were rather similar in all cases, with no extra peak at higher oxygen partial pressure, thus eliminating the presence of significant amounts of CuO phase in the films. As well, the intensity of the peak around 200 cm -1 presented a rather big difference between the deposited Cu2O films and the reference curve, which was possibly due to certain were also conducted, due to the low thickness of those films, the reflections were full of noise and difficult to distinguish useful information, they are not presented here. 4 To summarize, oxygen ratio in the carrier gas presented a huge impact on the surface morphology of Cu2O films deposited by MOCVD. A slight increase of oxygen ratio during the MOCVD process yielded Cu2O films with clear cubic shape grains, while at the same time the surface became rougher and discontinuous. Comparing the deposition results and optimization process with both precursors, it's much easier to achieve good quality, carbon free and homogeneous Cu2O films with CuF as the precursor. The later described Cu2O films were thus all deposited with CuF precursor.

IV.2 Tuning the texture of Cu2O thin films

In addition to obtaining thin film depositions with good homogeneity, we managed to tune the texture of the Cu2O films by modifying and optimizing the deposition conditions.

Textured Cu2O films have been mostly obtained by electrochemical deposition. 7,8 In MOCVD, it is well known that adding water vapor to the reaction can facilitate precursor decomposition, enhancing crystallinity and reducing carbon contamination. 9,10 J-Y Kim et al 11 reported that Cu metallic films were deposited through MOCVD using a humid flow. In their work, different deposition rates were achieved by controlling the humidity. In another work by A. Jain et al. 10 , a detailed explanation was given on the mechanism by which the presence of H2O facilitates copper precursor decomposition, involving a proton transfer between H2O and the organic precursor. But the effects of humidity on the deposition of Cu2O thin films in terms of morphology, grain orientation and electronic properties has not been studied in detail.

In this section of work, several approaches have been found to tune the grain structure in Cu2O films, namely, presence of Ag nanoparticles and addition of humidity in the carrier gas. Of the two approaches, the addition of humidity in the carrier gas has shown an obvious impact on grain orientation and morphology, and thus on the resulting electronic properties.

By adjusting the humidity in the gas flow we were thus able to tune the deposition speed and texture orientation of deposited Cu2O films. As well, an evolution of mobility with deposition temperature and carrier gas humidity was observed, reaching a maximum of 17 cm 2 /V.s. In addition, we used optimized Cu2O films to form p-n junctions with ZnO,

showing an excellent rectification with an on-off ratio exceeding 10 4 . In order to understand the impact of carrier gas humidity on morphology, the initial nucleation process was studied.

Finally, we used the classical nucleation theory and thermodynamics of film growth to explain the results obtained. All the above are shown and discussed in detail in the following sections.

Chapter 4. Deposition of highly crytsalline Cu2O films via AA-MOCVD -124 -

IV.2.1 Effect of the presence of Ag particles on film growth and morphology

The effects of noble metal on the formation of Cu oxides have been discussed in the work by L.Wang et al. 12 , in which the Cu2O cubic grain formation by oxidizing a Cu0.5-Au0.5 alloy substrate was studied. Also, in another work by C.C.Tseng et al. 13 , the optical and electrical properties of Cu2O films incorporating Ag deposited by magnetron sputtering was evaluated.

Therefore, in this work we have explored the effects of the presence of Ag in Cu2O film formation by AA-MOCVD.

To achieve the incorporation of silver, AgT was added to the CuF solution. Due to the costly nature of silver precursors, the exploration with higher precursor concentration was avoided. Copper precursor concentration was fixed at 10 mM and depositions were carried out at 335 °C for 3.5 hours. The quantity of Ag in the final Cu2O films was adjusted by varying the initial silver precursor concentration. The quantity of Ag in the film was estimated by EDS (using potential of 15 kV with a full scan of surface area about 3 μm × 3 μm). In all these samples, the small quantity of metallic Ag in the film could not be detected by XRD. Electronic properties, such as carrier concentration or mobility, of those films were not investigated due to that the sheet resistance of the Ag incorporated Cu2O films were so resistive that they exceeded the measurable range for the Hall effect equipment we have (sheet resistance maximum up to 10 M Ω/□). Thus only resistivity was able to be measured by four probe measurement, giving lowest resistivity value from pure Cu2O film, about 200 Ω.cm. Surprisingly, with more silver incorporated in the Cu2O film, the resistivity tended to increase, the reason behind is not clear yet.

IV.2.2 Effects of humidity on thin film orientation, morphology and electronic properties

In addition to the study of the effect of Ag incorporation on the growth and morphology of Cu2O films, the effect of carrier gas humidity was also systematically studied, including the To observe the effects of temperature on Cu2O deposition with both dry and wet carrier gas, parallel experiments were conducted with temperature ranging from 290 °C to 365 °C, as presented in Table IV-4. A constant carrier gas flow of 8 L/min was passed through the bubbler. To estimate the water molar ratio introduced in the total gas flow, we assumed that the water vapor originally present in the carrier gas was 0. The water consumption was 0.11 ml/min, which corresponds to a molar ratio of 1.37 %. The morphology of the Cu2O films are presented in observed above 1000 cm -1 . In addition to that, the strongest peak for sample D335 located at around 170 cm -1 , while for W335 the strongest peak position stayed at 210 cm -1 , again in agreement with the different texture of each sample.

To further study the effects of carrier gas humidity, besides varying the deposition temperature, the humidity was also tuned while fixing the deposition temperature at 335 °C.

The morphologies of Cu2O evolved as shown in the SEM images in intensity shift happened between the strongest and second strongest ones, which could be associated with the grain orientation change. Lastly, we could also observe that the peak at 300 cm -1 and 650 cm -1 decreased slightly, the later one even splitting into two visible peaks, both associated to Cu2O phase.

In decomposition of the precursor before reaching on the substrate. 15,16 In all cases, the deposition rate of Cu2O films was always faster with wet CG than in dry deposition condition.

Thus, water seemed to effectively activate the precursor reaction at the substrate but not its increase in mobility with increasing deposition temperature. With dry CG, the mobility of Cu2O films was enhanced from 2 cm 2 /V.s up to 13.7 cm 2 /V.s when deposition temperature increased from 305 °C to 365 °C. Similar trend was also observed in the case of humid CG deposition condition, carrier mobility being improved from 5 cm 2 /V.s to 17 cm 2 /V.s. As previously introduced in Chapter I, for pure Cu2O films deposited with MOCVD, the mobility was generally about 1 ~ 5 cm 2 /V.s, thus this value was rather good comparing to previous works using similar fabrication techniques [17][18][19] . This enhancement of mobility could be attribute to the bigger grain size obtained with humid CG deposition condition (maximum 300 nm measuring along the cubic edge). However, the sample with lowest resistivity was not the one deposited at highest temperature, which is attributed to the carrier concentration decrease over 350 °C, as shown in figure c).

In addition, for Cu2O films deposited with various water molar ratio at 335 °C, deposition rate and electronic properties were also studied. As shown in Figure IV-22 a), along with higher water molar ratio from 0 % to 1.12 %, an increase in deposition rate from 2.2 nm/min to 4.6 nm/min was observed. From there, growth rate became almost constant. The electronic properties of those films were again evaluated by Hall-effect measurement with Van de Paul four probe configuration. For water molar ratio from 0 % to 1.12 %, the resistivity values increased from 180 Ω.cm to 900 Ω.cm. A further increase in water molar ratio in carrier gas produced a resistivity drop back to 160 Ω.cm with water consumption of 0.11 ml/min. Conversely, carrier concentration showed a decreasing trend from 7.7 × 10 15 cm -3 for 0 % to 1.1 × 10 15 cm -3 for 1.37 % water molar ratio. Since the origin of the p type conductivity of Cu2O film relies on the presence of copper vacancies, carrier concentration in the films is very sensitive to the oxygen ratio change during MOCVD. 20,21 The oxygen introduced in the flow was fixed while the concentration of water vapor varied. Thus, the decrease of carrier concentration with increasing humidity could be explained due to the fact that in presence of water, precursor decomposition was largely accelerated. Thus, this increased supply of copper atoms created an oxygen deficient environment during the deposition process. Since the number of oxygen vacancies in Cu2O largely rely on the oxygen/copper ratio during the deposition, [21][22][23][24] thus the oxygen deficient environment led to a decrease of carrier concentration. However, with a further increase of the water ratio from 1.12 % to 1.37 %, a slight increase of carrier concentration from 1.1 × 10 15 cm -3 to about 3 × 10 15 cm - 3 was observed. Meanwhile, by reproducing two more samples at the same conditions we obtained lower concentration around 1.1 × 10 15 cm -3 (with mobility around 13 cm 2 /V.s).

This would mean that other factors may be affecting the carrier concentration of the films and that a more detailed study would be necessary to fine tune this parameter. On the other hand, the effect of water facilitated on precursor decomposition was rather saturated over 1.12 % water molar ratio, as shown in Figure IV-22 a). Meanwhile, with higher water molar ratio during MOCVD deposition, the carrier mobility of the produced Cu2O thin films was largely enhanced from 4 cm 2 /V.s to 17 cm 2 /V.s, which could be attributed to the bigger grains observed in Figure IV-17. In addition to the area conductivity mapping, the electrical behavior on certain points was also conducted. To improve the reliability of the results, both Cu2O films deposited at 305 °C (W305-08) and 335 °C (W335-08) were characterized. Full conductivity scans were firstly conducted over a surface area of 0.5×0. To further understand the mechanism behind the high conductivity observed in the grain boundaries, the crossection of the (111) oriented Cu2O film (on Si substrate, for the sake of easier sample preparation) was characterized with the ASTAR technique, a TEM-EBSD like phase mapping characterization tool with ultra resolution as introduced in Chapter II.3.2.3.

ASTAR (TEM-EBSD like technique) Automatic crystal orientation/phase mapping of Cu2O sample (W335-08) crossection with orientation (111). a). Reconstructed correlation grey scale index map. b). Phase map (with index-reliability). The red color corresponds to Cu2O phase, the blue color represents

CuO phase and the green color refers to the Si substrate.

Taking into account that CuO has a better conductivity than Cu2O and the position where the majority of the CuO phase located, it's rather straight forward to associate the existence of CuO in grain boundary with the better conductivity that has been observed. The possible effects from the shape of the AFM tip on contact area has been taken into consideration, but the radius of the AFM tip was merely 25 nm. This, together with the rather high resistivity of the Cu2O films, allow to consider the effects from the tip shape almost nnegligible. Band (VB). According to energy diagrams 26,27 , when the AFM tip was pointing above the Cu2O grains, a typical rectifying junction was formed between Pt/Cu2O/ITO, as shown in figure a). Thus, it allowed the easy pass of current along one direction while no current was presented with reverse bias. This would be the explanation of the weak rectifying curves shown in Figure IV-24 e) and f) on Cu2O grains.

Meanwhile, due to the presence of CuO between the grain boundaries, figure b), when the AFM moved above the grain boundary, current was running through both Cu2O and CuO .

Thus, no rectifying behavior was observed on the grain boundary as shown in Figure IV-24 e) and f). However due to the presence of CuO phase (lower resistivity than Cu2O), a higher conductivity was observed on grain boundaries.

Optimized Cu2O films presenting low resistivity and high carrier mobility were combined with ZnO to form p-n junctions. n type ZnO ~70 nm thick layers were deposited on top of the Cu2O films at 60 °C using a home made Spatial Atomic Layer Deposition (SALD) system (AP-SALD) system, 20,28,29 in collaboration with Mr. V. Nguyen from LMGP. In this approach, precursors were continuously injected in different locations of the reactor while the substrate oscillated at very short distance (50 -200 µm) from the head. Inert gas barriers located alternatively between precursor flows acted as effective barriers to keep precursor gases chemically isolated, thus conventional ALD cycles were replicated but with up to two orders of magnitude faster. Diethyl zinc ((C2H5)2Zn; DEZ) and water vapor (H2O) were used as precursors for zinc and oxygen, respectively.. More details can be found in the publication by N. Nguyen et al. 20 Chapter 4. Deposition of highly crytsalline Cu2O films via AA-MOCVD -142 -Diodes were made using the Cu2O films deposited with humidity condition at 335 °C (sample W335-08, see Table IV- Between the ZnO and Cu2O layer, no cracks were observed, thus it indicated a good adhesion.

The rectified behavior of the p-n junction is presented in Figure IV-27 c). In the I-V curve, the voltage bias was scanned from -1 V to +1 V and it showed an excellent rectifying shape.

The leakage current was merely 2.2 × 10 -6 mA at -1 V, with on current reaching almost 6 × 10 -2 mA at +1 V, thus giving an on-off ratio exceeding 10 4 . This indicates that with SALD, even at low deposition temperature (60 °C), a good p-n junction can be formed between ZnO and our optimized Cu2O layers.

IV.3 Studies on the reaction mechanism

IV.3.1 Study of the initial stages of thin film growth.

In order to explore the effect of humidity in the reaction mechanism and texture development, layers presented more nucleation sites, smaller grain size and smoother surfaces. After 50 min, the films deposited with dry CG didn't show a clear increase in grain size and instead more small grains were present. Conversely, Cu2O films deposited with wet CG presented a clear increase in grain size with longer deposition duration, and there were much less smaller grains that for the films deposited with dry CG.

In MOCVD there are three well-known growth modes at the initial stages, 30,31 namely , layer-by-layer growth, island growth and layer plus island growth. In the majority cases of thin film depositions by MOCVD, the last mixed growth mode is the most commonly seen, which is also called the Stranski-Krastanov (SK) growth. In this case, AFM results have shown that for dry CG deposition condition, a layer-by-layer growth mode took place, in which each layer was made of many small particles that presented a moderate growth in size upon longer deposition duration. Additionally, to further understand if such a texture orientation change is mainly due to the initial crystallization difference, a sample was deposited using humid CG during the initial This implies that with 10 min of deposition under humid CG condition, it is enough to induce a different texture. Thus, these results suggest that once the first few nanometers of films were deposited, by following the principle of surface energy minimization, the later growth was compelled to be stacked in a similar manner by keeping the same orientation. It would also implies that growth along the crystallographic faces in humid conditions is more favorable that nucleation of new grains, even when water is not present anymore. Thus, the addition of humidity in the CG has key effects on the initial stage of reactions and trough the water assisted precursor hydrolysis and reaction, favoring growth along the (111) direction, which in turns favored the formation or large crystals.

IV.3.2 Thermodynamic explaination

This section aims to explain the grain growth mechanism from the point view of a nonquantitative thermodynamic approach. During MOCVD there are several process taking place. To simplify the problem the assumption that processes occurring before precursor decomposition are relatively similar in all the cases has been made. This includes precursor transport, surface diffusion and precursor adsorbing during MOCVD deposition of Cu2O.

Therefore only precursor decomposition and Cu2O nucleation and growth are discussed in our thermodynamic model. In order to apply the classic thermodynamic nucleation theory, the process from CuF to Cu2O has been divided into two stages in our case. Firstly CuF is decomposed (thermally) and reduced into Cu metal atom (by its own reducing organic ligand). Then the metal Cu atoms are oxidized (by reaction with oxidizing reactants during deposition). Therefore, in this case, the thin film deposition rate mainly relies on the precursor decomposition rate on substrate (the first stage CuF decomposition), meanwhile the grain size of the obtained thin films depends on the nucleation and grain growth process (the second stage Cu atoms oxidization).

Even though the process does most likely not follow this separate and distinct steps, we can use this approach in our model since the Gibbs free energy change between CuF and copper oxide does not depend on the process, but only on the beginning and final states.

In the Cu2O thin film deposition process, we can thus define ∆G as the total Gibbs free energy from CuF precursor to form one mole of Cu2O in the MOCVD process. Thus, it equals to the Gibbs free energy of precursor decomposition into copper atom (∆G R ) plus copper atom oxidization with oxygen (∆G O ) to form Cu2O, as shown in Equation IV 22.

∆G = 2∆G 𝑅 + ∆G O (22)
Since ∆G R is related to the reactivity of the precursor, thus it's the dominant factor for deposition rate during MOCVD process. While ∆G O is associated with copper atom oxidation, thus it's the major impact on Cu2O grain nucleation and enlargement, which can be applied into classic nucleation theory as shown in Equation IV 23, 32

N s ~ exp (-∆G O /kT) ( 23
)
where Ns is the nucleation rate and k is the Plank constant and T is the temperature (in Kelvin).

To begin with, several terms and definitions need to be clarified. Firstly, in this work there is two terms related to oxygen ratio, PO2 and PO2 * , which corresponds to oxygen ratio in deposition chamber and oxygen ratio regarding to copper atoms separately. Thus, it's defined that PO2 * = PO2/MCu, where MCu is the quantity of copper atoms decomposed from the CuF precursor per unit time. In the majority of cases, when the supplying of copper atoms for oxidization is constant, those two terms are equal and both constant when they are applied in thermodynamic equations, such as in the work of L.Wang et al. 12 But in this work, since in MOCVD process the Cu atom supply can be largely altered and dependent on the precursor reactivity, thus those two terms need to be differentiated for the sake of clarity, therefore the oxygen ratio in Equation IV 25 should be PO2*. For ∆G R , if we don't consider the decomposition of the ligand here, based on the reaction CuF→ Cu + Ligand (volatile), thus Equation IV 24 is given. [CuF] is the precursor decomposition rate at standard condition.

[Ligand] is thus the reactivity of the ligand as a constant here. Also, since the reaction to obtain one mole of Cu2O is 2Cu + 1/2 O2 → Cu2O, thus Equation IV 25 is obtained. This is mainly a description of the crystallization and morphology process, in which ln [CuF] is the reactivity of copper metal organic precursor and it's directly related to copper atom supply, thus growth speed. ln [Cu] is the reactivity of copper atoms.

∆G 𝑅 = ∆G 0 𝑅 -RTln[CuF] + RTln [Ligand] (24)
(Governs growth speed)

∆G O = ∆G 0 𝑂 -2RTln[Cu] -1/2 RTln[ 𝑃 𝑂2 𝑀 𝐶𝑢 ] (25)
(Governs grain size)

Using Equation IV 3 and 4, the deposited Cu2O films with water vapor and Ag incorporation are explained as following.

Effects from H2O vapor: It is known that during a MOCVD deposition the presence of water vapor can facilitate the decomposition of metal organic precursor, especially in the case of metal oxide deposition. [33][34][35] Therefore, in our case humid gas led to a higher precursor reactivity, the term [CuF] increased, thus ∆G R was more negative indicating a faster deposition rate. Meanwhile, more copper atoms were supplied during oxidization and nucleation process, while the oxygen partial pressure in chamber PO2 was maintained, thus the oxygen/copper atom ratio term PO2/MCu became smaller and thus ∆G O was less negative.

Based on the nucleation theory, with less negative Gibbs free energy, the nucleation is less favored than grain growth, as confirmed by the results shown in Copper atom reactivity: With higher Ag precursor concentration during deposition, more Ag atoms were incorporated into the Cu2O film. As a result, bigger grains of Cu2O appeared also with a pyramid-like surface, as shown in Figure IV-11. In the reference work 12 , the effects of a noble metal on reducing the reactivity of copper atoms during oxidation process is presented. Therefore, as shown in Equation IV 4, the total Gibbs free energy was less negative. Thus, the nucleation rated was slowed down due to the lower Gibbs free energy, and growing big grains was favored. Eventually, with the Ag incorporation, thin films with larger grains were formed. A catalytic effect of Ag in the decomposition of the precursor is also possible but has not been considered in this discussion.

IV.4 Chapter IV conclusions

In this chapter, first of all the configuration of AA-MOCVD system in LMGP was optimized.

Based on that, the deposition parameters, including precursor choice, substrate and the use of plasma, were also adjusted to deliver homogeneous and carbon contamination free Cu2O thin films. In the process, CuF was chosen as the precursor instead of Cu(AcAc)2 due to a cleaner decomposition. With the optimized MOCVD system and Cu2O deposition condition, the studies on the effects of oxygen ratio, noble metal incorporation (Ag nanoparticles in this case) and carrier gas humidity were conducted. The results revealed that, during the deposition of Cu2O films in AA-MOCVD, a slight higher oxygen ratio, silver nanoparticles and higher humidity all could have huge impacts on the morphology and lead to an increase to grain size. As well, the cubic shape Cu2O grains emerged more clearly on the surface.

Later on, the effects of humid CG on initial thin film growth mode were explored, identifying the layer-by-layer growth mode for thin films deposited in dry CG while island like growth mode for depositions in wet CG. To better understand the results from those three factors, a simplified thermodynamic model was established and discussed. Based on that, the reason for the emergence of big grains at each condition was explained.

Among those three factors, the effects of carrier gas humidity on morphology, grain orientation and electronic properties were investigated in detail. With the wet CG during deposition, the deposition rate was largely enhanced and the Cu2O films with "pyramid-like" surface were obtained, with large grains of 300 nm on the cubic edge. Those types of thin films have been confirmed by EBSD, XRD and TEM to be well textured along the (111) index. In addition, through observing the cross-section of (111) oriented Cu2O film by TEM, it was discovered that those films were made of stacks of columnar grains. Furthermore, the electronic properties of deposited Cu2O films have been studied, in which higher mobility was found on average on (111) textured samples and maximum mobility reached 17 cm 2 /V.s.

In addition, the investigation of the conductivity of the Cu2O films was also carried out. With the assistance of electrical AFM, the current mapping of the Cu2O surfaces was obtained, indicating a higher current on the grain boundaries.

Chapter V Deposition and characterization of AgCuO 2 thin films

As presented in Chapter I, AgCuO2 has shown to have potential to be an excellent p type small bandgap (about 1 eV) semiconductor material. The first discovery of this phase was through solution reaction and resulting in powder form 1 . Although several works have reported its crystal structure 2 , single crystal resistivity measurement 3 and its electronic structure 4 , the optical properties (eg. transmittance, bandgap) and further electrical characterizations (e.g. Hall Effect measurement) have not yet been conducted. The lack of related reports is due to the difficulty of AgCuO2 thin films synthesis. To achieve such phase, the synthesis condition has to consider both the thermal instability inherent of Ag compounds and the high oxidation states in this metal oxide. Only recently a few studies have reported attempts to synthesizing AgCuO2 in thin film through electrochemical deposition, pulsed laser deposition or sputtering. [5][6][7] The electrochemical deposition method yielded rather rough, discontinuous films containing micron sized nanoplates of AgCuO2, on ITO substrate. 5 With the reported work with the PLD technique, the deposited thin film was not well crystallized due to the relatively lower deposition temperature (room temperature). 8 Sputtering films were usually not pure and presented very low crystallinity. V.1 Co-deposition of silver copper oxide via Plasma Enhanced AA-

MOCVD

As previously mentioned, the initial attempt was trying to use both silver and copper precursors in solution to achieve AgCuO2 phase by AA-MOCVD at relatively low deposition temperatures (less than 270 °C, to avoid decomposition of Ag-Cu mixed phases).

With the foundation of the previous work, by comparing the lowest decomposition temperature of the precursors available, Ag(hfac)triglyme (AgT) and Cu trifluoro acetylacetonate (CuF) were chosen as the precursors. The lowest possible deposition temperature from both precursors is at least 270 °C. Even though the temperature is still too high concerning the phase stability, there is still possibility to achieve AgCuO2 thin films at such temperature in an oxidizing environment generated by oxygen plasma (containing ozone and radical species).

Due to the limited availability of AgT precursor (also due to the high price of Ag), the maximum mixed precursor solution concentration used was 0.02 M, which was prepared by mixing 4.9 g of AgT and 3.69 g of CuF in 500 ml of ethanol and stirring for 10 minutes at room temperature. A crystal clear green-blue solution was achieved. A plasma power of 350 W was used with mixture of oxygen and argon as gas source. Corning glass was used as substrate for the deposition. Several depositions were performed at 250 °C and 270 °C for 1 hour, as shown in Table V-1. The obtained coatings were characterized by SEM and EDS analyses. Although XRD characterization was also used for phase identification, due to the low thickness and crystallinity of the films obtained at such deposition temperature, no reflections were obtained in the XRD patterns.

Samples were deposited with and without oxygen plasma at the same conditions. Table V-1 details the deposition parameters and the images of the deposited thin films. Judging from the intensity of the color, when using plasma during depositions the resulting samples were lighter, indicating thinner coatings (sample C1). Even though the oxygen ratio in plasma gas source was decreased, attempting to reduce the corrosion ability of the radicals, the deposited films were still rather thin (sample C2 and C3). Despite SEM characterization was attempted, the strong charging effect did not allow to obtain topographical information on the sample. By slightly tuning down the plasma power to 300 W, at the same deposition temperature for 1 hour coating, the sample obtained (C4) showed a much more intense dark color, thus indicating a probably thicker deposition. Conversely, with further increasing the plasma power with those deposition conditions, coatings were barely visible on the glass substrate.

In addition, for the sake of comparison, two samples were also coated at 250 °C and 270 °C without plasma, as shown in Table V-1 (samples C5, C6). Without oxygen plasma, even at 250 °C, a yellow colored deposition was observed. At 270 °C, the 1 hour duration deposition without plasma resulted in a much darker colored thin film. Thus, once again the corrosive effect from the radicals in high power oxygen plasma was rather obviously evidenced. Further tests to co-deposit silver copper oxide thin films were not conducted, due to the following several problems. First of all, to meet the temperature gap between the AgCuO2 phase stability window and precursor decomposition in MOCVD, precursors with much lower decomposition temperature (at least lower than 200 °C) would be required. Due to the fact that all commercialized precursors with such low decomposition temperature, especially for silver precursors, are extremely expensive (more than 100 euro/g). Meanwhile, in AA-MOCVD, large quantity of precursors are required for each deposition (at least 1 g of precursor per deposition). Furthermore, the corrosive radicals from the oxygen plasma even bring more difficulty for the deposition process. Thus, higher precursor consumption per deposition is needed, making such co-deposition approach economically unaffordable for us.

Therefore, even though it's theoretically achievable, co-deposition of AgCuO2 is not easy or affordable through Plasma Enhanced AA-MOCVD method with the current experimental conditions. For future work, an ozone generator is suggested to replace the role of oxygen plasma, so that only ozone species are produced without corrosive radicals. In addition, designing/finding new precursors with even lower decomposition temperature is necessary and the deposition temperature is suggested not exceeding 200 °C. Thus, other more convenient approaches to obtain AgCuO2 thin films were adopted in this work.

further increase of copper precursor concentration to 0.034 M, the Ag:Cu ratio in the alloy film was rather close to 1:1 ( actual value 1:0.91), as shown in figure c).

V.2.2 Electrochemical oxidation

In Chapter III, it has been shown that silver (I, II) oxide films can be achieved through electrochemical oxidation of conductive Ag films. Therefore, a similar approach was applied for the oxidation of AgCu alloys. As usual, the initial step was to perform a cyclic However, the cyclic voltammetry presented in Figure V-5 does not show any clear oxidation peak from 0 V to 1 V. From 1 V to -1 V, two reduction peaks were observed at -0.82 V and -0.95 V. From -1 V to 0 V, two oxidation peaks were observed at -0.43 V and -0.16 V.

Therefore, only two reversible processes seemed to appear after a first oxidation of the sample that are hard to assign. The morphologies of the electrochemically oxidized AgCu samples in 0. Finally, with even higher electrolyte concentration (NaOH 1 M), oxidation of AgCu films were conducted by applying 0.5 V and 0.7 V for 17 min. The morphology of the oxidized samples is presented in Figure V-10. At lower oxidation potential 0.5 V, clusters made with nanoplates like structures emerged out of the surface. Similar clusters can also be found on the surface of the sample oxidized at 0.7 V. Thus, even if the desired AgCuO2 phase is able to be formed through such oxidizing technique, the morphology of the obtained films will be rather rough and not homogeneous, which is not ideal neither for thin film property characterizations nor for device fabrication. So as shown, the oxidation of AgCu films through electrochemical approach was tested in different concentration of NaOH solution electrolyte with potential ranging from 0.5 to 0.7

V. The treated samples were characterized by GIXRD and no reflections were found corresponding to the desired silver copper mixed oxide phase. In addition, the oxidized samples presented quite rough morphology, which is in accordance with the results of the electrochemical oxidation of Ag films presented in Chapter III. Conversely, as previously reported, the oxygen plasma treatment was able to oxidize 100 nm Ag films rapidly, the morphology of the oxidized films presented a smooth and continuous manner. Therefore, the same oxidation technique, oxygen plasma treatment, was adopted and the results are presented in the next section.

V.2.3 Oxygen plasma treatment

In addition to the electrochemically oxidation, there are several other approaches that can be applied to oxidize the AgCu film. In this work, oxygen plasma treatment was also tested. As previously introduced with oxygen plasma treatment of Ag films, such oxidation technique is rapid and it doesn't have the requirement for a conductive substrate.

The oxygen plasma treatments were conducted with different plasma power (ozone intensity) and substrate temperatures. The morphology of the as-deposited and oxidized films is was conducted in ozone environment at room temperature and lasting for more than 10 hours. In addition, the ozone treatment of a mixture of metallic copper and silver power with 1:1 ratio was reported. The ozone treated powders only resulted in small amount of AgCuO2 phase. The major obtained phase was silver oxide but without copper oxide. In addition, the ozone oxidation of AgCu alloy films were also tested, and the results showed that it was mostly silver that was oxidized, in agreement with the results here. In our work, because of the thin film form of the alloy instead of the powers, less areas were exposed to the ozone species, which further increases the difficulty to achieve the full oxidation of the alloy film.

In addition, in this work, due to the corrosive radicals from the oxygen plasma, long duration exposure of samples to ozone species were not possible. Combining all those reasons, the oxygen plasma treatment of AgCu alloy films was not suitable to obtain AgCuO2 thin films neither.

V.3 Deposition of AgCuO2 thin films by Successive Ionic Layer Adsorption and Reaction (SILAR)

As previously introduced, several approaches have been tested to synthesize AgCuO2 thin films. Initially, the direct co-depositions were attempted via Plasma coupled AA-CVD. Then, the oxidation of the AgCu alloy films was also tested. Unfortunately, none of those methods was able to yield AgCuO2 thin films. As previously introduced, to obtain AgCuO2 phase, there are mainly two criteria, low deposition temperature and highly oxidizing environment.

Those two requirements can be easily satisfied in solution-based deposition techniques.

Therefore, to obtain AgCuO2 thin films, the possibility of using solution based coating techniques was investigated.

Successive Ionic Layer Adsorption and Reaction (SILAR) has been widely used for the deposition of high quality binary/ternary compound thin film coatings. [20][21][22] The reaction mechanism of SILAR, which relies on the adsorption and reaction of metal cations and negative anions, has been introduced in Chapter II. Since SILAR is a low temperature deposition technique, thus the concerns about AgCuO2 thermal stability can be eliminated.

The high oxidizing reaction environment in SILAR process can be easily created by adding oxidizing chemical reagents, such as H2O2, Na2S2O8, etc. Thus, this part of work will introduce the experimental setup and parameters used for the deposition of AgCuO2 thin films by SILAR method. In addition, the electronic and optical properties of the obtained AgCuO2 thin films will be reported for the first time in literature.

V.3.1 Experimental setup and parameters

The metal oxide thin film coating process known as SILAR is shown in Figure V-14. A specifically chosen substrate is immersed in metal and cation ion solutions in sequence, and buffer solution in between to rinse away the excessive ions on the surface. Different concentrations of AgNO3 (99%, Sigma Aldrich), and Cu(NO3)2.3H2O (98%, Sigma Aldrich)

were used as the metal precursors in the metal ion solution to meet the stoichiometry in AgCuO2 in the final films. Pure deionized water was used as buffer solution. To deposit metal oxide, NaOH was introduced as the anion solution. According to previous works about solution synthesis of AgCuO2 powders, the concentration of NaOH solution was chosen as 1 M. 1,2 Due to the high oxidation states in AgCuO2, sodium persulfate (Na2S2O8, 98%, Sigma Aldrich), was also mixed into the anion solution as a strong oxidizer. 3,23 During the deposition process, the immersion speed was set to 60 mm/s and withdrawal speed was fixed at 5 mm/s. The moving distance on Z axis and rotation angles of the different motors in the X-Y plane were adjusted and controlled through a central server connecting with a PC using Perfect Motion V.02 software.

V.3.2 Substrate preparation

The substrate is crucial for a successive SILAR deposition, thus different substrates were tested, including Corning glass, Si substrate and metal oxide thin films (Cu2O, ZnO and TiO2). Among these substrates, Cu2O/glass, ZnO/glass and TiO2/glass were prepared in LMGP. Below, the substrate preparation for SILAR deposition is described for each case.

D. Cu 2 O/glass substrate preparation:

The Cu2O films were coated on cleaned Corning glasses (with size about 2.5 cm × 3.75 cm).

The depositions were conducted in AA-CVD as introduced in Chapter IV. During the deposition, aerosol were formed with 0.01 M CuF precursor solution and they were carried into deposition chamber by CG1 (2 L/min) and CG2 (3 L/min) plus 5 L/min air. The ventilation value was set to 250, corresponding to 3 cm of liquid length in the pressure gauge.

The depositions were carried out at 335 °C with dry carrier gas for 40 min, reaching a thickness of approximately 70 nm.

E. ZnO/glass substrate preparation:

The ZnO layers were deposited by Atmospheric Pressure Spatial Atomic Layer Deposition (AP-SALD) with the help of Mr. V.H.Nguyen in LMGP, as in Chapter IV. Diethyl zinc ((C2H5)2Zn; DEZ) and water vapor (H2O) were used as precursors for zinc and oxygen, respectively. The flowrate for N2 bubbling through the DEZ and H2O bubbler were 30 sccm and 100 sccm, then diluted with N2 flow rates of 270 sccm and 350 sccm, respectively.. The sample oscillated under the injection head at a travelling speed of 10 cm/s. The depositions were conducted at 220 °C for 200 cycles on cleaned corning glass substrate, reaching a thickness of about 300 nm. Before using them for SILAR coating, the substrates were cut in pieces 2.5 cm × 3.75 cm in size.

F. TiO 2 /glass substrate preparation:

With the help of Mr. R.Parize in LMGP, the TiO2 thin layers were prepared by Atomic Layer Deposition (ALD) (F200 Fiji reactor from Cambridge Nanotech). The deposition parameters have been reported. 24 Sequential pulses of tetrakis dimethylamino titanium (TDMAT) and H2O of 0.1 s followed by a purge of 10 s after each pulse were used to alternately inject titanium and oxygen precursors, respectively. The TDMAT chemical precursor was heated at 75 °C. The deposition chamber was maintained at 200 °C during the whole deposition, and its pressure was set to 11 mTorr. 500 cycles of deposition were employed to grow about 30 nm thick TiO2 layers on Si substrate. The samples were annealed in air in an oven kept at 300 °C for 3 h for crystallizing the amorphous TiO2 layer into the anatase-TiO2 phase.

With the prepared substrate, before being employed in SILAR process, 10 min of acetone and 10 min of isopropanol ultrasonic bath were required. Then, the substrates were dried by nitrogen purges and thus ready for deposition.

V.3.3 SILAR deposition parameters

The metal precursor solution was prepared by mixing 0.42 g of AgNO3 (0.0025 mol) and different amounts (mCu) of Cu(NO3)2.3H2O in 50 ml D.I. water and followed by 5 minutes magnetic bar stirring. The cation solution was prepared by adding 2 g of NaOH (0.05 mol) and 1 g of Na2S2O8 (0.0042 mol) in 50 ml of D.I. water and the mixture was stirred for 10 minutes. The two buffer solutions used were 50 ml of D.I. water. To achieve the same liquid level, all those solutions were kept in glass beakers with the same size (total volume of 80 ml). The moving down range for the substrate was set to 77 mm. To be immersed in solution, the substrates were moved down with speed of 5 mm/s. To remove the substrate from the solutions, they were withdrawn at a speed of 60 mm/s. One cycle of SILAR deposition normally began with substrate immersion in cation ion solution for t1 s, then the position of the substrate was changed to the buffer solution 1 for t2 s. Thirdly, the cation ion covered substrate was taken out of buffer solution 1 and immersed in metal ion precursor solution and maintained for t3 s. Lastly, the substrate was rinsed in buffer solution 2 for t4 s. By repeating the cycles, the coating of AgCuO2 were achieved. mCu represents the mass of added Cu(NO3)2.3H2O precursor, t1, t2, t3 and t4 are the duration in each solution. The deposited films were dried by nitrogen purge and then in 100 °C oven for 1 hour. With other parameters remaining the same, those parameters were chosen to be optimized and the results are presented in the following sections.

V.3.4 Optimization of deposition condition

The first SILAR deposition was conducted adding 0. In figure b), the BSE image presents a rather clear element contrast between grains and substrate. On those grains, the electron intensity is much brighter than the surrounding substrate, which could be either due to a loose connection between these grains and the The discontinuity of the deposited particles on substrate could be possibly due to the insufficient time for metal ions to cover the substrate surface evenly during immersion.

Therefore, increasing the immersion duration in the metal ion solution was used to try to improve that. To optimize the morphology of AgCuO2 thin films (on Cu2O/glass substrate), the immersion duration (t3) in metal ion precursor was increased from 10 s to 27 s, as shown in Table V-4. On the other hand, in order to reduce the impurity in the AgCuO2 film, the copper ion concentration was decreased gradually while the silver ion concentration was kept constant at 0.05 M. The immersion duration in metal precursor solution (t3) was thus set to 20 s. With all other deposition conditions remaining the same, several depositions were conducted as shown in Table V-5. Due to the special condition of the samples (both substrate and the deposited AgCuO2 films contain copper ), the EDS analyses can possibly penetrate the deposited sample layers thus the copper element from the substrate can be taken into account as well. Therefore, the EDS analyses of the element composition of the films is not reliable.

To evaluate the deposited phase, GIXRD was the most appropriate available technique for phase identification in our case. (sample B5) copper ion concentration. But again, this pure phase AgCuO2 film was rather difficult to be reproduced. Even if no CuO phase appears on the XRD patterns, the existence of such impurity phase inside the AgCuO2 is still believed to be below the detection limit.

While by checking the reference peaks corresponding to AgO phase, no clear evidence proves the existence of detectable amount of AgO impurity in the AgCuO2 films.

Besides the XRD analyses, Raman characterizations were also conducted. AgCuO2 powder obtained in previous work by d. Muñoz-Rojas et al. 3 , was used as reference for Raman. The AgO powder was bought from Sigma Aldrich (99.5%), thus allowing the Raman characterization in LMGP. The Cu2O reference curve was taken from reference 25 and CuO reference peaks were taken from the reference 26 . The characterizations were conducted with a blue laser (448 nm wavelength as excitation line), ×100 objective and acquisition time of the Raman spectra are rather flat at higher wavenumbers (more than 1000 cm -3 ), indicating high quality depositions without significant amount of carbon contamination. Anipsitakis et al. 28 In their work, they pointed out that the Ag 1+ ions can be oxidized by the S2O8 2-cations into Ag 2+ state, as shown in Equation V 26.

Ag + + S2O8 2-→ Ag 2+ + SO4 *-+ SO4 2- ( 26
)
Even though it has been reported that Cu 2+ ions can be also oxidized by persulfate anions into Cu 3+ , 29,30 this oxidation state is not stable in solution and the oxidizing rate is slow. Thus, in our case, we considered that only silver cations were oxidized. Therefore, during the SILAR process when the Ag 1+ and Cu 2+ ions adsorbed to the substrate surface were immersed into the NaOH solutions containing persulphate, the oxidation and coprecipitating in such basic solution occurred, thus yielding the formation of AgCuO2, as shown in Equation V 27,

Cu 2+ + Ag 2+ + 4OH -→ AgCuO2 ↓+ 2H2O (27)
Due to the highly oxidizing environment, the first few layers of Cu2O could be oxidized into CuO. Thus, besides the origin of CuO impurity phase from excessive copper ion concentration, this could be an inevitable source. To understand the role of substrate and oxidizer during the SILAR deposition of AgCuO2 films, coatings with the following parameters in Table V-6 were conducted. Two substrates were tested under similar conditions, with cation solution only containing silver or both silver and copper ions, and In addition to the optical characterizations, the deposited AgCuO2 films were also characterized by Hall effect measurement. Before the characterization, the Pt tips were directly placed on top of the deposited films, the I-V plot presented a straight linear correlation, indicating an excellence ohmic contact between the tip and the AgCuO2 film.

Thus, in this work no metal contacts were pre-deposited before characterization. Even though the AgCuO2 films were characterized above thin Cu2O layers (about 60 nm thick), due to its high resisitivity (see Chapter IV), the contribution from the bottom Cu2O layer was negligible.

With the Hall effect measurement, p type conductivity of such phase was confirmed and the results of electrical measurements are presented in Figure V-26. In the figure, it shows a trend of decreasing resistivity with decreasing copper ion concentration, from 1.6 Ω.cm down to a minimum of 0.2 Ω.cm (with copper ion concentration 0.03 M, sample B4). In addition, with lower copper ion concentration, the carrier concentration of the deposited sample decreased from 1.7 × 10 18 to 2.5 × 10 17 cm -3 . Meanwhile, the mobility was enhanced from 0.2 to more than 10 cm 2 /V.s. Each data point has been repeated twice (sample deposition and characterizations), the value presented is the minimum value obtained and 10% error bar was added for all values. With copper ion concentration of 0.05 M, equal to the silver ions, the carrier concentration reached as high as 1.7 x 10 18 cm -3 , but very low mobility of 0.2 cm 2 /V.s was observed. With reduced copper ion concentration, the carrier concentration gradually decreased down to the order of 10 17 cm -3 while the mobility was enhanced, up to more than 10 cm 2 /V.s. With copper ion concentration at 0.03 M, both the With even lower copper ion concentration, rather high mobility (around 10 cm 2 /V.s) with carrier concentration in the order of 10 17 cm -3 still can be found. Depending on the copper ion concentration, the carrier concentrations were in the range of 10 17 and 10 18 cm -3 and mobility reached a maximum of 24 cm 2 /V.s. The resistivity of those SILAR deposition AgCuO2 films varied from 0.1 to 2 Ω.cm. Comparing with the previously reported CuMO2 (M= Al, Cr, Ca, etc) ternary compounds, AgCuO2 films exhibited rather low resistivity (See in Chapter I, Table .I.1). [33][34][35][36][37] Such low resistivity could be associate with the delocalized charge and mix valence state reported. 38 In addition, comparing with those compounds, much higher mobility and lower bandgap were found with AgCuO2, which expands its potential application in both thin film transistors and photovoltaics.

V.4 Chapter VI conclusions

To summarize, in this part of the work we were aiming at the synthesis of AgCuO2 thin films. To achieve that, co-deposition by AA-MOCVD with silver and copper precursors with oxygen plasma was tested. Due to the corrosive nature of the radicals contained in the oxygen plasma, depositions could be hardly obtained. In addition, AgCuO2 phase is thermally unstable over 200 °C in air, thus precursors for both copper and silver with lower decomposition temperatures are required (better to be less than 200 °C). Combining all the difficulties, another approach was adopted. Thus alternatively, oxidation of AgCu metal alloy films (deposited by MOCVD with Ag and Cu ratio 1:1) was tested by two approaches, electrochemical oxidation and oxygen plasma treatment. From both oxidation results, only AgO, CuO phases were obtained. The formation of AgCuO2 phase during oxidation was also extremely difficult with the experimental conditions available.

Therefore, the last alternative tested was synthesizing AgCuO2 films through solution based deposition techniques. As a widely used thin film coating methods, Successive Ionic Layer Adsorption and Reaction (SILAR) process was adopted. To achieve dense and continuous thin film deposition, different substrates were tested and Cu2O/glass gave the best deposition results. With the optimized condition, dense and continuous AgCuO2 thin films can be coated (about 100 nm thick), with surface roughness RMS about 8 nm. Besides that, by comparing the deposition results on Cu2O/glass and ZnO/glass, the SILAR reaction mechanisms for AgCuO2 were also investigated, but more work is needed to completely elucidate the reactions taking place.

The optical bandgap of the deposited layer was determined using the Tauc formula, and a value of 1.2 eV was obtained assuming a direct bandgap. By comparing the Tauc plots with both direct and indirect bandgap, AgCuO2 is more likely to be a direct optical bandgap material. In addition to that, through Hall effect characterizations, p type conductivity of such AgCuO2 films was confirmed and the films presented rather low resistivity reaching down to 0.2 Ω.cm. Majority of the samples could have carrier concentration in the range of 10 17 to 10 18 cm -3 and giving mobility over 10 cm 2 /V.s. Comparing with previously reported ternary copper based oxide compounds, AgCuO2 has lower resistivity, higher mobility and lower bandgap.

evenly heated, close environment. With AgO and CuSO4 as starting materials, the phase transformation process during hydrothermal reaction was studied by in situ X-ray diffraction in the ESRF synchrotron, in Grenoble. To understand the reaction process, different factors during the reaction, including pH, temperature and pressure, were studied. Lastly, based on the experimental results, a reaction mechanism for the hydrothermal synthesis of AgCuO2 is proposed.

VI.1 Experimental setup

The in situ XRD measurements were conducted in the Swiss-Norwegian Beam lines (BM01A) at the European Synchrotron Radiation Facility (ESRF), Grenoble, France. The diffraction data were collected in transmission mode (λ = 0.6999 Å, 17714.56 eV) using the PILATUS@SNBL platform. 6 The reaction cell is illustrated in Figure VI-1 a). The starting solution/suspension (prepared just before the experiment) can be loaded through one end while the other end connects with a pressure controller during the reaction (high-pressure liquid chromatography pump, R128703 Waters 515 Hplc Wat207000). Reaction temperature was calibrated and controlled thanks to a heat blower, as described in previous works [7][8][9] . To study the reaction mechanism and activation energy, the hydrothermal synthesis of AgCuO2 with AgO and CuSO4 as starting materials was reproduced, as previously reported in the work of D. Muñoz-Rojas et al. 5 To prepare the initial slurry (suspension), 0.262 g of AgO and 0.522 g of CuSO4.5H2O were mixed and stirred in 6 ml of deionized water for 5 min, then certain amount of KOH was added and the suspension was stirred for further 5 min. The slurry was then also grinded in a tender manner (i.e. to avoid clogging the capillary phases was observed in the XRD patterns. Lastly, for the hydrothermal reaction conducted without any KOH (0 M, sample A3), the XRD reflections of the obtained powder also corresponded to Cu4(SO4)(OH)6. Again, no XRD reflections associated with any silver containing compound were identified. The recorded raw diffraction data were treated by Matlab program to obtain 2D and 3D phase transition maps. In addition the treated data with "_.xy" format were converted into "_.raw" files (with wavelength 0.6999 Armstrong) with DIFFRACTION PLUS V.04. Then phase identification was carried out using the software EVA V.02. The details of data treatment will be presented in the appendix at end of the chapter. All diffraction signal acquisition duration in S1, S2 and S3 are 5 s, thus their reflection intensity are comparable.

In the first case we followed the standard reaction parameters (sample S1), which was also defined as standard reaction condition in this work. 5 The phase evolution and identifications are presented in Throughout the characterization, the intensity of AgO phase reflections decreased as the reaction proceeded. However, at the end of this hydrothermal reaction, neither the reflection corresponding to AgCuO2 nor any silver containing compound was observed. The absence of silver containing phase and the resulting Cu4(SO4)(OH)6 crystalline phase in solution were in agreement with the XRD results obtained from sample A2, synthesized in the lab. The explanation for the absence of silver containing phases can be found from the work of N.N.Lesnykh et al.. 11 As it's know that, metal oxides are partially soluble in basic solution to form metal hydroxide complex in solution. Here is the same for AgO, even though the solubility is rather low. However, in that work, it pointed out due to the presence of SO4 2- anions in solution, the formed [AgOH(SO4)] 2-ads complex was much more soluble than [Ag(OH)2]ads complex in 0.1 M NaOH solution. In our case, the SO4 2-ions were also included in the slurry (from CuSO4) and the KOH concentration was also in the same level (0.3 M), thus this could explain the dissolution of AgO. On the other hand, the Cu4(SO4)(OH)6 precipitated as the final product, which also confirmed the existence of such metal sulfate hydroxide complex. As shown in Equation VI 1 and 2, both Cu 2+ and SO4 2-ions are able to hydrolysis with water, however since H2SO4 is a strong acid, thus its SO4 2-ion is a weak basic, thus the overall CuSO4 solution presents acidic pH. In the reaction slurry, AgO was also added as reactant, thus it was dissolved in the acidic solution as shown in Equation VI 3. Therefore, the H + ions were consumed and balance in Equation VI 1 shifted to the Cu(OH) + complex. Thus, combining the SO4 In this starting slurry, there is always an equilibrium between the metal oxide and the metal hydroxide complex, as shown in Equation VI 6 and 7. 13 Even though, for AgO, such solubility is rather low. In order to achieve AgCuO2 thin film deposition in MOCVD, besides the work about silver and silver oxides, the copper oxide thin film optimization was also required. However, thanks to the long working experience in LMGP about copper oxide, the deposition of CuO was already well studied and optimized. Thus, we turned our interests to Cu2O thin film optimizations and attempted to achieve Cu2O films with better electrical properties.

During this part of work, to achieve Cu2O films with better homogeneity, the configuration of the AA-MOCVD system in LMGP was optimized. With a careful choice of the copper precursor (using Cu(II) trifluoroacetylacetonate (CuF)), high crystalline, homogenous and carbon-free Cu2O films were deposited. Based on that, the effects of oxygen ratio, silver nanoparticle incorporation and carrier gas humidity were investigated during the MOCVD.

With slight higher oxygen ratio (still in the range of Cu2O deposition), Cu2O films with rougher surfaces and cubic grains were deposited. In addition, to incorporate the Ag nanoparticles in Cu2O film, the silver precursor AgT was mixed into CuF solution with various concentrations. With a higher AgT concentration, more Ag nanoparticles were incorporated in the Cu2O films. With more Ag nanoparticle incorporation, Cu2O films with bigger size, (111) oriented grains and higher deposition rates were observed. Lastly, with a tunable humidity in the carrier gas, depositions were carried out. The incorporation of humid carrier gas during MOCVD had a great impact on the deposited Cu2O films. First, the surface of the films presented 'pyramids-like' structure and those types of films were identified with (111) texture by XRD, electron diffraction in TEM and EBSD. In addition, higher deposition rates were observed with more humid carrier gases. By the TEM cross-section view of the sample, big grains (width over 300 nm) with columnar structures were observed. From the point view of electrical properties, Cu2O films deposited with humid carrier gas presented high carrier mobility, reaching a maximum of 17 cm 2 /V.s. As well, the conduction characteristics of those (111) oriented Cu2O films were further examined by electrical AFM.

Through this, a higher conductivity was observed on the grain boundaries than on grains.

The ASTAR phase mapping with ultra-fine resolution revealed a higher distribution of CuO phase in the grain boundaries which might be responsible for this abnormally high conductivity.

As the last part of work in this section, the effects of humidity on the initial thin film nucleation were studied. Cu2O films deposited for short durations (10 min ~ 50 min ) with and without humidity carrier gas were characterized by AFM. From the characterization, General conclusion and perspectives -214 -much higher roughness of Cu2O films were found when deposited with humid carrier gas, indicating a Stranski-Krastanov growth mode. Meanwhile, without humid carrier gas, the Cu2O films tended to have layer-by-layer growth. Furthermore, based on the classic nucleation theories, the Gibbs free energy during the deposition process with different conditions (oxygen ratio, silver incorporation and humidity) was discussed. Through the model established, from the point view of thermodynamics, the larger grain sizes and high deposition speed were explained.

AgCuO2 thin film coating and electrical and optical property characterizations.

Finally, with the foundation of the work about silver, silver oxide and copper oxides, we moved towards AgCuO2 film deposition. First of all, an attempt to deposit AgCuO2 in AA-MOCVD was tested. To conduct such tests, a mixture precursor solution of CuF and AgT was used for the co-deposition. During the deposition process, the oxygen plasma provided the high oxidizing species. Those co-depositions were tested at the precursor decomposition temperature limit, however similar to the silver oxide deposition, the temperature incompatibility between precursor and AgCuO2 phase made such co-deposition extremely difficult. In addition, besides the oxidizing species provided from oxygen plasma, highly corrosive radicals were also emitted and they made the deposition process even more complicated. Considering both the durability and the cost of precursors, a lower cost solution thin film coating technique was adopted instead, namely was Successive Ionic Layer Adsorption and Reaction (SILAR). Metal ion precursor solution was made up with almost equal ratio of AgNO3 and Cu(NO3)2. The cation solution contained 1 M NaOH solution. By adding sodium persulfate in the solution, the strong oxidation ability was provided. Since SILAR deposition is a rather substrate sensitive technique, different substrates with/without seed layers were tested. As a result, Cu2O/glass substrate was able to deliver a high quality coating. Even though with ZnO/glass substrate coatings can also be obtained, continuous films were rarely achieved. The deposited films on Cu2O/glass were identified by GIXRD to be the AgCuO2 phase, even though with a small quantity of CuO as impurities. In the optimized case, a continuous and dense 100 nm thick AgCuO2 thin film was coated. The RMS value was as low as 8 nm, indicating a rather smooth surface. The deposited AgCuO2 films were then characterized by UV-Vis spectroscopy. With the transmittance data, the Tauc plots were used for estimating the bandgap with both direct and indirect formula. By comparing the obtained Tauc plots in those two cases, a 1.2 eV direct bandgap was obtained.

4 .

 4 Outre l'optimisation de Cu2O, le dé pôt d'AgCuO2 par MOCVD a é galement é té abordé . Pour ce faire, le dé pôt de films minces d'oxyde d'argent et d'argent a é té pré alablement optimisé . Pour cela, deux nouveaux pré curseurs d'argent, à savoir les triglymes Ag (hfac) phé nanthroline et Ag (hfac) ont é té synthé tisé s et complè tement caracté risé s. On peut obtenir des revê tements Ag de haute qualité avec les deux pré curseurs. Les films d'oxyde d'argent ont é té obtenus par oxydation é lectrochimique et traitement par plasma à l'oxygè ne des revê tements pré -dé posé s Ag.

  Figure I-1 d), which shows the density of states (DOS) for both Ag2Cu2O3 and Cu3O4.

Figure I- 1 .

 1 Figure I-1. Details about Ag2Cu2O3 phase. a) Schematic illustration of crystal structure. b) Thermal Gravimetric Analysis (TGA) shows rapid decomposition at 250 °C in reducing atmospheres. c) Band structure of Ag2Cu2O3 obtained by Extended H𝑢̈ckel Tight Binding (EHTB) calculation. d) Density of States (DOS) for Ag2Cu2O3 (left) and Cu3O4 (right) obtained from EHTB calculation. The shading areas refer to the contribution of the metal ions marked on the top corner. In both c) and d), dash line represents position of the Fermi level.46

Figure I- 3 .

 3 Figure I-3 b), the Ag 3d range contains a spin-orbit coupling doublet, Ag 3d3/2 and Ag 3d5/2.

Figure I- 3 .

 3 Figure I-3. XPS characterization of electrodeposited AgCuO2 (Ag2Cu2O4) phase. a) the full XPS spectra and the detailed spectra at a) Ag 3d, c) Cu 2p and d) O 1s regions. Open circles represent experimental data and solid and dashed lines are fittings.50

  Figure I-4 a), in which a layered structure with large surface area is presented. In Figure I-4 b), the diffraction pattern and crystal structure are presented.

Figure I- 4

 4 Figure I-4. a) SEM image of Ag2MnCuO4 crystals. b) the XRD pattern and inserted schematic crystal structure.51

  conductivity and Seebeck coefficients were measured. The electrical conductivity was measured by standard four probe Pt tip device and the Seebeck coefficient was characterized by thermopower. High electrical conductivity 3.2 × 10 -1 S/cm was obtained for AgCuO2, while AgCu0.5Mn0.5O2 presented much lower (3 orders of magnitude) conductivity, as shown in Figure I-5.

Figure I- 5 .

 5 Figure I-5. Thermopower characterization ∆𝑉 = 𝑓 (∆𝑇) for: a) AgCuO2 and b) Ag2MnCuO4. Four probes DC conductivity measurements conducted in room temperature as inserted images. 43 During the thermal power measurement, thermovoltage ∆𝑉 and temperature gradient ∆𝑇 were followed simultaneously from the decay of short pulse to 100 °C. Therefore, the Seebeck coefficient was extracted, as shown in Figure I-5. The thermopower measurements show that the nature of conductivity is hole-type (p) with positive Seebeck coefficient of value +2.46 for AgCuO2 and +78.83 μV/K for AgCu0.5Mn0.5O2. Such a high conductivity and low Seebeck coefficient for AgCuO2 (3 orders of magnitude higher conductivity than AgCu0.5Mn0.5O2) can be attributed to the charge delocalization in its mix-valence bands.

Figure I- 6

 6 Figure I-6. a) SEM image of a AgCuO2 particle from hydrothermal reaction contacted using ion assisted deposition of Pt. The insets show a close-up of the particle from different perspectives, as well as an EDX spectrum of the contacted particle showing the corresponding peaks for Ag and Cu. b) V versus I plots obtained for contacted particles using FIB-Pt nanocontacts. The inset shows the 4 probes contacting the patterned Al microelectrodes in the wafer, to which the metallic nano-wires, Pt in this case, are connected.36

Figure I- 7

 7 Figure I-7. a) Tauc plot using direct bandgap formula from transmittance data of a 460 nm thick Ag2Cu2O3 thin film at temperatures ranging from 4 K to 300 K and b) the corresponding Varshni fitting of the bandgap versus temperature. 53 In addition, from the work of E.Lund et al., 53 the Ag2Cu2O3 thin film electronic and optical properties were comprehensively studied. As shown in Figure I-7 a), presenting transmittance data of 460 nm thick Ag2Cu2O3 film, a bandgap about 2.2 eV was obtained. While in Figure I-7 b), a decreasing bandgap with increasing temperature was observed. In addition for bandgap determination, Photoluminescence (PL) characterizations were conducted at various temperatures ranging from 10 to 300 K and results are shown in Figure I-8 a). As it is shown, the PL curves of Ag2Cu2O3 demonstrate a low luminescence yields.

Figure I- 8 . 53 Figure I- 9 .

 8539 Figure I-8. Photoluminescence (PL) characterization of a 460 nm Ag2Cu2O3 thin film. a) Full PL curve between temperature 10 and 300 K using 325 nm as excitation line. b) Arrhenius plot of integrated PL intensity at 2.2 eV.53

Figure I- 10 .

 10 Figure I-10. Calculated species distribution versus pH in solution contains 3.75 mM AgNO3, 3 mM Cu(NO3)2, 0.5 M NH3.H2O at room temperature for a) Ag (I) and b) Cu (II) species. c) Inluence of pH on solubility of Ag2O, AgOH, CuO and Cu(OH)2 in solution. d) Cyclic voltammograms (CVs) measured with ITO electrode (1.5 cm 2 ) in such solution with scanning speed 10 mV/s. Even though Ag2Cu2O3 in thin film was reported by sputtering technique at low temperature (no more than 100 °C), due to the critical requirements for high oxidation environment to form AgCuO2 phase plus its thermal instability, currently there is no dense and continuous thin films reported. But, recently from a work of Q. Lu et al., 50 AgCuO2 nano plates were successfully deposited by electrochemical deposition on ITO substrate with incorporation of ammonia in the process as electron pair donor ligand. In order to chose a proper pH for an easier control of the stoichiometry of the final product, the influence of solution pH on Ag(I) and Cu(II) species were investigated and calculated in a solution with 3.75 mM AgNO3, 3 mM Cu(NO3)2 and 0.5 M NH3.H2O at room temperature, as shown in Figure I-10 a) and b).The concentration of Ag 1+ species decreases for pH values exceeding 5 and the Ag(NH3)2+ species concentration keep increasing until pH reaches 8 and stabilizes. While the Ag(NH3) + only exists when pH is between 5 and 8. In the case of Cu(II), due to the strong bonding between Cu ions and ammonia group, the ammonia ligand keeps increasing in the Copper ammonia complex when pH increases, therefore we see several plateaus at different pH ranges. Figure I-10 c) demonstrates the solubility evolution of both Ag and Cu oxide/hydroxide by the influence of pH in such solution. The silver species presents higher solubility in a general view, therefore to maintain the correct stoichiometry in AgCuO2 compound, slightly more Ag precursor should be added. Lastly, with 10 mV.s -1 scanning

Figure I- 11

 11 Figure I-11. a) XRD pattern of electrodeposited on ITO substrate. b) TGA analysis of AgCuO2 nano plates under N2 atmosphere. Within the 0.25 M NaOH solution, by applying potential in the range of H2O and NH3 oxidation, AgCuO2 nanoplates were coated on ITO electrode. The XRD pattern in Figure I-11 a) presents a rather pure AgCuO2 compound phase and the TGA analysis in figure b) shows a rapid decomposition over 250 °C of this compound.

Figure I- 12

 12 Figure I-12. a) and b) morphology of AgCuO2 nanoplates by SEM. c) is the SAED pattern of AgCuO2 plateand d) is the HRTEM image of a nanoplate.50

Figure I- 13 .

 13 Figure I-13. Cubic crystal structure of Cu2O. 7 The crystal structure of Cu2O is shown in Figure I-13. Cu2O crystalizes in a cubic structure (a = 4.2696 Å, Pn3 ̅ m), with unit cell containing 4 copper atoms inside, 1 oxygen atom at center and 8 more oxygen atoms at each corner, shared with 8 other unit cells, and thus giving 2 oxygen atoms per unit cell.

Figure I- 14 .

 14 Figure I-14. Diagram of a) Conduction band and valence band formation in Cu2O. b) Acceptor levels in Cu2O.

Figure I- 15 .

 15 As it is shown, in the Cu-rich/O-poor system (Figure I-15 a)), the overall lowest formation enthalpy is from formation of copper vacancy ∆𝐻(𝑉 𝐶𝑢 0 ) = 0.7 𝑒𝑉 . Another type of vacancy creation mechanism, i.e. neutral split vacancy, can be generated with higher formation energy ∆𝐻(𝑉 𝑠𝑝𝑙𝑖𝑡 0 ) = 1 𝑒𝑉. Meanwhile, the oxygen interstitials have formation energy 1.3 eV and 1.8 eV, separately. In the other case (Figure I-15 b)), the Cu-poor/O-rich diagram has also shown that the overall lowest formation energy corresponds to copper vacancies. 58 Therefore, in either deposition conditions, copper vacancy formation has always the lowest formation energy, thus it explains the origin of p type conductivity of intrinsic Cu2O. Another important principle of Cu2O p type conductivity is the self-compensation mechanism. As shown in Figure I-15, the formation energy of charged defects changes with regard to the Fermi level. When the Fermi level is over 1 eV, the formation energy of VCu 0 decreased to zero or even negative, implying that to compensate the increasing Fermi level, the crystal lattice produces positively charged copper vacancy. Such compensation mechanism is called self-compensation. Therefore, it explains the difficulty of doping Cu2Ointo n type semiconductor (the same principle is behind the difficulty to have intrinsic p type ZnO for example).58

Figure I- 15 .

 15 Figure I-15. Formation energy of different vacancies in copper rich a) and copper poor b) conditions. The signal c, d represents different interstitial sites inside the unit cell.58

  by fitting the resistivity of Cu2O films sputtered at different temperatures, two conduction path models were found. As shown in Figure I-16, at higher deposition temperatures the conductivity is mainly controlled by the acceptor level EA (VCu) with activation energy 0.15 eV, while at lower deposition temperatures, photon energy Wph with activation energy 0.03 eV becomes the dominant factor. It indicates that forCu2O films deposited at lower temperatures, the hopping mechanism induced from the activation of phonon energy level is the major mechanism for Cu2O conductivity, while at higher temperatures the conductivity is determined by the acceptor level. Besides the reason of deposition temperature, there are other factors that can largely affect the produced Cu2O films, such as morphology, grain size and defects, etc.

Figure I- 16 . 66 I. 3 . 1

 166631 Figure I-16. Temperature dependency of dark conductivity of a 270 nm thick Cu2O film. With higher deposition temperature, the conductivity is mainly controlled by the copper vacancy VCu 0.3 eV above the VBM while at lower temperature, the conductivity is limited by the hopping mechanism. 66 I.3.1 Application of Cu2O in thin film transistor

Figure I- 17 .

 17 Figure I-17. Diagram of Cu2O transistor a) top view and b) cross section view. c) the transfer curves of the fabricated Cu2O TFT with different active layer thickness and drain voltage VD -1 V. W and L are the width and length of the fabricated device (W: 1000 μm, L: 100 μm).67By fixing the drain to source voltage to -1 V, the drain current vs Gate voltage curves were measured and are presented in Figure I-17 b). As it's shown, with a 45 nm thick Cu2O film as the active layer, it shows a clean transfer function with low threshold slope (TS) around 1.6 V/dec. Furthermore, it reported the field-effect mobility (μFE) ~ 0.06 cm 2 /(V.s), on-off

Figure I- 18 .

 18 Figure I-18. Electrical characterizations of Cu2O based TFT fabricated at room temperature and annealed at 200 °C. a) IDS -VDS output curve; b) Transfer curve characterization IDS -VGS with drain voltage VD = -5 V and field effect mobility μEF VS gate voltage VGS.[START_REF] Kim | p-Channel oxide thin film transistors using solution-processed copper oxide[END_REF]

  (PEC) cell. It's a technique which combines solar energy collection and water electrolysis onto a single photoelectrode. The PEC cell relies on minority carriers generated by light absorption of the electrode and then being pumped into the liquid, thus a potential difference is created between the junction of semiconductor electrode and liquid, as shown in Figure I-19 a). 69 Thus, a redox reaction can be conducted at the vicinity of the electrode and one of the most studied reaction is for reducing H + into H2, namely, hydrogen generation. To achieve that, the energy band position of the semiconductor material should be carefully chosen. For an ideal semiconductor material for water electrolysis, its energy band should have negative potential versus hydrogen reduction and positive potential regarding to oxygen generation. With a comprehensive review of the energy band diagram of different semiconductors, Cu2O offers features. As shown in Figure I-19 b), Cu2O has an ideal energy band position for PCE cell hydrogen generation. Comparing to hydrogen reducing reaction potential, conduction band of Cu2O lies 0.7 V below while the valence band stays above the potential for oxidation reaction to produce O2. 69 Thus, it's an attractive material for H2 generation without external voltage bias or low voltage bias. But Cu2O as the photoelectrode has problems of stability in solution and tends to be reduced into metallic Cu. Thus, finding a proper protection method for the Cu2O from reduction or oxidation in solution is required.

Figure I- 19

 19 Figure I-19. a) working principle of a PEC cell. b) Energy diagram of Cu2O verses the potential of redox reactions of H2O. 69 In the work of A.Paracchino et al. 70 , on top of electrodeposited Cu2O, different combinations of protection layers were tested and bilayers of n type Al doped ZnO and TiO2 are confirmed for an effective prevention of the surface from oxidation. Pt was electrodeposited functioning as a catalyst, as shown in Figure I-20 a).

Figure I- 20 . 2 )Figure I- 21 .

 20221 Figure I-20. PCE cells made from 1.3 um electrodeposited Cu2O layer, bilayers of Al doped ZnO and TiO2 as the protective layers and Pt nanoparticles as catalyst. a) schematic presentation of the cross section and b) SEM image of the top-view. 70 With (111) oriented Cu2O surface, the photo electrochemical responses were characterized with bare Cu2O electrode and with the bilayer protection, the results are shown in Figure I-21. With bared Cu2O as the electrode, a photocurrent of -2.4 mA cm -2 was generated at 0.25 V regarding to the reversible hydrogen electrode (RHE), in a 1 M Na2SO4 electrolyte at pH 4.9. The corresponding J-V curve is presented in Figure I-21 a), which shows a rapid decay of photocurrent. The major reason for this rapid photocurrent decrease can be attribute to the reduction of Cu2O electrode in solution, shown in Equation 2, 𝐶𝑢 2 𝑂 + 2𝑒 -+ 2𝐻 + → 𝐶𝑢 + 𝐻 2 𝑂(2)

  photocurrent of PEC cells with Cu2O can be enhanced up to -0.47 mA cm -2 . The energy band position of Cu2O and CuO with respect to NiO is shown in Figure I-22. Therefore, even though both Cu2O and CuO are suffering from instability in solution, CuO seems to be more stable and this can inspire new choices or structures for PEC cells to solve the stability issue and meanwhile obtaining higher photocurrent.

Figure I- 22 .

 22 Figure I-22. Energy band diagram of Cu2O and CuO versus NiO as the catalyst. Therefore, in the work of Y.Yang et al., 72 bilayer of Cu2O/CuO was used at the photo electrode, CuO as the outside layer in contact with solution. The Cu2O layer is synthesized through a repeated double-potential pulse chronoamperometric (r-DPPC) deposition method. Cu2O/CuO bilayers were acquired through thermal oxidation of the Cu2O. The PEC cell structure is shown in Figure I-23. With such combined bilayer structure, the light spectrum able to be absorbed was broadened, longer electrical field for efficient charge separation was obtained and the overall electrode could have better stability than a single Cu2O phase.

Figure I- 23 .

 23 Figure I-23. Schematic presentation of energy diagram of Cu2O/CuO bilayer photo electrode. The photocurrent density with regarding to RHE is shown in Figure I-24 a), in which the pure Cu2O photoelectrode (red line) has much lower current density (-0.21 mA cm -2 ) and decaying much faster than the other two. Due to a smaller bandgap of CuO and better stability in solution, as previously mentioned, the current density with CuO electrode delivers much higher current density and more stable performance. But with the Cu2O/CuO as electrode, as expected the photon current was largely enhanced up to -2.47 mA cm -2 . In addition, the stability of the composite electrode is largely improved as well. Moreover, with such composite structure more light absorption and larger charge separation space, the conversion efficiency is largely improved than Cu2O as electrode and better than pure CuO, as shown in Figure I-24 b).

Figure I- 24 .

 24 Figure I-24. Comparison of the performance of photoelectrode made from Cu2O (red line), CuO (dark line) and Cu2O/CuO bilayer (blue line). a) comparison of the current density versus potential and b) photo conversion efficiency of different photo electrode versus potential.

  Figure I-25 a) and b), they illustrate planar Cu2O/ZnO and Cu2O/ZnO nanorods solar cell structures. Figure c) is the energy diagram of such Cu2O/ZnO p-n junction, in which xp represents the full depletion length in the p part, xn is defined as the depletion length in ZnO and Vbi is built in potential in such photovoltaic devices.

Figure I- 25 .

 25 Figure I-25. Schematic presentation of a) Cu2O/ZnO bilayer structure, b) Cu2O/ZnO nanorods structure and c) energy diagram. 79 During this study, by varying the Cu2O layer thickness deposited by electrochemical deposition, the short circuit current Jsc , open circuit voltage Voc and fill factor FF were evaluated. The depletion length x in both p and n parts can be calculated by applying Poisson's Equation at the p-n junction interface and giving the build in potential ratio relation following Equation 3:

  3 to 3 μm. For bilayer solar cell structures, the open circuit voltage evolution is clearly presented in Figure I-26 a). Before reaching the required thickness 2.7 μm to form a full depletion layer, the Voc increased rapidly with regarding to a thicker film, in accordance with the theory prediction. While in Figure I-26 b

Figure I- 26 .

 26 Figure I-26. Evaluation of Cu2O/ZnO bilayer and Cu2O/ZnO nanorods solar cell performance depending on the absorber layer Cu2O thickness. a) open circuit voltage Voc , b) short circuit current Jsc and c) fill factor FF evolution with regarding to Cu2O thickness change.[START_REF] Musselman | Incompatible Length Scales in Nanostructured Cu2O Solar Cells[END_REF] On the other hand, in the Cu2O/ZnO nanorods structure, the Voc remained stable and Jsc decreased with thicker Cu2O layer. The stable voltage value implied that the distance between the rods plays an important role in the depletion zone formation. In another words, if the distance between the rods is smaller than the required full depletion length, the final Cu2O/ZnO rods structure couldn't form a p-n junction with a complete deletion zone, regardless the thickness of the top layer. While the decrease of the short circuit current in Figure I-26 b) could be due to inefficient charge collection with much thicker Cu2O top layer.

a

  mapping of the performance for Cu2O/ZnO bilayer cell with various thickness parameters was established, as shown in Figure I-28.

Figure I- 27

 27 Figure I-27. a) scheme of Cu2O/ZnO bilayer solar cell structure with gradient layer thickness. b) the cross section view of bilayer structure and corresponding elemental analysis by EDS.[START_REF] Panigrahi | Oxide-Based Solar Cell: Impact of Layer Thicknesses on the Device Performance[END_REF]

Figure I- 28 .

 28 Figure I-28. Mapping of Cu2O/ZnO bilayer solar cell performance with gradient thickness. a) Jsc, b) Voc, c) Maximum output power (Pmax) and d) fill factor (FF). e) summary of Voc and f) Jsc versus Cu2O thickness evolution.[START_REF] Panigrahi | Oxide-Based Solar Cell: Impact of Layer Thicknesses on the Device Performance[END_REF]

  Figure I-28, a) presents the Jsc mapping of the gradient structure. Figure b) shows the trend of Voc on the gradient surface. In both cases, solar cells with thicker ZnO and Cu2O layers at the top-left corner have generally better performance, which can be confirmed by figure c) and d). However by checking the color bar, it' easy to find that the optimal performance point was not the ones with the thickest layers. This result are presented with a clear view in figure e) and f). In figure e), the optimal Voc 0.6 V achieved was with about 250 nm ZnO and 600 nm thick Cu2O. With even thicker Cu2O layer, the Voc decreased slightly, which is in accordance with the previous results shown Figure I-27. Thus, since the electronic properties of Cu2O films deposited by Spray Pyrolysis and by MOCVD are rather similar, those thickness versus performance maps is also a good indication for Cu2O based solar cell devices fabricated by MOCVD.

Figure II- 1

 1 Figure II-1. a) Working principle of PI-MOCVD. b) Schematic presentation and c) photo of PI-MOCVD system. INSET: Ag coating obtained with this system in which a homogeneous 5 cm diameter coating can be observed. 1 and 2 are the injector lines, 3 is the gas line for oxygen and 4 is the deposition chamber. The photo and schematic presentation of the PI-MOCVD system used in this work is shown in Figure II-1. Figure a) illustrates the process of deposition in PI-MOCVD, where injected droplets are immediately evaporated in vacuum injection lines, thus vapor phase precursor are transported by carrier onto the hot substrate, where deposition takes place. The PI-

Figure II- 2 .

 2 FigureII-2. Basic working principle of thin film deposition in AA-MOCVD.5

Figure II- 3 .

 3 Figure II-3. Original configuration of the AA-MOCVD system used.

Figure II- 4 .

 4 Figure II-4. Details of the final optimized AA-MOCVD configuration.

II. 1 . 3 1 . 3 . 1

 13131 Figure II-4 b) presents a close up view of the carrier gas mixer, which has been designed to avoid big droplets in the precursor flow and condensation on the tube to have direct impact on deposition quality. The big and fast droplets in flow are mostly removed by hitting the Si substrate that acts as a filter and the condensations above are collected through the tube at the bottom side while smaller droplets are introduced in the reaction chamber. Film deposition II.Preparation of starting precursor solutions: In the case of Ag, several different precursors were evaluated: a). Ag acetate (AgAc, 99%, Sigma Aldrich). 0.02 M solutions were prepared by adding 1.67 g (0.01 mol) of AgAc in 500 ml of butanol-01 or ethanol, with continuous magnetic stirring with a speed of 300 cycle/min. AgAc can slowly be dissolved in alcohol solvents, but a long stirring duration is required. To facilitate this dissolving process, both ethanolamine and ethylene diamine can be added as dispersants. The solution presented a clear transparent color. The molecular structure of silver acetate is presented: b). Ag nitrate (99.5%, Sigma Aldrich). Solutions with 0.02 M concentration were prepared by mixing 1.7 g (0.01 mol) of AgNO3 in 500 ml of ethanol. The mixture was stirred with magnetic bar with speed 300 cycles/min for 30 min. The fully dissolved solution was clear yellow. The structure of the precursor is shown as below: c). 2,2,6,6-Tetramethyl-3,5-heptanedionato Ag(I) [Ag(TMHD)] (99.9%, Strem Chemical). Another silver precursor solution with 0.01 mol/L concentration was prepared by adding 1.46 g (0.005 mol) of Ag(TMHD) in 500 ml of ethanol (or butanol-01). The mixture was stirred for more than 60 min with magnetic bar at 300 cycles/min, dispersants such as ethanolamine was needed to obtain a clear grey solution. The molecular structure of the Ag(TMHD) is shown below: d). Ag(hexafluoro acetylacetonate) phenanthroline (AgP, new precursor obtained during

  dispersant (for 500 ml of 0.01 M solution, 2.5 ml of ethylene diamine). The solution became clear blue right after the dispersant was introduced. After 10 mins of further stirring, the solution was ready for deposition. The precursor molecule structure is shown below: b) Cu(II) trifluoro acetylacetonate (CuF) (97%, STREM Chemical). The precursor solutions with 0.01 mol/L concentration were prepared by adding 1.85 g (0.005 mol) of CuF into 500 ml alcohol solvent (ethanol in this case) with magnetic bar stirring for 10 min with a speed of 300 cycle/min. The precursor rapidly dissolved in alcohol solvent and the solution presented a crystal clear green blue color after 2~3 minutes. The molecule structure is shown below:

Figure II- 5 .

 5 Figure II-5. Scheme of the pressure gauge used in the system with a 5 ° tilted capillary filled with red liquid.

Figure II- 6 .

 6 Figure II-6. Scheme of the electrochemical setup with a potentiostat and the electrochemical cell with the working, counter and reference electrodes. 7 There are several protocols that are currently used when performing electrochemistry. The most commonly used modes are linear sweep voltammetry (LSV), cyclic voltammetry (CV), chronoamperometry (CA) and chronopotentiometry (CP). Among those different working modes, the LSV and CV modes are normally used to identify electrochemical processes. In LSV mode, as shown in Figure II-7 a), the potential linearly increases from potential E0 to E1 (forward scan) and back to E0 (backward scan), E0 normally representing the open circuit voltage. Through recording the current versus voltage, the oxidation and/or reduction potential at which electrochemical processes occur can be identified. Conversely, in CV mode the potential is linearly increased from E0 up to E1 and then falls back to negative potential E2 and finally returns to E0, as presented in Figure II-7 b). Through the full cycle characterization, not only the oxidation/ reduction potentials can be distinguished, but also the reversibility of the reactions can be evaluated, as shown in Figure II-7 c). Each peak on

Figure II- 7

 7 Figure II-7. a) Potential versus time in LVS mode. b) Potential versus time in CV mode and c) is the correspondent CV curve (voltammogram) with single electron transfer, current versus potential.

Figure II- 8 .

 8 Figure II-8. Experimental setup for electrochemical oxidation of Ag/AgCu films. In this work, electrochemistry was used for oxidizing conductive metallic Ag or Ag-Cu films and for the direct deposition of AgCuO2 thin films from solution. For the oxidation purpose, a typical three-electrode cell configuration was used, as shown in Figure II-8, in which the Ag/Cu films deposited by MOCVD were the working electrode and Pt was used as counter electrode (left), with a AgCl/KCl 3.5 M reference electrode (right). The three electrodes were placed in a three neck flask containing a 1 M NaOH solution as the electrolyte. The

Figure II- 9 .

 9 FigureII-9. Working principle of atmospheric plasma.8 In this work, the atmospheric plasma generator was used to produce a highly oxidizing atmosphere.9,10 The oxygen plasma head (from AcXys technologies) is placed in the AA-MOCVD deposition chamber, as illustrated in Figure II-10.

Figure II- 10 .

 10 Figure II-10. Scheme of the setup used to carry out oxygen plasma oxidation of metallic Ag/Cu films.

  Ionic Layer Adsorption and Reaction (SILAR) deposition is a solution based technique, widely used for metal oxide and sulfide thin film coatings. Commonly, there are at least three types of solutions used in the SILAR process, namely metal ion solution, cation solution and buffer solution. The buffer solution normally consists of deionized water or a water-alcohol mixture . The deposition mechanism relies on the electric static affinity between the ions in solution and the substrate 11 . The thin film deposition mechanism in SILAR is illustrated in Figure II-11, the reactions between metal ions and negative cations lead to the precipitation and thus the deposition of the targeted material, as shown in Equation II 5.

Figure II- 11 .

 11 Figure II-11. Schematic illustration of thin film coating during one SILAR cycle. (a) cation adsorption; (b) rinsing; (c) anion reaction; and (d) rinsing. 11In this work, for the first time, the SILAR deposition technique had been used to synthesize AgCuO2 thin films. The experimental setup used is mainly composed of a motor in the Z axis to move the substrate up and down and a rotating platform on the X-Y plane to change the position of the three solutions in a sequential order. More detailed experimental process, parameters and results will be discussed in Chapter V.

(

  v=1) of a vibrational mode with energy Einc-Escat. If the scattered photon has more energy than the incident photon, it is the anti-Stokes Raman process, and the molecule has relaxed from an excited vibrational state (v=1) to its ground state (v=0). The energy of the vibration is equal to Escat-Einc. Thus, the Stokes and anti-Stokes parts of a Raman spectrum can be symmetrically observed on both sides of the Rayleigh line but with different intensities. In practice, the more commonly measured in Raman spectroscopy is the Stokes spectrum which has a higher intensity. It is to be noted that the intensities of the Raman lines are very weak in comparison with the very intense Rayleigh line (at least ×10 -6 ). The energy lost (Stokes) or gained (anti-Stokes) by the photons in the scattering process is called the Raman shift.Raman shifts () are commonly expressed in wavenumbers (cm -1 ) such as  = 1/. The Raman spectrum corresponds to the wavenumber dependence of the Raman scattered intensity at a given incident wavelength. It is characteristic of a molecule or a material for a given crystallographic structure.

Figure II- 12 .

 12 FigureII-12. Schematic representation of the Rayleigh and Raman scattering, and of the IR absorption processes.14

Figure II- 13 .

 13 Figure II-13. Scheme of the Bragg diffraction. The two scattered X-Ray beams undergo constructive interference when 2dhkl sinθ = nλ. (image from: https://fys.kuleuven.be/iks/nvsf/experimentalfacilities/x-ray-diffraction-2013-bruker-d8-discover) A typical X-Ray diffractometer setup is schematically presented in Figure II-14. Based on the desired information to be obtained, different XRD configurations can be applied. The Bragg-Brentano geometry (see Figure II-14) is the most typically used configuration.

Figure II- 14 .

 14 Figure II-14. Schematic illustration of the Bragg-Brentano geometry. (image from: http://li155-94.members.linode.com/myscope/xrd/background/machine/) With such a setup, only the diffraction planes parallel to the sample surface can undergo constructive interference, thus resulting in diffraction peaks. As shown in Figure II-15, crystal A will diffract in Bragg-Brentano geometry while crystal B won't be able to satisfy the diffraction condition thus not appearing in the diffraction pattern.

Figure II- 15 .

 15 Figure II-15. Scheme to illustrate that only crystallographic planes parallel to sample surface diffract in Bragg-Brentano geometry.

  beam is almost parallel to the sample surface, thus the crystal planes parallel to the sample surface diffract with much weaker signal. Contrarily to the Bragg-Brentano geometry, under GI geometry, the grains with crystal planes not parallel to the sample surface are the main contributors to GI diffraction signal. In this work, most GIXRD patterns were acquired with a Bruker C40 D8 with incidence and exit angles maintained around 0.5 °. There are also part of the GIXRD patterns collected with a RIGAKU Smart lab diffractometer.

Figure II- 16 .

 16 FigureII-16. Schematic illustration of Grazing Incidence XRD characterization. ω is the incidence angle and s is the vector perpendicular to the diffraction plane.

Figure II- 17 )

 17 FigureII-17) taking place at different penetration depths and providing complementary information. As it is shown, due to the strong energy of the primary electron beam, Auger Electrons are firstly excited at the very surface of the sample, next are the secondary electrons (few nm range) which are used for imaging the sample. In a deeper penetration depth, back scattered electrons and characteristic X-Rays are also generated, both carrying information on the sample composition.

Figure II- 17 .

 17 FigureII-17. Schematic illustration of the primary electron beam interacting with the surface of the sample.20

Figure II- 18 .

 18 Figure II-18. Scheme of a typical Scanning electron microscope. (image from: http://saturno.fmc.uam.es/ web/superficies /instrumentacion /Instrumentation.htm) The scheme of a typical scanning electron microscope is shown in Figure II-18. Actually, it resembles a conventional optical microscope, but in optical microscopes light is confined and focused by optical lenses, while in SEM electrons are instead confined by a magnetic field created by surrounding coils. There are different types of electron beam generation,

Figure II- 19 . 21 II. 3 . 2 . 2

 1921322 Figure II-19. Schematic presentation of the EBSD configuration and example of a Cu diffraction pattern. 21II.3.2.2 Atomic Force Microscopy (AFM)

Figure II- 20 .

 20 Figure II-20. Illustration of the AFM working principle. 22 In addition to the normal topographical information acquisition, the conductivity mapping of the sample surface was also carried out in this work. Samples were deposited on conductive ITO/glass substrate. To ensure a good connection, a 100 nm thick Au contact was evaporated onto the ITO, as shown in Figure II-21. During the characterization, a conductive Pt coated Si tip was used under contact mode. Meanwhile, a potential was applied

Figure II- 21 .

 21 Figure II-21. Illustration of electrical Atomic Force Microscope characterization.

  characterization, there are two operational modes commonly used, namely imaging and diffraction mode. Their schematic configuration is shown in Figure II-22. A JEOL-2016 LaB6 working at 200 kV was used in this work.

Figure II- 22 .

 22 FigureII-22. Schemes of imaging and diffraction mode in TEM.23

Figure II- 23 .

 23 Figure II-23. Optical measurement of a) the total transmittance (TT) and b) the total reflectance (TR).

Figure II- 23

 23 Figure II-23 presents the configurations used for the measurements in our work, figure a) total transmittance and figure b) total reflectance. Absorption can as well be measured in two different ways. The first one involves placing the sample inside the ARTA. In this method, the measurement is more accurate but extra care has to be taken not to damage the spectroscope components. The second method to obtain the absorption Ta (used in this work) is by deducting it from Tt and Tr according to Equation II 9.

  allowed transition, n=2, direct forbidden transition, n=3/2, indirect allowed transition, n=1/2, indirect forbidden transition, n=3.II.3.4 Electrical characterization II.3.4.1 Four probe measurementsFour point probe is a widely used technique to rapidly measuring the resistivity of semiconducting materials. The basic setup is illustrated in Figure II-24. By passing a current through the two outer probes(1, 4), voltage is measured between the two inner probes(2,

Figure II- 24 .

 24 Figure II-24. Schematic illustration of four point probe measurement.

  FigureII-25. To measure the carrier concentration and mobility, the semiconductor sample is exposed to a magnetic field in the Z axis and a current is applied between the corners 1 and 3, namely I13. Due to the Lorenz force applied on different charged particles (electrons and holes), a potential difference can be created from the corner 2 and 4, namely V24.

Figure II- 25 .

 25 Figure II-25. Illustration of the Van Der Pauw configuration for Hall effect measurement. 32 By also measuring I24 and V13, an average value of carrier concentration and mobility can be obtained. Measurements in this study were conducted with a Nano Metrics model at room temperature by applying 0.5 T of magnetic field in such Van Der Pauw configuration. Due

Chapter 3 .

 3 Thin film deposition of silver and silver oxide -71 -

Figure III- 1 .

 1 Figure III-1. Reflux system used for the synthesis of Ag metalorganic precursors.

  Through a collaboration with Dr.Patrizia Rossi's group from University of Florence, Italy, the molecular and crystal structures of AgP and AgT were studied by means of single crystal X-ray diffraction. Measurements were carried out with an Oxford Diffraction Excalibur diffractometer using the Cu-Kα radiation (λ = 1.54184 Å) for AgP and Mo-Kα radiation (λ = 0.71073 Å) for AgT. Data collection and data reductions were performed with the program CrysAlis Pro. Finally, absorption correction was performed with the program ABSPACK.29

2 .

 2 Single crystals of AgT were twinned. Two domains were taken into account during refinement. The fluorine atoms bound to the carbon atom C5 (see Figure III-2 below) of the -hfac anion in AgP are disordered. Such disorder was modelled by introducing three positions for each fluorine atom (refined occupancy factors: 0.44, 0.41, 0.15). In AgT all the CF3 moieties are disordered (see Figure III-3 below). Such a disorder was modelled by introducing two positions for each fluorine atom bonded to C4, C5 and C10 (with refinedoccupancy factors: 0.62/0.38 for the atoms bonded to C4; 0.83/0.17 for the atoms bonded to C5; 0.76/0.34 for the atoms bonded to C10). In addition all the atoms of the C4(F3) moiety were set in double position (refined occupancy factors:0.55/0.45). For all the non-hydrogen atoms in AgP and AgT, anisotropic thermal parameters were used. For the hydrogen atoms in AgP, they were found in the Fourier difference map and their positions were freely refined, while their thermal parameters were refined accordingly to the bound atoms. The hydrogen atoms in AgT were introduced in calculated position and refined in agreement with the carbon atom to which they are bound. Chapter 3. Thin film deposition of silver and silver oxide -73 -

  )C(1)C(3)O(2), centroid of the central ring of phen as defined by C(9)C(10)C(11)C(12)C(16)C(17), is 3.60(1)Å) and the angle between the CT-CT line and the mean phen plane is 71.2(2)°. Finally, the dimers stack along the a axis in a head-to-tail arrangement (Figure III-2 b), which brings the carbon atom C(10) of a symmetry related (-x, -y, 2-z) phen molecule 3.350(8) Å apart from the silver ion, with a distance less than the sum of the Van der Waals radii (3.81 Å) 42 . In addition, inter-dimer π stacking interactions involving the phen molecules provided by two contiguous dimers are present (the mean phen planes are 3.5 Å apart). Finally, dimers propagate along the b axis (Figure III-2 c) through a weak C(7)H(7)...O(1)' intermolecular interaction (' =x,1+y,z; D...A and H...A distances 3.628(6) and 2.75(6) Å, respectively, DHA angle 161(6)°). The check CIF reports of both precursors submitted to Cambridge structure data center are presented in the Appendix at the end of the chapter.

Chapter 3 .

 3 Thin film deposition of silver and silver oxide -76 -

Figure III- 2

 2 Figure III-2. a) Ellipsoid representation of the dimer found in the solid state structure of AgP (hydrogen atoms have been omitted for sake of clarity, for the same reason just one model of the disordered CF3 moiety was shown). b) along the b axis direction showing the intra-and inter-dimer π-stacking interactions. Hydrogen atoms were omitted (top); c) along the a axis direction showing the CH...O intermolecular interactions. For the sake of clarity just one model of the -hfac fluorine atoms bound to C5. The solid state structure of compound AgT was also determined by means of single crystal X-ray diffraction. In the asymmetric unit of AgT one [Ag(triglyme)2] + cation and one [Ag(hfac)2] -anion are present (Figure III-3). In the two ions, the silver atom shows different coordination spheres. In fact, in the [Ag(triglyme)2] + cation, the silver atom Ag(2) is

  [[Ag(hmten)][Ag(hfac)2]; hmten = hexamethyl-triethylene tetra-amine]. 44 The two -hfac anions lay in two planes nearly perpendicular to each other (the angle between the two mean planes defined by O(1), O(2), C(1), C(2), C(3) and O(3), O(4), C(6), C(7), C(8) is 84.8(2)°).Finally, no relevant intermolecular interactions are present in the crystal packing of AgT.

Figure III- 3 .

 3 Figure III-3. View of the [Ag(triglyme)2]+ cation and [Ag(hfac)2]-anion in AgT (ellipsoids are drawn at 30% probability). For the sake of clarity, just one model of all the -hfac fluorine atoms (which were disordered and refined isotropically) is shown. For the same reason the hydrogen atoms are not shown in the image.

Figure III- 4 .

 4 Figure III-4. FT-IR spectra of AgP a) and AgT b) in nujol (peaks from nujol marked with arrows).

1 H and 13 C

 13 NMR of AgP and AgT were recorded in CDCl3 (TableIII-4). The 1 H NMR spectrum of AgP (Figure III-5) has been attributed in accordance to the assignments of references.45,46 The spectrum shows a singlet (5.79 ) associated with the ring proton of the hfa ligand, while the phenanthroline shows peaks in the typical range of aromatic ligand from 7.6 to 9.1 . The a protons are observed as a singlet at 9.04 , the b proton as a multiplet at 7.79  , the c protons as a doublet of doublets at 8.40  and, finally, the e protons as a doublet at 7.89 . The phen peaks have been assigned following the notation reported as a note in TableIII-4. The 1:1 (hfa:phen) ratio found through evaluation of intensities of pertinent resonances finds counterpart in the ratio obtained through single crystal X-ray diffraction, as shown in Figure III-5.

Figure III- 5 .

 5 Figure III-5. 1H NMR spectrum of AgP in CDCl3.

Figure III- 6 .

 6 Figure III-6. 13C NMR spectrum of AgP in CDCl3. In Figure III-6, the 13 C NMR spectrum of AgP shows resonances associated with the coordinated hfa ligand, which consist of a singlet ( 86) for the CH group, quartets ( 118)

  In Figure III-7, the spectrum shows a singlet at 5.81  associated with the ring proton of the hfa ligand. Different resonances are associated with the triglyme framework. The methyl groups (a protons) are always observed as a singlet at  = 3.60. The b and c protons are observed as a multiplet centered at  = 3.62. The protons d are associated with a singlet at  = 3.71. Evaluation of intensities of pertinent resonances points to a 1:1 (hfa: triglyme) ratio in accordance with single crystal X-ray diffraction.

Figure III- 7 .

 7 Figure III-7. 1H NMR spectrum of AgT in CDCl3. In Figure III-8, the 13 C NMR spectrum of AgT is similar to that found for the AgP adduct in regard to the hfa ligand. Resonances associated with the coordinated hfa ligand consist of a singlet ( 86) for the CH group, a quartet ( 117) for the CF3 groups and a quartet ( 177) for the CO groups. The quartets are due to first order (CF3; 1 J = 289 Hz) and second order (CO; 2 J = 32 Hz) coupling with the CF3 fluorine atoms. The coordinated triglyme signals are observed at  61.21 (s, OCH3, a),  70.11 (s, OCH2, b),  70.58 (s, OCH2, c) and  71.42 (s, OCH2, d).

Figure III- 8 .

 8 Figure III-8. 13C NMR spectrum of AgT in CDCl3. The thermal characteristics of the present adducts were evaluated by Thermogravimetric Analysis (TGA). The TGA profile of AgP (Figure III-9 a) presents a two-step weight loss (as clearly shown in the derivative curve), in the range 190-230 °C and 230 -280 °C, respectively, with a residue of 37% up to 400 °C. This value does not match the stoichiometric Ag weight % in the precursor, namely, 21.8%. Thus it can be concluded that the precursor does not present a clean decomposition, with some organic species still remaining at 400 °C. Conversely, the TGA curve related to AgT (Figure III-9 b) shows a single (see derivative) rapid weight loss in the range 150 -250 °C, yielding a 22.5% residue when heating to 400 °C. The stoichiometric Ag weight portion in the precursor being 21.9%, the value obtained from TGA is thus well tuned with the expected percentage of metallic silver in the compound, implying a clean and complete decomposition of the precursor. By comparing the TGA results for those two precursors, AgT appears as a better silver metalorganic precursor.

Figure III- 9 .

 9 Figure III-9. Thermogravimetric Analysis (TGA) and its derivative curves (DTG) of precursors, between 30 °C and 400 °C, of a) AgP and b) AgT. TGA was conducted under N2 atmosphere with heating rate 5 °C/min. III.2 Deposition of silver coatings via MOCVD .

  5 have been tested. Only AgAc, AgP and AgT yielded Ag coatings. In this section, the results of Ag coatings deposited with the three types of Ag precursors are given.

Figure III- 10 .Chapter 3 .

 103 Figure III-10. TGA and DGA of AgAc in a) nitrogen and b) air atmosphere with heating speed 20 °C/min. 47 Among all the Ag precursors, AgAc is reported for the first time in this work as a metal organic precursor for MOCVD deposition of Ag coatings. The thermal properties of AgAc have been investigated in the work of BM. Abu-Zied et al. 47 As shown in Figure III-10,AgAc decomposes rapidly at about 295 °C in both nitrogen and air atmosphere with a heating speed of 20 °C/min. In order to dissolve AgAc in alcohol solvent such as ethanol, two dispersants can be used, namely, ethylenediamine and ethanolamine (as introduced in the Chapter II).48,49 We used solutions prepared with each of the dispersants. While no Ag coatings can be obtained with ethylenediamine, solutions containing ethanolamine yielded Ag coatings using similar conditions 320 °C). The hypothetic mechanism behind this different behavior is shown in FigureIII-11. As it's well known, for ammine complex there are normally 2 coordination numbers for the silver ion and they present a linear coordination (angle between N-Ag-N is 180 °).50 Meanwhile, the molecule structures of the two dispersants are rather similar, except that in ethylenediamine, two ammine groups are presented at each end of the molecule while in ethanolamine, only one ammine group is present on one side of the molecule, as shown below:

Figure III- 11 .

 11 Figure III-11. Illustration of silver precursor complex group formed with a) ethanolamine, forming single silver precursor molecule containing structure, and b) with ethylenediamine, creating linear chain like complex with multiple silver precursor molecules included. "L" represents ligand in precursor. The depositions of Ag coating using AgAc as precursor were carried out at a pressure of 0.1 mbar and at 330 °C on cleaned glass substrates. With ethanolamine as the dispersant, AgAc is able to give Ag coatings with 500 injected droplets (about 12 ml) and the SEM images of the Ag coatings with different flow rates are shown in Figure III-12. The circular shape of depositions is typically the sign of incomplete vaporization of the precursor or the precursor solution, which could be due to the low volatility of the AgAc precursor. With less carrier gas flow rate during the deposition, the size of grains in Ag coatings tends to decrease from

Figure III- 12 .

 12 Figure III-12. SEM images of Ag coatings deposited in PI-MOCVD at 330 °C for 5 minutes, using AgAc as precursor concentration 0.02 M 12 ml, with ethanolamine as dispersant (about 0.05 M ). Coatings were deposited on glass substrate with different flow rate of carrier gas: a) 2 L/min, b) 1 L/min and 0.5 L/min. With longer deposition duration, more continuous Ag films with 100 nm thickness could be deposited as shown in Figure III-13. These films were deposited with 0.02 M AgAc precursor solution, also at 330 °C on glass substrate. Figure a) and b) are the silver coating deposited for 8 minutes duration (20 ml solution consumption), presenting a yellow brown color and connected Ag grain networks. In Figure c) and d), the Ag film deposited for 10 minutes is presented (with 25 ml solution consumption). With a longer deposition duration, the silver grains became denser and more continuous on the substrate, and the films presented a silver white color was. The resistivity of this silver film reached down to 40 μΩ.cm.

Figure III- 13 .

 13 Figure III-13. Conductive Ag coatings using 0.02 M AgAc precursor in PI-MOCVD system. a), b) with 20 ml deposition duration about 8 minutes and c), d) 10 minutes deposition with 25 ml solution consumed. Image b) is the photo corresponding to coating a) and image d) is the photo of the coating in c).

3 . 1 .

 31 Figure III-14. Similar effects of enhanced deposition by ethanolamine is shown here with precursor AgP, in which a) shows a pink Ag coating with small grains (size about 20 nm) without ethanolamine as the dispersant. In figure b) a darker Ag coating deposited with ethanolamine as dispersant is shown, presenting higher deposition rate and bigger grain size (averagely about 55 nm). Using AgT as the precursor with the same deposition conditions, thanks to a much better solubility of AgT in alcohol solvents, no dispersant was necessary. The deposited Ag coating is presented in figure c), which shows a blue color coating with big round silver grains on top (maximum up to 100 nm diameter). The morphology is rather different from the silver coatings shown in figure b).

Figure III- 14 .

 14 Figure III-14. SEM images of Ag coatings deposited with AgP precursor: a) without ethanolamine and b) with ethanolamine as dispersant. c) SEM image of the coating obtained with AgT (no need to add dispersant) in PI-MOCVD system at 320 °C with 1000 injection pulses (0.01 M about 12 ml solution consumption).

Chapter 3 .

 3 Thin film deposition of silver and silver oxide -88 -clusters, as shown in Figure III-15 a) and b). Clusters with diameter of about 6 μm can be seen. EDS in figure c) confirmed the existence of Ag.Trying to get insight into the mechanism behind the different results obtained at different temperatures, we made a test with precursor solution sprayed on a hot plate at 300 °C using a sprayer. The spray droplets moved rapidly upon arriving in the hot plate, leaving practically no decomposition products on the heated glass substrate. Combined with the fact that MOCVD is a very surface sensitive technique, the absence of deposition might be due to a low adhesion between the glass substrate and the Ag particles produced, which could be carried away by the relatively high gas flow before having the chance to grow no the substrate surface. It has to be noted that the PI-MOCVD reactor had the initial configuration (see chapter II) for this tests, and thus the fact that the substrate holder was upside down could imposed an extra difficulty.

Figure III- 15 .

 15 Figure III-15. Ag clusters deposited by AA-MOCVD using AgAc as precursor with concentration of 0.01 M (and including ethanolamine with a concentration of 0.02 M) at 210 °C for 1 h. a) and b) are the SEM images of the silver clusters. Inset in a) is a picture of the glass substrate. c) is the EDS for Ag element identification (the other elements are from the glass substrate).

  of depositions using the different precursors, AgT showed higher deposition efficiency than AgP (see SEM images in Figure III-19). For a deposition duration of 0.5 h, AgP yielded a yellow coating while AgT gave an orange color with a higher particle density (Figure III-19 a) and c)).

Figure III- 16 .

 16 Figure III-16. Images of silver coatings: a) P1 (yellow), b) P2 (dark in the picture but with a purple hue to the naked eye), c) T1 (orange), d) T2 (as for P2, dark in the image but having a purple hue to the naked eye).Increasing the deposition duration to 1.5 h, coatings from AgP and AgT both presented purple color but the grain density for the latter was still higher than for the former (Figure III-19 b) and d); Pictures of the films are shown in Figure III-16). The color of the films is due to both thickness, but also to particle diameter and shape, since Ag nanoparticles present plasmonic absorption at different wavelengths depending on both the size and the shape of the particles (Transmission spectra of the films are shown in Figure III-17).52,53

Figure III- 17 .

 17 Figure III-17. Transmittance spectra of Ag films deposited with different precursors AgP and AgT with deposition duration 0.5 h (P1 and T1) and 1.5 h (P2 and T2).

For

  the deposition conditions used, the particles in the coating are not connected and thus the samples present high resistance values, namely a sheet resistance of 5 ×10 8 Ω/cm 2 measured by 4 probes. As well, the adhesion between the Ag coating and glass substrate is poor and the films can be easily removed. The thickness of the coatings is on average equivalent to the diameter of the nanoparticles (see Figure III-18).

Figure III- 18 .

 18 Figure III-18. SEM Crossection of sample T2 (deposited using AgT precursor for 1.5 h). a) Secondary electrons b) Back/scattered electrons. The height of the coating is about 35 nm.

Figure III- 19 .

 19 Figure III-19. SEM images of silver coatings deposited at 320 °C by AA-MOCVD using: AgP precursor with deposition duration a) 0.5 h (P1); b) 1.5 hrs (P2); and AgT precursor with deposition duration c) 0.5 h (T1); d) 1.5 hrs (T2). By analyzing the different particle diameters, a narrower size distribution of Ag grains in P1 is revealed from comparing Figure III-20 a) and b). The size of grains in P1 is mainly located around 19 nm, while particles obtained in T1 range broadly distributed from 10 nm to 22 nm.

Figure III- 20 .

 20 Figure III-20. Size distribution of Ag nanoparticles deposited from AgP and AgT: P1 (a) and T1 (b). As shown in Figure III-21 a) and b), the coatings obtained with both precursors present well crystallized Ag nanoparticles, but from the sample deposited with AgP (P2), a thin shell surrounding the Ag nanoparticles was observed. Also, as already hinted by TEM images (Figure III-21 c) and d)), Ag nanoparticles obtained from AgP gave more regular round shapes than those obtained from AgT (T2), which are elongated. Electron diffraction patterns of the different particles show diffraction rings which only correspond to metallic Ag. Thus a smaller average grain size for Ag coatings with AgP precursor can be attribute to such a shell layer, which possibly limited the grain growth during the silver deposition process.

Figure III- 21 .

 21 Figure III-21. Transmission Electron Microscopy images and electron diffractions for Ag particles deposited using AgP (P2), a), c), e) and AgT (T2), b), d), f).To further evaluate the purity and crystallinity of the coatings obtained from the two precursors, Raman spectroscopy and X-Ray Diffraction were used and the results are shown in Figure III-22. Figurea) shows the Raman spectrum obtained for silver particles deposited using AgP. In this case, strong peaks related to the C-OH group in the range 1200 cm -1 and 1600 cm -1 can be observed, in agreement with the presence of the carbonaceous shell observed by TEM. On the other hand, coatings deposited with AgT showed a flat band curve in this high wavenumber range.

Figure III- 22 .

 22 Figure III-22. Characterization of silver coatings deposited with the different precursors. a) RAMAN spectroscopy for carbon residues detection; b) Normalized GIXRD patterns of silver coatings deposited on glass substrate. The carbon contaminations found for AgP, both by Raman and TEM, are due to the incomplete decomposition of the precursor under the deposition conditions used, and are in agreement with the TGA results shown above. Removal of the organic coating could not be obtained by annealing. For temperatures above 300 °C, the Ag nanoparticles in the coating tent to fuse together forming bigger rounded particles (See the image in Figure III-23).

Figure III- 23

 23 Figure III-23. a) SEM image of the as deposited Ag coating (using AgP). Images b) and c) show SEM images of the coating after annealing at 400 °C in air for 1 h.

  cyclic voltammetry was used to evaluate the oxidation potentials of the different oxidation reactions at a certain condition. In this case, silver films (about 100 nm thick) were deposited on corning glass by PI-MOCVD at 330 °C for 10 minutes using 0.02 M AgAc as precursor solution. As it's shown in Figure III-24, a voltammogram was acquired by applying the potential sequence 0 V → 0.7 V → -0.3 V → 0 V (all the potential values in the chapter are given versus Ag/AgCl 3.5 M KCl reference electrode) and a scanning speed of 10 mV/s.

Figure III- 24 .

 24 Figure III-24. Cyclic voltammetry of a conductive Ag film (100 nm thick)/ Corning glass, voltage range: 0 V → 0.7 V → -0.3 V → 0 V, in 1 M NaOH solution with scanning speed 10 mV/s.The silver films deposited at the same condition (as previously mentioned) were oxidized in 1 M NaOH electrolyte, by applying 0.7 V potential with different durations, as shown in Table III-7. Given the high oxidation state of Ag in AgCuO2,58 this potential value was chosen to ensure high oxidation states for Ag. The same Ag film was cut into several pieces with roughly similar area (about 0.25 cm 2 ), thus those Ag thin film samples were roughly identical.

Figure III- 25 .

 25 Figure III-25. Figure a) shows the original as deposited Ag film (with a low resistivity of about 40 μΩ.cm). With electrochemical oxidation for 500 s, the morphology of the oxidized sample shows particle like and discontinuous surface in Figure III-25 b). With longer

Figure III- 25 .

 25 Figure III-25. Scanning Electron Images of Ag/ Ag oxides film morphology evolution with oxidation duration. A). Original Ag film; b). after 500 s (E1); c). after 1000 s (E2); d). after 1300 s (E3).

Figure III- 26 .

 26 Figure III-26. GIXRD patterns of Ag film oxidized electrochemically in 1 M NaOH solution with 0.7 V potential showing the phase evolution with time. Dark curve represents Ag film as deposited in PI-MOCVD and is oxidized with duration 500 s (blue), 1000 s (pink) and 1300 s (green).

Chapter 3 .

 3 Thin film deposition of silver and silver oxide -98 -iCAP 7400, in collaboration with Prof. Christophe Vallé e, from the Laboratoire des technologies de la Microé lectronique (LTM)). Air was primarily tested as the gas source at maximum plasma power (500 W) and the emitted plasma spectra is shown in Figure III-27 a). Surprisingly enough, with air as the inlet gas, a very weak trace of ozone signal was detected (at 774.2 nm) and the emitted spectra peaks are similar to those obtained when using pure N2 as the gas source, as shown in Figure III-27 b). This absence of ozone generation with air is further studied by varying the ratio of oxygen mixed with Ar. Generation of ozone presents a quasi linear dependence with the concentration of oxygen in inlet gas, as illustrated in Figure III-27 c). Lastly, to understand the relationship between plasma power and ozone generation with pure oxygen as gas source, different powers were analyzed and the result is shown in Figure III-27 d).The figure shows that a minimum power of 300 W is needed to be able to generate measurable ozone species intensity. With higher powers, the intensity of the ozone signal increases.

Figure III- 27 .

 27 Figure III-27. Atmospheric plasma analyzed by ICP-OES with different gas source and voltages. a) 500 V with air inlet gas. b) 500 V with N2 inlet gas. c) O2 as inlet gas with voltage from 200 V to 500 V. d) at 500 V, different concentration of O2 combining with Ar.

Figure III- 28 .

 28 Figure III-28. SEM images of Ag films a) before and b) after oxygen plasma oxidation for 3 min with 300 W at 70 °C.

Figure III- 29 .

 29 Figure III-29. Phase evolution of Ag films oxidized by oxygen plasma under power 300 W and 70 °C with different durations. Dark line represents as deposited Ag film, red represent Ag film under plasma 1 min and blue refers to Ag film under oxidation for 3 min. All samples used are cut from the same Ag film.The 100 nm thick silver film was deposited on corning glass via PI-MOCVD at 330 °C with AgAc as precursor, with concentration 0.02 M and 25 ml volume consumption. To avoid the

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -113 -discussed. By adjusting the different deposition parameter, Cu2O films with controlled texture have been achieved. High quality textured Cu2O films have been characterized and implemented in diodes, obtaining on-off ratios exceeding 10 4 . To further study the electrical properties of the of (111) textured Cu2O films, they were characterized by Conductivity

  Figure II-3), with precursor flow coming from the bottom to the top and reacting on the hot plate above (which faces down). While conducting the initial depositions with such configuration, it was very hard to obtain good homogeneity of the deposited Cu2O films. As it's well-known, in MOCVD, a stable laminar precursor flow is crucial to obtain homogeneous depositions. Under such configuration, the precursor flow needs to go from bottom to the top, against gravity. Even though the large particles and droplets can be avoided in this case, it's rather challenging to obtain a homogeneous flow, as shown in Figure IV-1 a).

Figure IV- 1 .

 1 Figure IV-1. Depositions of Cu2O via AA-MOCVD using our homemade system at 330 °C for 1 hour a) in its initial configuration, roughly 60 nm thick; b) in the optimized configuration with thickness about 100 nm.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD-114 -our AA-MOCVD system, it was rather difficult to obtain optimized and reproducible samples.

Figure IV- 2 .

 2 Figure IV-2. Different configurations in the process of AA-MOCVD optimizationTo avoid some of these issues, the first modification was to reverse the orientation of the deposition chamber, turning it by 180°. In this configuration, the entrance of the flow into the deposition chamber is located at the top of the chamber, the heating plate staying at the bottom, facing up. The ventilation outlet was also placed at the bottom, to the right, below the hot plate (Figure IV-2 a).

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -115 -As it is shown in FigureIV-2 a) and b), initially there was a part of the system in which the flow was horizontal before entering the chamber. When tests were performed with such configuration, due to the gravity, a separation between the precursor and the carrier gas was observed (i.e. a layer of precursor solution formed on the horizontal tube). Thus when the flow entered the deposition chamber, it wasn't laminar, and it even carried big droplets from the condensation on the glass tube. Thus, in order to avoid this carrier gas and flow separation, this section of glass tubing was titled from 0° to approximately 45° with respect to the horizontal (FigureIV-2 b). With such a tilt, it also largely reduced the chance for the condensed solution droplets in the tube to end falling into the deposition chamber. The tilt also enabled a smooth and homogeneous introduction of aerosol droplets into the deposition chamber, while avoiding bigger size droplets that form following condensation in the tube walls.Lastly, the location of the carrier gas #2 (CG#2) inlet was also carefully optimized: at the beginning of the tube (Figure IV-2 a), in the middle of the tube (Figure IV-2 b) and at the end of the tube (Figure IV-2 c). Placing it at a bigger distance away from the CG#1 inlet, the bigger droplets in the flow have more time to condense and return back to the solution chamber instead of reaching the hot substrate. In addition, it's more safe to separate the oxidizing gas far from the aerosol chamber, where hot flammable solvent is present.

Figure IV- 3 .

 3 Figure IV-3. Details of the final optimized AA-MOCVD configuration.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -116 -gas mixer was designed fulfill two functions: introducing CG#2 and collecting condensed droplets. As shown in Figure II-4, a square Si substrate was cut to precisely fit into the bigger round gas tube, just beneath the arrival of the flow through the main thinner tube. As the solution mist comes down from the top, the heavier, faster droplets impacted on the Si substrate, falling along its edges to be finally collected at the bottom of the large tube. CG#2 was introduced from the bottom and well mixed inside the bigger gas tube with the solution mix before reaching the deposition chamber through the thinner tube at the bottom (see Figure II-4 b). Thus with such gas mixer, hole free samples with better homogeneity were obtained, with good reproducibility. As a result, nice and homogeneous Cu2O films could be deposited as shown in Figure IV-1 b.

Figure IV- 4 .

 4 Figure IV-4. Cu2O films deposited at 320 °C a) and 370 °C b) with Cu(AcAc)2 0.01 M concentration for 1 hour in AA-MOCVD. (The black dots on the Cu2O films are carbon tape residues) With different oxygen ratio during deposition at 350 °C, the morphology of the samples was rather smooth and similar, as presented in Figure IV-5. The deposited Cu2O films had thickness around 100 nm. The grain size was very small (about 20 nm) but with dense structure. Samples deposited at 375 °C looked very similar to samples deposited at 320 °C, as shown in Figure IV-4 a) and b). With Raman Spectroscopy, the deposition quality of all samples was further examined, and the results are shown in Figure IV-6.

Figure IV- 5 .

 5 Figure IV-5. SEM images of Cu2O samples deposited with various oxygen ratios (total flow 10 L/min made of oxygen and nitrogen) at 350 °C. a) 25% of O2. b) 40% of O2. c) 80 % of O2.

Figure IV- 6 .

 6 Figure IV-6. Raman spectroscopy curves of a) Cu2O samples deposited at 350 °C and b) deposited at 375 °C for 1 hour, with oxygen ratio ranging from 25% to 80% using Cu(AcAc)2 as precursor.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -118 -

Figure IV- 7

 7 Figure IV-7 a). As shown in the inserted photo, a white layer was observed on the glass substrate which might correspond to the several new emerged peaks. When the deposition

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -119 -

Figure IV- 7 .

 7 Figure IV-7. Raman spectroscopy curves of a) Frost microscope glass slide (Inset: optical image of the substrate before and after plasma treatment) and b) corning glass slide before and after 30 min 500 W oxygen plasma treatment.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -120 -

Figure IV- 8 .IV. 1 . 3 . 1

 8131 Figure IV-8. Raman spectra of copper oxide films deposited with 0.6 (black) and 0.3 (red) oxygen ratio as gas source (total flow 15 L/min with N2 and O2). Precursor Cu(AcAc)2 concentration 0.05 M with 1 hour deposition at 300 °C.IV.1.3Depositions of Cu2O films using CuF as precursor

Figure IV- 9 .

 9 Figure IV-9. Cu2O coatings using CuF as precursor at 330 °C deposited on a) Frost glass and b) Corning glass. c) Optimized Cu2O film on corning glass.

Figure

  Figure IV-10 presents the SEM images and Raman spectra of coatings obtained with different oxygen ratio. For 12.5% O2 ratio (D1), a dense film with smooth surface and small sized grains was obtained (Figure IV-10 a)). By slightly rising the oxygen ratio in the flow to 13.75% (D2), the surface of the deposited films was rather similar, being still dense and continuous, as shown in Figure IV-10 b). When the oxygen ratio was further increased to 15% (D3), Figure IV-10 c), the films presented a non-continuous, rough surface with clear cubic Cu2O grains. All those cubic like grains were stacked together with only the corner of the cube pointing above. Those films were characterized by Raman spectroscopy which confirmed the only existence of Cu2O, as shown in Figure IV-10 d) below. As shown, the

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -122 -orientation of the polycrystalline grains in the films. Even though the XRD characterizations

Figure IV- 10 .

 10 Figure IV-10. SEM of Cu2O samples deposited at 335 °C, for 1 hour with different oxygen ratio: a). 12.5% O2; b). 13.75% O2; c). 15% O2. With increasing of oxygen ratio during deposition, more cubes of Cu2O were formed. d) Raman spectra for Cu2O films deposited with different oxygen ratio. Comparing with the Raman curves from Cu2O samples deposited with Cu(AcAc)2 precursor (in Figure IV-6), the Raman spectra from Cu2O depositions with CuF precursor did not show a measurable quantity of carbon contamination. The resistivity of the different films was measured using four point probe and all gave values between 215 and 250 Ω.cm, as indicated in TableIV-2. In addition, as shown in the same table, the deposition rate has the trend of enhancing with higher oxygen ratio during deposition, even though the thickness for FigureIV-10 c) is a range between 120 and 197 nm due to the discontinuity and roughness of the film, which is in accordance with the work of G.G. Condorelli et al.4 To summarize, oxygen

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -125 -

Figure IV- 11 .

 11 Figure IV-11. Cu2O film deposited with CuF 10 mM concentration a) without Ag precursor(L0) at 335 °C for 3.5 hours and samples with various AgT precursor concentration added: b) 10 mM CuF: 1 mM AgT (L3); c) and d) 10 mM CuF: 1.25 mM AgT (L4), the inserted image is the backscattering electron image; e) and f) 10 mM CuF: 2.5 mM AgT (L5), ), the inserted image is the backscattering electron image. As presented in Figure IV-11, the morphology of Cu2O film deposited without adding Ag precursor is shown in figure a), with a rather smooth and dense surface. In figure b), itpresents the morphology of Cu2O film with small quantity of AgT precursor added (sample L3) but the silver content is below the detection limit of EDS. The morphology of samples L1 and L2 is rather similar with L3, in which more clear cubic shape of the Cu2O grains emerged on the surface. Figurec) and d) show the SEM images of coatings obtained with further increased silver precursor mixed into the deposition solution (detectable Ag atomic ratio 0.3% to Cu). In figure c), a much rougher surface is presented with visible bigger grains.In addition, several nanoparticles scattered across the sample surface and locate at the lower points between the bigger grains. With the inserted backscattering electron image at high magnification in figure d), the nanoparticle was identified as a heavier element than copper, which is most likely to be silver.With the highest concentration of silver precursor added, the surface of the thin film is also full of scattered silver nanoparticles between the grains, as shown in figure e). It is more clear in the backscattering electron image in the inset. The EDS showed Ag ratio of 1% in the film. With higher magnification image in figure f), cubic shaped grains with sizes of about 200-300 nm and clear grain boundaries are presented.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -126 -Thin film thickness was obtained from SEM cross-section measurement, and is detailed in Table IV-3. It is clear that the addition of Ag precursor yielded higher deposition rates. By calculating the average deposition rate as a function of the concentration of Ag precursor (see Figure IV-12), we observed that the deposition rate was enhanced from 2.1 nm/min to 3.5 nm/min.

Figure IV- 12 .

 12 Figure IV-12. The average deposition rate with different silver quantity incorporated. To check the crystallinity evolution of Cu2O with the amount of silver, Raman Spectroscopy characterization was conducted. As presented in Figure IV-13 a), Raman curves show that all films deposited with Ag incorporation are still composed of Cu2O phase without any trace of CuO. Interestingly, with more Ag used, a change in peak intensity has been found for the strongest peak 165 cm -1 and the second strongest one at 210 cm -1 , which can be attributed to a change in grain orientation or an increase in film crystallinity. In Figure IV-13 b), the XRD patterns of these samples also confirmed the orientation shift.Without adding Ag precursor (sample L0), the main orientations obtained were assigned to (1,1,0) and (1,1,1). While as more silver precursor were mixed into the deposition solution, thus more Ag particles were incorporated, the intensity of the reflection corresponding to (1,1,0) decreased while the reflection corresponding to (1,1,1) increased, especially obvious in sample L5.

Figure IV- 13 .

 13 Figure IV-13. Cu2O films deposited with various quantity of silver precursor mixed: a) Raman spectra, and b) normalized XRD patterns.

  morphology, crystalline structure and electronic properties. During the deposition, the solution consumption rate was controlled by the CG1 flow (2 L/min N2) and was maintained around 1.5 ml/min. CG2 was then introduced as oxidizing gas (5 L/min air + 3 L/min N2) through the gas mixer mentioned in Chapter II (FigureII-4). The deposition chamber was slightly below atmospheric pressure due to the ventilation with the liquid pressure gauge reaching about 3 cm. The substrates were heated at temperatures ranging between 290 and 365 °C. Glass (Corning) substrates were used for the depositions. Standard deposition duration was 220 min. Shorter deposition times (< 60 min) were also used to evaluate the initial nucleation mode.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -128 -To introduce the humidity into the carrier gas, a bubbler containing deionized water at room temperature was added to the CG2 line, as shown in Figure IV-14Error! Reference source not found.. The CG2 flow was thus divided in two separate lines, including i) CG2a gas going directly to the reactor; and ii) CG2b bubbling through water (50 ml) during deposition process. The glass bubbler is a glass bottle with inner diameter about 6 cm. Both flow speed were adjusted to maintain a total CG2 flow rate of 8 L/min with the certain gas composition.By increasing the flow rate through CG2b (i.e. by increasing the water consumption rate) we could control the humidity (water molar ratio) in the carrier gas. Thus, water consumption and the water molar ratio in the flow during deposition was controlled and measured. The study was based on the samples deposited under different conditions, as shown in Table IV-4, with various deposition temperatures and water molar ratios.

Figure IV- 14 .

 14 Figure IV-14. Scheme of the modified AA-MOCVD system allowing the introduction of humidity in the carrier gas flow.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -129 -

  Figure IV-15. At lower deposition temperatures, namely 290 °C, and both in dry and wet conditions, discontinuous films were obtained in which cubic crystals were found scattered on the substrate along with a thin Cu2O layer below, as shown in Figure IV-15 a) and e). The only difference was that in wet deposition condition, the size of cubic grains were much bigger, with edge lengths reaching about 300 nm, as compared with 50 nm for grains in dry deposition condition.

Figure IV- 15 .

 15 Figure IV-15. Scanning Electron Microscopy images of films deposited at different temperatures in dry and wet conditions. a) 290 °C (D290);b) 305 °C (D305); c) 350 °C (D350); d) 365 °C (D365) and with humid carrier gas with temperatures: e) 290 °C (W290, titled surface); f) 305 °C (W305); g) 350 °C (W350); h) 365 °C (W365). In the deposition series with dry CG at temperature ranging from 305 °C to 365 °C, the morphology of the Cu2O films is shown from Figure IV-15 b) to d). Upon increasing the deposition temperature above 290 °C, continuous films were formed at 305 °C and the cubic feature of Cu2O was not obvious above the surface, shown in Figure IV-15 b). As deposition

Figure IV- 16 .

 16 Figure IV-16. XRD patterns of samples deposited at different temperatures with a) dry CG and b)humid CG; c). Evolution of texture coefficient with deposition temperature for both dry and wet CG; d). Raman spectra of samples deposited with dry and humid CG.To further characterize thin film quality and phase composition, Raman spectroscopy and XRD were used. The results are given in Figure IV-16, in which figure a) shows the XRD patterns of samples deposited with dry CG at different temperatures. At 290 °C, no reflection peaks were obtained, thus indicating a low degree of crystallization. At 305 °C, the film was already crystalline and a unique reflection corresponding to the (111) index could be

Figure IV- 17 .

 17 Figure a)shows the surface of a normal Cu2O film deposited at 335 °C with dry carrier gas. It presented a dense and smooth surface, grain size was around 150 nm. When 2 L/min gas bubbled through water (with 0.38 % water, W02), the morphology is shown in figure b), the surface of Cu2O film formed more visible grains out of the plain. Each grain had a wider size on bottom and smaller on the top with a triangle shaped plateau. The bubbler flow was further tuned up to 4 L/min (with 0.75 % water, W04), 6 L/min (with 1.12 % water, W06) and 8 L/min (with 1.37 % water, W08). The morphology was rather similar in those conditions, thus presented in figurec). In the figure, the triangle plateau shrank to a point, thus surfaces with pyramids structure were formed. In addition, bigger cubic grains emerged with maximum size up to 400 nm.

Figure IV- 17 .

 17 Figure IV-17. Cu2O films deposited by AA-MOCVD at 335 °C with different amount of introduced water molar ratio: a). 0 % (dry, D335), b). 0.38 % (W335-02) and c). 1.37 % (W335-08). The crystallinity of those samples was also checked by Raman Spectroscopy and X-Ray Diffraction. The patterns are given in Figure IV-18, in which the figure a) shows an evolutionof Raman peaks with regarding to the water molar ratio in gas flow. The very bottom Raman spectrum is a normal Cu2O thin film deposited with dry CG (Sample D335). As we can see, in a normal Cu2O sample deposited by AA-MOCVD, the first and second strongest peaks located at 170 and 210 cm -1 separately, both peaks were intense and sharp, indicating a high crystallinity. While moving to the curves above, more humidity was incorporated during the sample deposition process, with several changes worth noticing. First of all, intensity of the two tiny peaks located at around 150 cm -1 was enhanced. Other than that, an obvious peak

  Figure IV-18 b), the XRD patterns confirmed the sample grain orientation change. The major reflection for the normal Cu2O sample (D335) corresponded to the (110) orientation (in plain). While with more humidity added during deposition, the main reflection peaks were shifted to the (111) index. In all, with a more humid CG during AA-MOCVD deposition, the morphology of resulting Cu2O films evolved from flat surface corresponding to (110) orientation to "pyramids-like" rough surface with cubic grains aligned to (111) orientation, meanwhile they still remained as dense and continuous films.

Figure IV- 18

 18 Figure IV-18. a).Raman spectra and b) XRD patterns of samples deposited with different amount of water molar ratio. In order to give a direct view of different grain orientation on the surface, samples were characterized by EBSD on an area of about 2 × 5 µm 2 (Figure IV-19). Samples deposited in dry CG condition, exhibited a general surface grain preferential orientation along (101) (plane equivalent to (110)). The alignment of cubic grains on the surface is illustrated in figure a), with a diagonal facet pointing out of the plane. In the SEM image in figure b) and the reconstructed EBSD mapping in figure c), the grains corresponding to (101) were marked in green color, while the white spots were due to the lack of signal from such a rough surface. Such (101) orientation was also confirmed by the inverse pole figure shown in Figure IV-19 d), which showed a broad circle at the center of the diagram (implying an wide orientation distribution around (101)). On the other hand, deposition with humid CG yielded crystals mostly oriented along (111), as illustrated in Figure IV-19 e), with the corner of the cubic structure pointing out of the

Figure IV- 19 .

 19 Figure IV-19. EBSD characterization of Cu2O samples deposited without (D335) and with (W335) humid carrier gas at 335 °C for 220 mins. Deposition with dry carrier gas: a). Secondary electron image of D335 with titled angle of 70°; b). Recreated image based on EBSD characterization, grains oriented along (110) are marked in green; c). Orientation distribution of D335 in inverse pole figure. Sample deposited with humid carrier gas: d). Secondary electron image of W335 tilted 70°; e). Recreated EBSD image, grains oriented along (111) marked in blue; f). Orientation distribution of W335 in inverse polefigure.

Figure IV- 20

 20 Figure IV-20. a) Low magnification TEM cross-section image of a Cu2O film with (111) texture deposited at 335 °C on corning glass. b) Corresponding SAED pattern . c) Close up image showing a columnar growth developing from smaller crystals at the interface. Transmission Electron Microscopy (TEM) was also used to evaluate Cu2O film quality and morphology evolution along the thickness of the films. The results are presented in Figure IV-20, in which the cross-section and SAED pattern of sample W335 are presented. Figure IV-20 a) and b) present a low magnification cross-section image of the film and the corresponding electron diffraction pattern, respectively.

Chapter 4 . 7 Figure

 47 Figure IV-21. Effects of deposition temperature using dry and humid CG on the properties of Cu2O thin films. a) deposition rate; b). Resistivity; c). Carrier concentration and d). Carrier mobility. (Samples deposited with dry CG are represented with red squares and samples deposited with wet CG are marked with blue triangles.) Besides the effects of carrier gas humidity on morphology and grain orientation, the evolution of deposition rate and electronic properties with different deposition temperatures in dry and humid conditions were also compared, as presented in Figure IV-21. As shown in Figure IV-21 a), by dividing the thickness over deposition duration, the obtained deposition rate increased with deposition temperature rising from 305 to 350 °C . This was likely due to a more efficient precursor decomposition thanks to enhanced thermal activation of the reaction. At 365 °C the deposition rate decreased, most likely due to a partial pre-

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -136 -decomposition at the highest temperature used. The trend of deposition rate versus temperature was rather similar in both cases. In Figure IV-21 b), the resistivity of Cu2O films deposited in both dry and wet conditions at different temperatures is presented. Samples deposited at 290 °C could not be characterized due to the high resistivity presented as a result of the discontinued nature of the film. As it is shown, in both cases resistivity decreased with higher deposition temperatures. Thin films deposited with humid CG presented higher resistivity than the equivalent films deposited with dry CG. The minimum resistivity was obtained with the deposition with dry carrier gas at 365 °C, about 60 Ω.cm. The lowest resistivity for Cu2O thin films deposited with humid carrier gas was around 110 Ω.cm. As shown in Figure IV-21 c), the charge carrier concentrations for Cu2O films deposited with dry and wet carrier gas were compared. For Cu2O films deposited in dry condition, the carrier concentration stabilized at about 4 x 10 15 cm -3 . While for samples synthesized in humid condition, slightly lower carrier concentrations of about 1.7 x 10 15 cm -3 were obtained for all the temperatures. Furthermore, the evolution of mobility in both deposition conditions are presented in Figure IV-21 d). The figure showed, as expected from film morphology, an

Figure IV- 22 .

 22 Figure IV-22. Effect of humidity incorporation on Cu2O thin film a) growth rate and b) on the electronic properties at 335 °C.

Chapter 4 .

 4 Figure IV-21). Meanwhile, in Figure IV-23 b) and d) are the conductivity mapping of the corresponding surface areas, in which the darker color represents higher current measured, thus better conductivity. Comparing the AFM image and the correspondent electrical mapping, it's rather clear to point out that the grains have less conductivity than the grain boundaries, as most of the dark points are located between the grains.

Figure IV- 23 .

 23 Figure IV-23. AFM images of (111) textured Cu2O film (W335-08) a). with scan size of 500 nm c) with scan size of 5 μm and corresponding conductivity mapping b) and d).

  5 μm 2 on both Cu2O films, as shown in Figure IV-24 a) and c) for W305-08 and Figure IV-24 b) and d) for W335-08. The dark red areas Chapter 4. Deposition of highly crytsalline Cu2O films via AA-MOCVD -139 -correspond to the parts with higher current measured, thus indicating better conductivity. The conductive areas are mostly located along the grain boundaries, as the area shown with red marks in figure a) and b). By selecting the spots on grain and grain boundary, the AFM tip moved to the positions as marked in blue cross. At each spot, the applied potential varied from -5 V to +5 V and the current was recorded, thus the I-V curve for W305-08 is shown in Figure IV-24 e) while for W335-08 is presented in Figure IV-24 f). In both Figure IV-24 e) and f), weak rectification behaviors were observed on the Cu2O grains, while on the grain boundary, current could pass through either directions thus without showing any rectification. The resistivity in the grain boundaries was again much lower than in the grain. To verify the overall rectifying behavior, 9 mm 2 size Au contacts were deposited directly above the Cu2O films. By measuring the I-V curve, no overall rectification was found as expected.

Figure IV- 24

 24 Figure IV-24. a) and b) are AFM images of (111) textured w305 and w335, with 60 nm total height difference; red marks show areas of high conductance; c) and d) are the current mapping with voltage bias -2 V. Blue crosses indicate where IV shown in (e) and (f) were obtained.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -140 -In previously mentioned Cu2O films characterized by TEM SAED and XRD (Figure IV-20 and Figure IV-18), the quantity of CuO phase was below the detection limit. Meanwhile with the ultra-fine resolution from the ASTAR technique, CuO phase was observed, as presented in Figure IV-25. In figure a), the reconstructed correlation index map is rather similar to the cross-section image shown in Figure IV-25, showing the columnar growth of the films. The gray scale map as presented was obtained by plotting the value of the correlation index for the succession of the electron diffraction patterns at each location.In figure b), a phase mapping of the Cu2O cross-section is given. The map confirms that the film is mainly composed of Cu2O phase (colored in red). Meanwhile, the blue colored CuO phase was found and mostly located between the column Cu2O grains. The existence of CuO phase in grain boundary has previously been confirmed by XPS in the work of J. Deuermeier et al.25 Especially at the interface between substrate and the Cu2O film, where lots of smaller Cu2O grains have been reported in TEM image, more traces of CuO phase are observed.

Figure IV- 25

 25 Figure IV-25. ASTAR (TEM-EBSD like technique) Automatic crystal orientation/phase mapping of Cu2O sample (W335-08) crossection with orientation (111). a). Reconstructed correlation grey scale index map. b). Phase map (with index-reliability). The red color corresponds to Cu2O phase, the blue color representsCuO phase and the green color refers to the Si substrate.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -141 -

Figure IV- 26 .

 26 Figure IV-26. The possible origin of rectifying behavior in electrical AFM measurements of Cu2O films a) on grains with rectification and b) on grain boundary without rectification. In order to explain the different rectifying behavior when the Pt coated AFM tip was located on grains and grain boundaries, the corresponding energy diagrams of these two different locations of the film are drawn in Figure IV-26. Since Cu2O is a p type semiconducting material, thus the conductivity is mainly dominated by the hole conductivity in the Valence

  4). A p-n junction was fabricated following the structure of ITO/Cu2O/ZnO/Ag contact, as shown in Figure IV-27 (Ag contact not in the image). Figure IV-27 a) and b) present the top and cross-section view of the p-n junction. The top ZnO layer is about 70 nm thick with good conformity. The dense Cu2O film is about 900 nm thick.

Figure IV- 27

 27 Figure IV-27. p-n junction with structure Ag/ ZnO (60 °C, 70 nm) / Cu2O (111)/ITO. a) SEM images of top and b) cross-section view of the p-n junction; c). Rectifying I-V curve. (Scale bars in images correspond to 1 µm).

  depositions were performed with short durations (10 to 50 min). The morphology, mainly the roughness of the different samples, was characterized by Atomic Force Microscopy (AFM), the results are presented in Figure IV-28. The same deposition conditions for samples D335 and W335 in Table IV-4 were used for this study.

Figure IV- 28

 28 Figure IV-28 a), b) and c), presents the morphology of Cu2O films deposited for 10 min, 20 min and 50 min with dry carrier gas. Meanwhile, in figure d), e) and f) are depositions under humid condition with the three different durations. By comparing the morphology of the depositions under dry and wet conditions, we can see that with dry CG the growth of initial

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -143 -

Figure IV- 28 .

 28 Figure IV-28. AFM 3D images 1 × 1 μm 2 of samples deposited at dry carrier gas condition with different duration: a) 10 min; b) 20 min; c) 50 min. Samples with wet carrier gas and different duration: d) 10 min; e) 20 min; f) 50 min.g) sample was deposited with 10 min humid gas and the switching to dry deposition for 1 h. h) Roughness of samples deposited with different durations in wet and dry carrier gas conditions; i) X-Ray Diffraction patterns of samples with dry, wet carrier gas and sample starting with 10 min humid gas and followed by 40 min dry gas deposition.Conversely, films deposited under wet CG seemed to develop a SK growth mode, considering bigger grains as the crystallization islands. Especially after 50 min where big

10 min and followed by 40

 40 min deposition using dry CG (in total 50 min deposition). The morphology of this sample is presented in Figure IV-28 g). Its surface is rather similar to the film deposited with humid CG for 50 min (Figure IV-28 f), with big crystals being clearly observed (the same applies for the roughness value obtained, see Figure IV-28 h). XRD was also performed for the different samples (Figure IV-28 i), and the results confirmed that the presence of humidity induced the existence of a different and clear texture for such thin samples (thickness around 80 nm). XRD patterns also revealed the crucial role of water in the initial stages of thin film growth, showing that the preferential orientation of the sample obtained from humid-then-dry CGs was textured along (111), similarly to the sample deposited with humid CG, even though the intensity of reflections was slightly lower.

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -145 -

Chapter 4 .

 4 Deposition of highly crytsalline Cu2O films via AA-MOCVD -146 -

Figure IV-17 and Figure IV- 22 .

 22 With more water vapor incorporated during deposition, the average deposition rate was enhanced and the grain size was increased.

Chapter 5 .

 5 Deposition and characterization of AgCuO2 thin films -155 -

Figure V- 1 .

 1 Figure V-1. SEM images of co-deposition samples in Plasma Enhanced AA-MOCVD for samples a) C4, b) C5 and C6. SEM images of samples C4, C5 and C6, are presented in Figure V-1. The three samples presented smooth surfaces with small grains. The conductivity of sample C6 (about 40 ~ 50 nm thick) was close to metallic films with sheet resistance 0.03 Ω/□, while the other two were rather resistive. The EDS analyses of those samples were conducted with 10 kV voltage and the results are presented in Figure V-2. Comparing all EDS analyses in figure a), b) and c), an excessive of Cu was always found in the deposited samples. Comparing figure a) and c), the 300 W plasma didn't significantly change the element ratio in the resulted thin films (at 270 °C roughly about Ag: Cu= 1: 4). At even lower deposition temperature, less copper precursors were decomposed thus an increase of silver atomic ratio was observed in Figure V-2 b).

Figure V- 2 .

 2 Figure V-2. EDS analyses of co-deposited samples of samples a) C4, b) C5 and c) C6.

Figure V- 3 .

 3 Figure V-3. EDS analyses of samples a) M1, b) M2 and c) M3. The morphology of the optimized Ag0.5Cu0.5 films is shown in Figure V-4, in which figure a) is the secondary electron image and figure b) is the backscattering electron image of sample M3. They show continuous AgCu metal alloy films with rather homogeneous element distributions. It's difficult to judge the elemental distribution from the backscattering electron image with the particles on surface from figure b), thus further EDS element mapping for both Cu and Ag atoms was conducted. The results are shown in Figure V-4 c) and d), an overall homogeneous element distribution was thus confirmed. These Ag0.5Cu0.5 metallic alloy thin films were then used for further oxidation tests to obtain the AgCuO2

Figure V- 4

 4 Figure V-4. a) Secondary electron image and b) Back scattering electron (BSE) image of an AgCu alloy film (sample M3) with homogeneous Cu, Ag element distribution in c) and d).

  voltammetry on the samples to detect the different electrochemical processes taking place and the voltages at which those occur. Similar as for Ag films, a 3 electrode cell was used with Pt foil (size about 1 cm 2 ) as counter electrode, Ag/AgCl 3.5 M KCl as reference electrode and AgCu alloy films as working electrode in 1 M NaOH electrolyte solution. The cyclic voltammetry of AgCu alloy film was obtained in the potential window 0 V→ 1 V→ -1 V → 0 V with a scanning speed of 10 mV/s. The obtained voltammogram is presented in Figure V-5.

Figure V- 5 .

 5 Figure V-5. Cyclic voltammetry of AgCu alloy film with scanning speed of 1 mV/s, from 0 → 1 → -1 → 0 V. Scanning was conducted with Pt as counter electrode and 0.35 M AgCl/KCl as reference electrode in 1 M NaOH solution.Thus, we used previously reported electrochemical works of D.Muñoz-Rojas et al.,12 and several electrochemical oxidation experiments were designed with parameters as shown in Table V-3. The electrochemical oxidation of AgCu films was conducted under different potentials ranging from + 0.5 V to + 0.7 V. Besides, the effect of electrolyte concentration (NaOH solution) was also studied. As for the AgCu alloy thin film samples, even though the deposition parameters to obtain Ag0.5Cu0.5 composition films were optimized, unfortunately, due to the reproducibility problem of the PI-MOCVD system, the later deposited AgCu films have less copper element incorporated.

  2 M NaOH solution at various potentials are shown in Figure V-6. In figure a), 0.5 V was applied for 17 min (sample EA01), on the surface both big cubic and needle like structures were observed. With higher potential, 0.6 and 0.7 V (sample EA02 and EA03), the needle-like particles became more prominent as shown in Figure V-6 b) and c). In figure d), with longer oxidation time at 0.7 V (sample EA03b), the sample surface was mostly covered by needle structured grains.

Figure V- 6 .

 6 Figure V-6. SEM of electrochemically oxidized AgCu film in 0.2 M NaOH solution with different potentials: a) 0.5 V (sample EA01), b) 0.6 V (sample EA02), c) 0.7 V (sample EA03) for 17 min and d) 0.7 V (sample EA03b) for 34 min. In addition to the morphology characterization, the GIXRD patterns of the different samples were recorded and are shown in Figure V-7. Since those electrochemically oxidized sampleswere silver rich films, in the XRD of the original AgCu alloy film, only a weak peak reflection corresponding to Cu phase was obtained. In all XRD patterns, the reflection with the strongest intensity corresponds to metallic Ag, indicating a majority of Ag phase in all these samples. When 0.5 V were applied for the oxidation (EA01), a reflection corresponding to Ag2O phase appeared. Similar XRD peaks corresponding to CuO were also observed with 0.7 V potential (sample EA03). Meanwhile, in the other two cases, the XRD patterns didn't show any change comparing with the original XRD reflections. In addition, no peak reflections corresponding to Ag2Cu2O3 or AgCuO2 phase was observed in these oxidation conditions.

Figure V- 7 .

 7 Figure V-7. GIXRD patterns of electrochemically oxidized AgCu film in 0.2 M NaOH solution at different potentials ranging from 0.5 to 0.7 V.In addition to the oxidation of AgCu in 0.2 M NaOH solution, oxidations were also carried out with higher electrolyte concentration, namely, 0.5 and 1 M. The morphology of the samples obtained in 0.5 M NaOH solutions are presented in Figure V-8, in which at lower potentials, 0.5 and 0.6 V, both cubic and need like particles are visible on the surface as shown in figure a) and b). With increased potential to 0.7 V, the needle-like grain structure became prominent on the surface, as presented in Figure V-8 c). In Figure V-8 d), at 0.7 Vwith longer oxidation duration (34 min), the needle like grains didn't appear on the surface, which could be because of reproducibility issues with our oxidation setup.Concerning the phase change during the electrochemical oxidation process, FigureV-9presents the emergence of a reflection at about 32° for Ag2O in all cases. With longer duration, the reflection peak intensity corresponding to Ag2O increased. Thus, comparing with previous works on CuO nano structures,[17][18][19] those needle like grains are more likely CuO grains. Meanwhile, considering the cubic structure of Ag2O grains, those cubic grains on surface are more likely to be the Ag2O identified in XRD.

Figure V- 8 .

 8 Figure V-8. AgCu films electrochemically oxidized in 0.5 M NaOH solution with potentials, a) 0.5 V (EB01), b) 0.6 V (EB02), c) 0.7 V for 17 min (EB03) and d) 0.7 V for 34 min (EB03b).

Figure V- 9 .

 9 Figure V-9. GIXRD of AgCu films electrochemically oxidized in 0.5 M NaOH solution with various potentials. Similar experimental setups were used as introduced before.

  The GIXRD patterns of those samples are shown in Figure V-11, in which reflections corresponding to Ag2O phase appeared, similar to the previous results. Since in SEM, only the nanoplates clusters are observed, thus inform the XRD data, these structures can be associated with the Ag2O phase. The change of the grain shape is more like due to a higher electrolyte concentration.

Figure V- 10 .

 10 Figure V-10. Morphology of AgCu film oxidized electrochemically in 1 M NaOH solution by applying potential a) 0.5 V (EC01) and 0.7 V (EC02) for 17 min.

Figure V- 11 .

 11 Figure V-11. GIXRD of AgCu film oxidized electrochemically in 1 M NaOH solution by applying potential a) 0.5 V and 0.7 V for 15 min.

  presented in in Figure V-12. Figure V-12 a) shows the the surface morphology of the asdeposited AgCu alloy film. By applying 300 W power of pure oxygen plasma on heated substrate (at 75 °C) for 3 min, the morphology did not change much, as shown in Figure V-13 b). With increased plasma power to 400 W, after 3 min treatment (at 75 °C), hollow structures were visible on the sample surface, as presented in Figure V-12 c). Figure V-12 d) shows the morphology obtained after treating the alloy for 3 min with a plasma of 300 W, as shown in Figure V-13 b), but at higher temperature (200 °C). In this case, there is a significant change in film morphology, with big particles visible on top of the film.

Figure V- 12 .

 12 Figure V-12. SEM images of a) as-deposited AgCu film, and films oxidized at 75 °C for 3 min using a 300 W (b) and 400 W (c) plasma (PL1 and PL2). d) Film oxidized at 200 °C for 3 min using 300 W plasma (PL3).The GIXRD patterns of these samples are shown in Figure V-13. Similar with previous case, the intensity of reflections corresponding to Cu was rather weak. By comparing the phase change between different plasma power and treatment temperature, the latter seems to play a more critical role for increasing the reactivity of the metal atoms thus obtaining a strong signal from Ag2O, as shown by the sky blue curve. But unlike the simple case with Ag film only (introduced on Chapter III), except the Ag2O phase no other phases were obtained. Thus, no reflections corresponding to silver copper mixed oxide phase was observed. In a work of D.Muñoz-Rojas et al.,11 the solid phase transformation from Ag2Cu2O3 powders to AgCuO2

Figure V- 13 .

 13 Figure V-13. GIXRD patterns of AgCu film as deposited and treated by oxygen plasma with different power and temperatures.

Figure V- 14 .

 14 Figure V-14. Schematic presentation of the SILAR process for the deposition of metal oxides thin films. The image to the left shows the actual setup used.

  Using these deposition parameters, sample A0 was coated on Cu2O/glass substrate. The morphology of the coating is presented in Figure V-15 a) and b). The secondary electron image in figure a) shows a sample with big grains scattered on the surface. The grain size reaches 300 nm, as shown in the inserted close-up image.

Figure V- 15 .

 15 Figure V-15. Characterizations of first synthesized AgCuO2 film by SILAR deposition (sample A0): a) SEM secondary electron image and corresponding b) SEM backscattering image. c) GIXRD pattern.

Chapter 5 .

 5 Deposition and characterization of AgCuO2 thin films -172 -substrate (thus inducing charging effects) or the to the heavier elements being more concentrated in those scattered particles.The GIXRD pattern of sample A0 is presented in Figure V-15. As it is shown, the strongest reflection corresponds to the Cu2O phase (from the substrate). Besides that, several unique reflections corresponding to AgCuO2 phase were observed. Those unique reflections are locating around 16 °, 31.5 °, 32.3 ° and 41.5 °. Thus, AgCuO2 phase could be successfully synthesized through the SILAR approach. Meanwhile, in addition to the desired phase, rather clear reflections associated with CuO phase were also observed in the XRD pattern. Thus, the optimization of AgCuO2 thin film deposition in SILAR was carried out from two aspects, morphology and phase purity.

  Figure V-16 b) and cross section in d)). In addition to that, the inserted back scattering image shows a uniform element composition on the surface. With even longer duration, the surface still remained continuous but the BSE image showed rather clear contrast, which is more likely due to the rough morphology in this case (sample A3), as shown in Figure V-16 c).

Figure V- 16 .

 16 Figure V-16. SEM images of AgCuO2 samples deposited by SILAR with different immersion times in metal ion precursor solution. a) 10 s (A1), b) 20 s (A2), c) 27 s (A3) and d) cross section of b).

  Similar to the grains presented in Figure V-15, they have a platelet-like shape but with much small grain size (about 70 nm in length).Due to the vertical geometry of the SILAR deposition system, any particles in the solution could have the possibility to attach on the substrate surface, thus there are several bigger clusters remaining on the deposited film. The corresponding BSE image is placed on the top-right corner, in which a rather uniform element distribution across the surface is observed. The slight white and black contrast in the backscattering electron image is more likely attributed to the morphology of the sample.With slightly decreased copper ion concentration down to 0.4 M, the surface morphology of the deposited film (sample B2) has a more flat surface as shown in figureb). The shape of the grains changed from platelet-like to more cubic-like particles, with length around 80 nm.Meanwhile, the inserted BSE image shows a rather homogeneous element distribution. With a further decrease in the the copper ion concentration (sample B3) is presented in figure c), similar platelet-like structures were also observed on the sample surface, but with a much bigger size (with length up to 200 nm). The film still showed a dense and continuous aspect.The inserted BSE image also shows a homogeneous element distribution. Lastly, with further lowering the copper ion concentration, a smooth thin film surface was obtained, as presented in figured). In this case the film is made up of much smaller particles (about 70 nm) than for samples deposited with higher Cu concentration (sample B4). The backscattering electron image also shows an even element distribution between silver and copper on the surface. Due to a smoother surface morphology, less contrast is present in the image.

Figure V- 17 .

 17 Figure V-17. SEM images of AgCuO2 films deposited by SILAR with different copper ion concentrations. a) 0.04 M (B2), b) 0.035 M (B3), c) 0.03 M (B4) and d) 0.028 M (B5). In addition to SEM, the surface of sample B5, being the most continuous and smooth, was also characterized by AFM, as shown in Figure V-18, in which the recreated 3 D morphology is presented. As it is shown the grain is rather dense and continuous with surface roughness RMS about 8 nm.

Figure V- 18 .

 18 Figure V-18. Reconstructed 3D AFM images of AgCuO2 films deposited by SILAR using optimized conditions (Sample B5) with RMS value about 8 nm. To identify the phases deposited, GIXRD characterizations were again conducted and the results are shown in Figure V-19. In all the XRD patterns, reflection corresponding to

Figure V- 19 .

 19 Figure V-19.GIXRD patterns of process of optimization the purity of AgCuO2 films by reducing the copper ion concentration. Even though, in GIXRD the accurate composition of each phase can't be obtained from GIXRD patterns, by comparing the intensity of the reflections corresponding to impurity phase with the peak intensity of AgCuO2, the trend of the amount of impurities in the films can be presented relatively. As shown in Figure V-19, with gradually decreasing the copper ion concentration in solution, the peak reflection intensity corresponding to CuO decreases until fully disappears for the samples deposited with 0.03 M (sample B4) and 0.028 M

400 s using 2

 2 cycles. The obtained Raman spectra for samples B1 and B4 are presented in Figure V-20. As it is shown, the Raman spectra of AgCuO2 films deposited by SILAR have only three major peaks located at 220.2, 439 and 521.4 cm -1 , which is in agreement with the reference curve obtained from pure AgCuO2 powders. In the Raman spectra of sample B1, a peak at around 150 cm -1 was observed which is associated with the Cu2O phase.

Figure V- 20 .

 20 Figure V-20. RAMAN spectra of the synthesized AgCuO2 with comparison to the powder references. Meanwhile, the similar peak didn't emerge in B4 Raman spectra. Combining their morphologies presented in Figure V-17 a) and c), sample B4 possibly had a denser or thicker film than B1. By comparing Raman spectra from the thin films with the AgO, CuO reference spectra, no visible peaks were observed corresponding to the two impurity phases, indicating a lower sensitivity of Raman characterization as compared with GIXRD. For both samples,

Chapter 5 .V. 3 . 5

 535 Deposition and characterization of AgCuO2 thin films -178 -Therefore, combining all the optimized parameters, the deposition of continuous and almost pure phase AgCuO2 films in SILAR could be achieved with the following parameters. The precursor solution contained 0.05 M Ag 1+ and 0.03 M Cu 2+ . The cation solution is composed of 1 M NaOH and 0.08 M Na2S2O8. The immersion and withdraw speed in Z axis are set to 4 mm/s and 60 mm/s. The immersion duration in cation solution (t1), buffer solution 1 (t2), metal ion solution (t3) and buffer solution 2 (t4) is 15 s, 10 s, 20 s, 10 s, respectively. To obtain a 100 nm thick AgCuO2 film, 70 cycles of deposition are required. Understanding the mechanism of AgCuO2 thin film coating by SILAR process.

Figure V- 21 .

 21 Figure V-21. Morphology of coated samples by SILAR on Cu2O/glass substrate: a) deposition with only Ag 1+ ion precursor, without oxidizer (sample C1); b) deposition with only Ag 1+ ion precursor with oxidizer (sample C2); c) coatings with both metal ion precursors, without oxidizer (sample C3); d) standard thin film coatings for comparison (sample C4). The morphology of coatings on ZnO/glass substrate: e) deposited with only Ag 1+ ion precursor without oxidizer (sample Z1); f) deposited with both metal ion precursor with oxidizer (sample Z4).In Figure V-21 e), the morphology of sample Z1 is shown. The sample was deposited with only Ag 1+ precursor and without oxidizer. On the surface, the bottom ZnO layer has been largely etched away by the strong basic environment, thus the porous structures left behind.On top of the porous layer, the coated grain particles scatter on the surface. In the corresponding BSE image, much higher intensity is observed on these grains. Lastly, Figure V-21 f) presents the morphology of sample Z4, which was deposited with both ion precursors and oxidizer. On the sample surface, big plate like grains (size about 300 nm) are dispersed on the substrate. Unlike in figure e), no porous structures were obtained in the bottom layer. By further checking the BSE image, a very weak electron signal is presented on the substrate, while grains appear very bright, which indicates that the ZnO layer was possibly totally dissolved in this case. Since the substrate was ZnO, the EDS analyses of the deposited grains presented Ag:Cu atomic ratio very close to 1:1.To obtain the information of the phase composition of those samples, GIXRD was conducted and the results are shown in Figure V-22. The XRD of SILAR coatings on Cu2O/glass substrate are presented in Figure V-22 a). As it is shown, for sample C1, the major reflections correspond to Ag2O phase, which is rather reasonable since the sample was deposited without oxidizer. Meanwhile, and quite surprisingly, the XRD reflections for sample C2

Figure V- 22 .

 22 Figure V-22. GIXRD patterns of SILAR coatings on a) Cu2O/glass and b) ZnO/glass substrate with different conditions.Since this sample C2 was coated without Cu 2+ ion precursor, the formation of AgCuO2 indicates that part of Cu2O on the substrate was also participating in the SILAR reaction, which might explain the difficulty to reproduce pure AgCuO2 phase thin films with no CuO contamination. For sample C3

Figure V- 23 .V. 3 . 6

 2336 Figure V-23. Comparing Raman spectra of SILAR coated AgCuO2, Ag2Cu2O3 with powders synthesized through hydrothermal reactions from reference work. To further confirm the phase of those SILAR coated thin films, sample Z4 (Ag2Cu2O3 identified by GIXRD), C4 ( AgCuO2 identified by GIXRD) and C2 (AgO+AgCuO2) were characterized by Raman spectroscopy, the results are presented in Figure V-23. In this Raman characterization, a green laser (with wavelength 514.5 nm) was applied with integration duration 600 s × 2 cycles. For AgCuO2 sample (C4), comparing with the Raman spectra shown in Figure V-20, with green laser, stronger signal and less noise were presented and the same three peaks at the same positions were observed. The reference Raman curve for Ag2Cu2O3 was obtained by characterizing the powders synthesized based on reference work31 . The Raman signals from Ag2Cu2O3 powder or SILAR coating were much weaker and less visible. However, three peaks at 311.5, 494.6 and 548 cm -1 were found in both the reference powder and SILAR coating. For sample C2, the Raman spectra presented peaks from both AgCuO2 and AgO phases, confirming the formation of AgCuO2 on the Cu2O substrate (with only silver precursor).V.3.6 Optical and electronic property of AgCuO2 films obtained by SILAR

Figure V- 24 .

 24 Figure V-24. Optical characterizations of AgCuO2 films coated with various copper ion concentration by SILAR process: a) Transmittance curve and b) reflectance curve. The red transmittance curve corresponding to sample B4 presents the lowest transmittance in visible range, but the onset of transmittance starts at the same wavelength as the other samples, about 420 nm. Since the thickness of all the deposited film is almost the same, the most possible reason is due to a rougher surface (as presented in Figure V-17 c) above). Besides the transmittance, the reflectance curves of those samples are given in Figure V-24 b). The Cu2O seed layer gives an average 35% to 40% reflectance throughout the whole spectrum but reaches down to a minimum at about 10% at 590 nm. Meanwhile, the AgCuO2 films coated by SILAR exhibit higher reflection in the visible light range, ranging from 20% to 30%. In addition to those characterizations, the Tauc formula (assuming both direct and indirect bandgap) was applied on the transmittance curve in order to estimate the bandgap of AgCuO2, as shown in Figure V-25. In figure Figure V-25 a), the Tauc plot was obtained assuming a direct bandgap. The curve with black empty diamond represents the Tauc plot for the Cu2O seed layer used in this work, with direct bandgap about 2.5 eV. With the layers coated on top of Cu2O by SILAR, the onset of absorption was largely decreased. By following the linear part and plotting their interception with the X axis, these curves all yielded a bandgap of about 1.2 eV, regardless of the adjustment of the copper ion concentration,. Meanwhile, figure b) presents the Tauc plot assuming an indirect bandgap. As it is shown, the absorption coefficient versus the photon energy is not a straight line. Therefore, the SILAR deposited

Figure V- 25 .

 25 Figure V-25. Optical bandgap of AgCuO2 film deposited by SILAR using Tauc formula with different copper ion concentration used: a) plot as direct bandgap and b) as indirect bandgap.

Chapter 5 .

 5 Deposition and characterization of AgCuO2 thin films -186 -carrier concentration and mobility reached an optimal value, namely, 1.55 x 10 18 cm -3 for carrier concentration and 24 cm 2 /V.s for the mobility. This increase in electrical conductivity can be related with a higher purity AgCuO2 phase, as previously reported in Figure V-19.

Figure V- 26 .

 26 Figure V-26. Electronic properties of AgCuO2 film deposited by SILAR using Hall effect measurement with different copper ion concentration used. Values with 10% error bar.

Figure VI- 1

 1 Figure VI-1. a). Image of the capillary in situ reactor used: A: Sapphire capillary. B: Connections for external tubing. C: Sliding end connection. D: Parallel rails for preventing strain on the capillary. E: Goniometer head connection pin. 10 b). In situ XRD experimental setup at line BM01A, including X-Ray source, 2D detector, heater, capillary cell, goniometer and CCD camera.

Figure VI- 2 .

 2 Figure VI-2. Normalized GI-XRD reflections of samples A1, A2 and A3 characterized in the lab.To understand the hydrothermal reaction process of AgCuO2, similar experiments were repeated and characterized by in situ XRD in the synchrotron, as shown in detail in TableVI-2. For the sake of simplicity, here we define the standard hydrothermal reaction conditions as the equimolar mixture of Ag and Cu precursors reacting in 1 M KOH solution at 90 °C with 250 bar of pressure. All suspensions were prepared right before the synchrotron characterization, no more than 10 minutes. To study the effect of pH during the hydrothermal reaction process, the same concentrations of KOH solutions than the ones used in the lab were used in sample S1, S2 and S3 in the synchrotron.

Figure VI- 3 .

 3 Figure VI-3 c), the identified phases from the reflections are presented. The starting precursors of the reaction were AgO and CuO (obtained from CuSO4 precipitation in 1 M KOH solution in the starting slurry). The reflection at low diffraction angle (about 7 °) was confirmed to correspond to AgCuO2 phase. In addition, several other reflections with increasing intensity also matched AgCuO2. On the other hand, the intensity of reflections corresponding to Ag2O2 (AgO) and CuO phases rapidly decreased until almost disappeared from the diffraction. Meanwhile, a tiny amount of Ag2O phase from AgO reduction was also observed at diffraction angle 14.5 °. With a good solubility in water, the reflections of other chemicals, like Na2S2O8 and KOH were not present in the diffraction pattern. Even though the experiment was stopped after 250 seconds, the whole reaction was complete in only about 75 s. In the end, a powder was obtained with a majority of AgCuO2 phase, in agreement with the XRD patterns from the synthesis in lab shown in Figure VI-2 (sample A1).

Chapter 6 .

 6 In-situ XRD characterization of formation mechanism of AgCuO2 phase in hydrothermal reaction -198 -

Figure VI- 4 .

 4 Figure VI-4. Phase evolution of reaction S2, with 0.3 M KOH and 250 bar at 90 °C: a) 2 D, b) 3 D diagram and c) phase identification of selected XRD patterns. Lastly, in the case of the reaction without no added KOH, the in situ XRD characterization was also conducted (sample S3), and the results are shown in Figure VI-5. In figure a), a low background signal was presented in the 2D diagram. In addition, several short lines were observed during the reaction, indicating the existence of intermediate phases during the hydrothermal reaction without KOH added. Similar evidence of intermediate phases can be also found around 23 °. Lastly, the intensity of all the new formed lines were rather weak, which can be attribute a lower quantity of precipitation. Those results are further presented with a more clear view in the 3D diagram, as shown in figure b). At low diffraction angle, a reflection emerged right after the characterization began and then about 100 s afterwards

( 32 )

 32 Therefore, to conclude, the pH of the solution plays a critical role in AgCuO2 phase formation. Different pH during reaction can lead to the different reaction mechanism between CuSO4 and AgO thus resulting in different final products. From the absence of intermediate phase or shifting of peaks during the formation of AgCuO2 in 1 M KOH solution, possibility of solid phase transformation as the AgCuO2 formation mechanism in the alkaline solution is thus excluded.

Figure VI- 5 .

 5 Figure VI-5. Phase evolution of reaction S3, with no KOH with 250 bar pressure under 90 °C: a) 2 D, b) 3 D diagrams and c) phase identification of selected XRD patterns.

Chapter 6 .

 6 In-situ XRD characterization of formation mechanism of AgCuO2 phase in hydrothermal reaction -202 -

Figure VI- 6 .

 6 Figure VI-6. Extracted 2D XRD patterns from reaction (sample S1b) in 1 M KOH solution at 90 °C with 3 bar pressure.Combining the results obtained from the different reactions discussed so far, we propose the reaction mechanism for the hydrothermal synthesis of AgCuO2, as illustrated in FigureVI-7.At the beginning of the reaction, dissolved CuSO4 is mixed with a 1 M KOH solution. At such basic pH, CuO precipitates12 , which was also confirmed by the reflections at the beginning of the characterization as shown in FigureVI-3 c). Therefore, the actual reaction precursors in 1 M KOH solution are CuO and AgO in suspension.

𝐴𝑔𝑂 + 2𝑂𝐻 - 1 + 34 )Figure VI- 8 .

 1348 Figure VI-8. Evolution of AgCuO2 (111) reflections for reactions made in 1 M KOH solution with 250 bar pressure at a) 50 °C a), b) 70 °C, c) 80 °C, d) 90 °C and e) 120 °C. At 70 °C (sample T2), by comparing the reflection intensity, the reaction was completed after about 800 s. For reactions at 80 °C (sample T3) and 90 °C (sample T4), the obtained evolution of reflection intensity were rather similar. Reactions at both temperatures were completed after around 180 s. Lastly, Figure VI-8 e) presents the reflections at the highest reaction temperature (120 °C, sample T5) during characterization. By comparing the intensity change, only 30 s were required to complete the reaction. Therefore, since the

Table of content

 of

Table I -

 I 1. Optical and electronic properties of reported commonly seen p and n type metal oxide materials.

	p type metal	deposition/	resistivity	carrier	mobility	bandgap	refs
	oxides		synthesis	(Ω.cm)	density (cm -	(cm 2 /V.s)	(eV)
			technique		3 )
	SnO		PLD	n.a	2.5×10 17	2.4	2.7	24,25
	NiO		rf-MS	0.1~0.45	10 17 ~10 18	5~28	3.6~4	26
	Li doped NiO spray pyrolysis 0.4~1.98	1.9 × 10 17 ~	2.2~12	~2.9	27
					1.2×10 18
	ZnIr2O4	PLD	~25	n.a	n.a	2.97	28
	spinel type
	oxide
	Cu2O		rf-MS	~100	~10 15	60	2.7	16
	SrCu2O2	PLD	20	6.1×10 17	0.46	3.3	29
	CuGaO2	PLD	16.7	1.7×10 18	0.23	3.6	30
	CuCrO2	ALD	0.1~2.2	n.a	n.a	3.09	31
	CuCr0.93Mg0.	spray pyrolysis ~1	2×10 19	0.2~0.3	3.1	32
	07O2
	CuAlO2	thermal heating 1~10	1×10 17	10	3.5	33,34
	CuAlO2	MOCVD	0.5	1.8×10 19	0.16	3.75	35
	AgCuO2	Hydrothermal	0.01~0.05	n.a	n.a	n.a (~1	36
			(single crystal)				eV)
	n type metal
	oxides
	ZnO		AP-SALD	5 × 10 -2	10 19 ~10 20	5.5	3.3 ~ 3.5 37
	Hydrogenate	rf-MS	2.8× 10 -3	4.4×10 19	47.1	3.3 ~ 3.5 38
	d ZnO
	Al	doped	rf-MS	1.4× 10 -4	~10 21	36.8	3.5 ~ 3.7 39
	ZnO (AZO)
	Indium		Reactive ion	~10 -4	~10 26	~10	~ 3.9	40
	doped	Tin	plating
	Oxide (ITO)
	Fluorine	MS	6.71× 10 -3	1.46×10 20	15	3.8
	doped	Tin
	Oxide (FTO)

41

Note: Pulsed Laser Deposition (PLD), Atomic Layer Deposition (ALD), Metal Organic Chemical Vapor Deposition (MOCVD) and radio frequency Magnetron Sputtering (rf-MS).

Table I -

 I 2. Electronic and optical properties of Cu2O synthesized through different techniques.

	Deposition	Carrier	Mobility	Resistivity	Bandgap	References
	techniques	Concentration	(cm 2 /V.s)	(Ω. cm)	(eV)
		(cm -3 )
	AP-SALD	~ 8.5× 10 15	~ 7	~190	2.47	61,62
	MOCVD	~ 3.3× 10 15	~ 5	~ 1000	2.5	63
	PLD	~ 1× 10 15	~ 20	~ 200	n.a	22
	rf-MS	~ 1× 10 15	~ 60	~ 100	2.7	16, 66
	Electrochemical	~ 10 14	~ 1	~ 10 4	1.99	64
	deposition
	Spin coating	~ 1-2 × 10 15 ~ 18 to 5 n.a	2	65
	Thermal	10 13	>100	n.a	2	17
	Oxidation

II.1 Deposition of Ag and Cu2O thin films by Metal Organic Chemical Vapor Deposition (MOCVD).

	H., MacManus-Driscoll, J. L. & Berry, J. J. Improved fill factors in solution-processed Chapter II Experimental setup
	ZnO/Cu2O photovoltaics. Thin Solid Films 536, 280-285 (2013). Chapter II Experimental setup	--
	82. In this chapter, an introduction to the experimental techniques and equipment will be Illiberi, A., Scherpenborg, R., Wu, Y., Roozeboom, F. & Poodt, P. Spatial atmospheric atomic
	layer deposition of alxzn1-xo. ACS Appl. Mater. Interfaces 5, 13124-13128 (2013). II.1 Deposition of Ag and Cu2O thin films by Metal Organic Chemical Vapor Deposition presented. This introduction is mainly composed of four parts, including Metal Organic
	(MOCVD). Chemical Vapor deposition techniques, thin film oxidation setups, solution approaches to --
	II.1.1 Pulsed Injection MOCVD thin film deposition and different types of characterization techniques used in this thesis --
	II.1.2 Aerosol Assisted MOCVD work. The thin film deposition/oxidation techniques aim to the synthesis of AgCuO2 thin --
	II.1.3 Film deposition films, as introduced in the thesis objective. The synthesized films have been characterized --
	II.1.3.1 Preparation of starting precursor solutions: from the point view of morphology, elemental composition, phase identification, and optical --
	II.1.3.2 Substrate cleaning and electronic properties.	--
	II.1.3.3 Preparation of the MOCVD systems	--
	II.2 Oxidation techniques: electrochemical oxidation and oxidation with atmospheric
	oxygen plasma	--
	II.2.1 Electrochemical deposition and oxidation As one of the most widely used thin film deposition techniques, MOCVD can be used to --
	II.2.2 Oxygen plasma oxidation deposit a wide range of metallic or metal oxide coatings. In general, the MOCVD deposition --
	II.2.3 Deposition of AgCuO2 films by the SILAR method process relies on the delivering of precursor in vapor phase onto the reaction zone where a --
	II.3 Characterization techniques substrate is placed. With enough heating temperature or other means of reaction activation --
	II.3.1 Thin film crystallinity: RAMAN spectroscopy, XRD (i.e. plasma, light), metalorganic precursors can be decomposed and thus deposition takes --
	II.3.1.1 Raman Spectroscopy place. Depending on the property of the metal atom and the oxidation atmosphere in the --
	II.3.1.2 X-Ray diffraction (XRD): Bragg-Brentano and Grazing Incidence. deposition chamber, either metallic or metal oxide films can be coated with high quality. --
	II.3.2 Morphological characterization: SEM: EDS, EBSD, AFM Such type of thin film coating technique has the advantages of being scalable, low cost and --
	II.3.2.1 Scanning Electron Microscopy (SEM) being able to yield homogeneous dense coatings. As a consequence, MOCVD is a widely --
	II.3.2.2 Atomic Force Microscopy (AFM) applied technique in industry. After decades of developing such thin film coating technique, --
	II.3.2.3 Transmission Electron Microscopy (TEM) MOCVD has derived into several branches. Depending on how the precursor is fed into the --
	II.3.3 Ultraviolet-Visible (UV-Vis) spectroscopy deposition chamber, MOCVD is classified into several types, including conventional Hot --
	II.3.4 Electrical characterization Wall MOCVD (HW-MOCVD), Cold Wall MOCVD (CW-MOCVD), Aerosol Assisted --
	II.3.4.1 Four probe measurements MOCVD (AA-MOCVD) and Pulsed Injection MOCVD (PI-MOCVD), etc. 1 In this work, --
	II.3.4.2 Hall Effect measurements two types of MOCVD approaches were applied for the deposition of Ag and Cu based thin --
	II.4 References films, namely Aerosol Assisted MOCVD (AA-MOCVD) and Pulsed Injection MOCVD (PI---
	MOCVD

The structure of the precursor is presented in Chapter III. e). Ag(hexafluoro acetylacetonate) triglyme (

	AgT, new precursor obtained during this
	thesis) 6 . Compared with the previous two precursors, better solubility was found during the
	process of solution preparation. 0.01 M of precursor solution could be rapidly prepared by
	adding 2.45 g (0.005 mol) of AgT in 500 ml of ethanol with magnetic stirring for only 5 min,
	a clear grey colored solution being obtained. The molecular structure is presented in Chapter
	III.
	As for copper, two different copper precursors were used in this work:
	a) Cu(II) acetylacetonate (Cu(AcAc)2) (98%, STREM Chemicals). Starting solutions with
	a concentration of 0.01 M were prepared by adding 1.3 g (0.005 mol) of precursor in 500 ml
	of alcohol solvent (butanol-01 or ethanol). The mixture was stirred until complete dissolution
	using a magnetic bar with speed 300 cycle/min. Since Cu(AcAc)2 has low solubility in
	alcohol solvents, at least 0.01 M ethylene diamine (99%, Sigma Aldrich) was added as

Table III -

 III 1. Reported silver metal organic precursors used in metalorganic CVD. AgP, phen=1,10-phenanthroline) and [Ag(triglyme)2] +-[Ag(hfac)2] -(AgT, triglyme = 2,5,8,11-tetraoxadodecane). In addition, all the ligands used in the new precursors, namely, Hhfac, triglyme and phenanthroline, are not very expensive, thus reducing the final precursor cost. The synthesis and characterization of these new silver complexes is described in the next sections, along with the results of their use as precursors for the deposition of Ag coatings via AA-MOCVD with alcohol solvents.

	Metalorganic	CVD	Solvents	Deposition	Ref
	Precursor	variant		Temperature
				/ °C
	Ag(hfac)COD	Pulsed-	Ethanol,	150-400	11
		spray-	methanol,
		evaporation	monoglyme
		CVD
	(hfac)Ag(BTMSA)	HW-CVD	none	~ 250	22
	(hfac)Ag(SEt2)	AA-	Toluene	n.a.	18
		MOCVD
	[Ag(O2CC3F7)(PPh3)2] AA-	THF	310	19
		MOCVD
	Ag(fod)(PEt3 )	PI-	None (Liquid	220	24
		MOCVD	precursor)
	(hfac)Ag(VTES)	CW-CVD	None (Liquid	220	25
			precursor)
	(tfac)AgP(OEt)3	CW-CVD	None (Liquid	180	23
			precursor)
	(C2F5COO)Ag(PMe3) HW-CVD	n.a.	200	21
	Ag(hfac)(PPh3)	AA-	THF	310	17
		MOCVD
	Ag(O2CtBu)(PMe3)	HW-CVD	n.a.	180	26
	Silver carboxylates	PI-	Mesytilene	300
		MOCVD
	Hfod: 6,6,7,7,8,8,8-heptafluoro-2,2-dimethyl-3,5-octanedione BTMSA: bis(trimethylsiyl)acetylene; SEt2: Ethyl sulphide;
	VTES: vinyltriethylsilane; PEt3: triethylphosphine; P(OEt)3: triethylphosphite; PMe3: Trimethylphosphine; PPh3:
	Triphenylphosphine; O2CtBu: 2,2,6,6-tetramethyl-3,5-heptanedionate
	In this work, besides deposition tests performed with commercially available Ag precursors
	including AgAc, AgNO3 and Ag(TMHD), two new silver metalorganic precursors were used.

27

COD: 1,5-cyclooctadiene; H-hfac: 1,1,1,5,5,5-hexafluoro-2,4-pentanedione; H-tfac: 1,1,1-trifluoro-2,4-pentanedione;

The new precursors were designed and synthesized thanks to a collaboration with the group of Prof. Graziella Malandrino's in the University of Catania, Sicily. The new precursors, with general formula Ag(hfac)-L, were synthesized using different ligands: [Ag(hfac)(phen)] (III.1.

1 Synthesis and characterization of new silver metalorganic precursor Starting

materials: Ag2O (99%, Strem Chemicals), H-hfac (99%, Strem Chemicals), triglyme (99%, Sigma Aldrich), 1,10-phenanthroline (anhydrous 99%, Strem Chemicals), dichloromethane (99.8%, Sigma Aldrich), and n-pentane (99%, Sigma Aldrich) were used as received for the synthesis.

Table III -

 III 2.Single crystal crystallographic data and refinement parameters for AgP and AgT.

	Chapter 3. Thin film deposition of
	silver and silver oxide
		AgP	AgT
	Chemical formula	[Ag(hfa)(phen)]	[Ag(triglyme)2] + [Ag(hfa)2]
	Empirical formula	C17H9AgF6N2O2	C26H38Ag2F12O12
	Formula weight	495.13	986.30
	Temperature (K)	100	250
	Wavelength (Å)	1.54184	0.71073
	Crystal system, space group	Triclinic, P-1	Triclinic, P-1
	Unit cell dimensions (Å, °)	a = 7.4309(3) b = 9.2288(4)	a =11.4965(6) b = 12.9561(7)
		c = 13.1426(6)	c = 13.0344(9)
		α = 102.579(4)	α = 96.821(5)
		 = 105.172(4)	 = 91.795(5)
		γ = 103.608(4)	γ = 92.350(4)
	Volume (Å 3 )	807.10(6)	1924.8(2)
	Z, Dc (mg/cm 3 )	2, 2.037	2, 1.702
	 (mm -1 )	10.823	1.127
	F(000)	484	984
	Crystal size (mm)	0.30x0.27x0.20	0.28x0.25x0.19
	 range (°)	5.165 to 72.336	4.111 to 29.400
	Reflections collected / unique 8748 / 3045	13020 / 7438
	Data / parameters	3045 / 295	7438 / 519
	Goodness-of-fit on F 2	1.115	1.061
	Final R indices [I>2(I)]	R1 = 0.0407, wR2 = 0.0934	R1 = 0.0637, wR2 = 0.1714
	R indices (all data)	R1 = 0.0565, wR2 = 0.1169	R1 = 0.0876, wR2 = 0.1942
	-74 -

Table III -

 III

3. Selected bond distances (Å) and bond angles (°) for AgP and AgT.

Table III -

 III

4

.

1

H and

13

C NMR of AgP and AgT.

Table III -

 III 5. Ag coating deposition conditions and results summary.Even though Ag coating deposited using AgNO3 has been reported in previous works by N.Bahlawane et al.11 and S.D.Ponja et al.10 , in this work we were not able to obtain Ag coatings

	Chapter 3. Thin film deposition of
	silver and silver oxide
	Precursors tested AgNO3 AgAc	Ag(TMHD) AgP	AgT
	tested deposition	300	-	250 -400	250 -400	250 -350	250 -350
	temperature (°C)	450
	Need			N		Y	Y	N	N
	dispersant ? (In
	alcohol solvents)
	Deposition	of	N		Y	N	Y	Y
	Ag	in	PI-
	MOCVD ?
	Deposition	of	N		Y	N	Y	Y
	Ag in AA-
	MOCVD ?
	from AgNO3. For Ag(TMHD), different types of solvents were tested including M-Xylene,
	Toluene, butanol-01 and ethanol. With deposition temperatures between 250 and 400 °C, no
	-84 -

Ag coating was obtained. Thus, only results of Ag coatings by AgAc, AgP and AgT are discussed here.

  Table III-6. Both precursors were carefully grinded before being dissolved in ethanol, reaching 0.01 M concentration at room temperature.

	Chapter 3. Thin film deposition of
	silver and silver oxide
	-89 -

Table III -

 III 6. Deposition parameters (all reactions carried at 320 °C) using the new Ag precursors.

Table III -

 III 7. Electrochemical oxidation of Ag films with different durations, under 0.7 V, in 1 M NaOH solution

	Oxidation duration (s) 500	1000	1300
	Sample names	E01	E02	E03

  Furthermore, the crystal structures of both precursors were solved by means of single crystal X-ray diffraction.TGA analyses have shown that AgT is a precursor with a more rapid and clean decomposition than AgP. The new precursors present a good solubility in ethanol and have

	been tested in an Aerosol Assisted CVD system. Depositions by AA-MOCVD revealed that
	the two precursors, used under similar CVD parameters, produced coatings with different
	surface morphologies. Coatings deposited from AgP are made of Ag nanoparticles
	presenting a carbonaceous shell. Conversely, Ag coatings from AgT consist of pure Ag
	nanoparticles and present high crystallinity. In conclusion, our results show that the new
	AgT adduct is a good precursor for depositing metallic Ag coatings via AA-MOCVD using
	alcohol solvents.

using less expensive and toxic ligands, namely, [Ag(hfac)phen] (AgP) and [Ag(triglyme)2] + [Ag(hfac)2] -(AgT). FTIR spectroscopy has confirmed the formation of such molecular complexes.

Table IV -

 IV

2. Cu2O depositions with CuF precursor under different oxygen ratios for 1 hour.

  Table IV-3 summarizes the experimental conditions of each sample deposited with different quantity of Ag incorporated. Copper and silver precursors were mixed in the same alcohol solvent and co-deposited at the same time on Corning glass substrate.

Table IV -

 IV 3. Cu2O samples deposited with different quantity of Ag incorporated.

	Sample name	L0	L1	L2	L3	L4	L5
	Cu : Ag ratio in	Pure	10:0.625	10:0.833	10:1	10:1.25	10:2.5
	solution
	Ag ratio in film	None	Not detectable		0.3%	1%
	(atom ratio)
	Thickness	462	682	738	742	774	779
	(nm)

Table IV -

 IV 4. Sample deposition parameters and samples identification. Original water molecule in carrier gas set as constant zero.

	CG2b	Temperature	290	305	320	335	350	365
	(L/min)	(°C)
		Introduced
		H2O	molar
		ratio in total
		gas flow (%)
	0	0		D290	D305	D320	D-335	D-	D-
								350	365
	2	0.38					W02
	4	0.75					W04
	6	1.12					W06
	8	1.37		W290 W305 W320 W335	W-	W-
							(=	350	365
							W08)

  Taking advantage of the ASTAR technique, we were able to scan the cross-section of the Cu2O sample with ultra high resolution . As a result, tiny amounts of CuO phase were discovered in the grain boundaries, especially in the bottom part of the film, which helped to explain the electrical characteristics observed in the electrical AFM mapping. With optimized Cu2O (111) films, a p-n junction was also fabricated with structure of ITO/Cu2O/ZnO/Ag contact, and an excellent rectifying curve was obtained with on-off ratio exceeding 10 4 .

	Chapter V Deposition and characterization of AgCuO2 thin films	--
	V.1 Co-deposition of silver copper oxide via Plasma Enhanced AA-MOCVD	--
	V.2 Oxidation of Ag-Cu metal alloy film	--
	V.2.1 Deposition of Ag0.5Cu0.5 alloy films	--
	V.2.2 Electrochemical oxidation	--
	V.2.3 Oxygen plasma treatment	--
	V.3 Deposition of AgCuO2 thin films by Successive Ionic Layer Adsorption and Reaction
	(SILAR)	--
	V.3.1 Experimental setup and parameters	--
	V.3.2 Substrate preparation	--
	D. Cu2O/glass substrate preparation:	--
	E. ZnO/glass substrate preparation:	--
	F. TiO2/glass substrate preparation:	--
	V.3.3 SILAR deposition parameters	--
	V.3.4 Optimization of deposition condition	--
	V.3.5 Understanding the mechanism of AgCuO2 thin film coating by SILAR process.	-
	177 -
	V.3.6 Optical and electronic property of AgCuO2 films obtained by SILAR	--
	V.4 Chapter VI conclusions	--
	V.5 References	--

  7,9,10 With the working experiences on silver and copper oxides in MOCVD as introduced previously, this chapter will introduce the work on exploring new approaches to obtain better quality AgCuO2 thin films in terms of morphology and crystallinity. There are several thin film coating techniques that have been tested in this work, including initially planned codeposition of silver copper oxide film with Plasma Enhanced MOCVD. Other approaches explored in this work include oxygen plasma treatment of Ag-Cu alloy films deposited by MOCVD, electrochemical oxidation of Ag-Cu alloy films or by the solution method Successive Ionic Layer Adsorption and Reaction (SILAR). Finally, optical and electrical characterizations of AgCuO2 films are reported for the first time.

Table V -

 V

	Sample name Picture of sample Deposition	Plasma	power	Oxygen ratio
		temperature (°C)	(W)		in plasma gas
					source
	C1	270	350		50%
	C2	270	350		20%
	C3	270	350		10%
	C4	270	300		50%
	C5	250	0
	C6	270	0

1. Co-deposition details in Plasma Enhanced AA-MOCVD, with precursor ratio AgT: CuF = 1:1 for 1 h.

Table V -

 V 3. Electrochemical oxidation of AgCu film parameters.

	Sample names Electrolyte NaOH	Potential	for	Oxidation	Obtained phases
		concentration (M)	oxidation (V)		duration (min)
	EA01	0.2	0.5		17	Ag2O, Cu
	EA02		0.6			Ag, Cu
	EA03		0.7			Ag2O, Cu
	EA03b		0.7		34	Ag, Cu
	EB01	0.5	0.5		17	Ag2O, Cu
	EB02		0.6			Ag2O, Cu
	EB03		0.7			Ag2O, Cu
	EB03b		0.7		34	Ag2O, Cu
	EC01	1	0.5		17	Ag2O, Cu
	EC02		0.7			Ag2O, Cu

  successive SILAR coatings. Especially with Cu2O layer, the deposition rate was much faster than with any other substrate. Thus, the optimization of AgCuO2 film deposition was mainly conducted with Cu2O/glass substrate. The depositions on ZnO/glass by SILAR will be presented in the section dealing with reaction mechanism below.

	Chapter 5. Deposition and characterization
	of AgCuO2 thin films
	47 g (0.0025 mol) of Cu(NO3)2.3H2O
	in the mixture precursor solution. t1, t2, t3 and t4 were set to be 15 s, 10 s, 5 s and 10 s. Under
	this deposition conditions, different substrates including corning glass, Si substrate, TiO2
	-171 -

/glass, ZnO/glass and Cu2O/glass were tested with 70 deposition cycles. Due to a low adhesion between the metal ions and most of the mentioned substrates, the coated layers on surface tent to fall back in the solutions. Only ZnO/glass and Cu2O/glass could give

Table V -

 V 4. Parameters of AgCuO2 surface morphology optimization.

	Sample names	A0	A1	A2	A3
	Immersion	5	10	20	27
	duration in metal
	ion precursor t3 (s)

Note: Ag 1+ and Cu 2+ ion concentration were fixed to 0.05 M. t1, t2, t4 remained the same with values 10 s, 15 s, and 10 s.

Table V -

 V 5. Phase purity optimization by adjusting the concentration of copper ions in solution.The morphology of the samples deposited with different copper metal ion concentrations were characterized by SEM and the images are presented in Figure V-17. As it is shown, figure a) is the SEM image of the film deposited with equal ratio of copper and silver ion concentration in solution, presenting a dense surface with small grains (size about 100 nm).

	Sample names	B5	B4	B3	B2	B1
	Phase optimization: Ag 1+	0.028 M	0.03 M	0.035 M	0.04 M	0.05 M
	fixed at 0.05 M, Cu 2+
	precursor concentrations (M)

  The reaction mechanism for the deposition of AgCuO2 phase in SILAR process is discussed here. During the SILAR process, the substrate and the oxidizer played important roles. The oxidizer used in this work was Na2S2O8, which was added into the anion high pH NaOH solution. As it is reported in the work of D.Zhao et al.,27 Na2S2O8 is a very strong oxidant which can be activated through several means, thermally (over 60 °C), metal ion or by base activation. The oxidation of Ag 1+ ions with Na2S2O8 has been reported in the work of G.P.

Table VI -

 VI 2. in situ XRD of samples under different conditions.

	Sample names temperature	KOH	solution	Pressure	Obtained phase
		(°C)	concentration (M)	(bar)
	S1	90	1		250	AgCuO2
	S2		0.3			Cu4(SO4)(OH)6
	S3		0			Cu4(SO4)(OH)6 with
						intermediate phases
	S1b		1		1	AgCuO2

  2-ions and H2O, this complex was precipitated into Cu3(SO4)(OH)4, as reaction shown in Equation VI 4. As the reaction proceeded, thus more protons were consumed by AgO, the more supplying of Cu(OH) + complex, thus the Cu4(SO4)(OH)6 precipitation was observed, as the reaction shown in Equation VI 5. During the reaction, the water hydrolysis could even be accelerated by the applied 90 °C temperature, thus acceleration of the precipitation process. Meanwhile, since there was always excessive CuSO4 in solution, the formed intermediate phases, such as Cu3(SO4)(OH)4, possibly redissolved in acidic solution.

	Chapter 6. In-situ XRD characterization of formation
	mechanism of AgCuO2 phase in hydrothermal reaction
	𝐶𝑢 2+ + 𝐻 2 𝑂 ↔ 𝐶𝑢(𝑂𝐻) + + 𝐻 +	(28)
	𝑆𝑂 4 2-+ 𝐻 2 𝑂 ↔ 𝐻𝑆𝑂 4 -+ 𝑂𝐻 -	(29)
	𝐴𝑔𝑂 + 2𝐻 + → 𝐴𝑔 2+ + 𝐻 2 𝑂	(30)
	3𝐶𝑢(𝑂𝐻) + + 2𝐻 2 𝑂 + 2𝑆𝑂 4 2-→ 𝐶𝑢 3 (𝑆𝑂 4 )(𝑂𝐻) 4 ↓ +𝐻𝑆𝑂 4 -	(31)
	4𝐶𝑢(𝑂𝐻) + + 2𝐻 2 𝑂 + 3𝑆𝑂 4 2-→ 𝐶𝑢 4 (𝑆𝑂 4 )(𝑂𝐻) 6 ↓ +2𝐻𝑆𝑂 4 -
	-200 -

Chapter 2.Experimental setup

In this work, given that the final objective was to deposit AgCuO2 coatings by MOCVD,

Datablock AgP-ellipsoid plot:

General conclusion and perspectives

Cu2O thin film optimization in AA-MOCVD.

Acknowledgements

thanks to the Ministè re de education nationale, de l'project is thanked for supporting my research expenses. In addition, I am very appreciate for the KIC Innoenergy program. This thesis work is supported by Ecole Doctorale for Industrial, Materials, Mechanical and Process Engineering (I-MEP2) and Ministè re de l'Education nationale, de l'Enseignement supé rieur et de la Recherche in France. As well, KIC Innoenergy PhD program has provided financial support for the motilities during the PhD thesis. This PhD thesis work was supervised by Dr. David Muñoz-Rojas from Laboratoire des maté riaux et du gé nie

Datablock AgP-ellipsoid plot:

V.2 Oxidation of Ag-Cu metal alloy film

Due to the difficulty of co-depositing AgCuO2 with our AA-MOCVD system and the precursors available, a two-step process to achieve AgCuO2 was explored, as was done for Ag (see Chapter III). As it has been reported in the work of D.Muñoz-Rojas et al., 11,12 the AgCuO2 phase can be transformed through oxidation of Ag2Cu2O3 power either electrochemically or by ozone. Thanks to the high diffusion mobility of Ag atoms [13][14][15][16] , the oxidization of Ag0.5Cu0.5 alloy films was tested to obtain the AgCuO2 phase. Thus, in this work, Ag-Cu alloy films were firstly co-deposited by PI-MOCVD system (since less precursor is consumed than with AA-MOCVD) with an adjusted atomic ratio close to 1:1, and then further oxidized.

V.2.1 Deposition of Ag0.5Cu0.5 alloy films

The metallic alloy films were deposited at 320 °C with pressure 0.01 mbar and 1000 injection pulses (25 ml solution consumption) on corning glass substrates. During the deposition, the carrier gas for precursor transportation was 1 L/min Ar while the oxidizing gas was 1 L/min of oxygen. To obtain Ag0.5Cu0.5, the precursor solutions were initially prepared by mixing an equimolar amount of both precursors (0.0835 g of AgAc and 0.185 g of CuF, 0.5 mmol) in 25 ml ethanol solvent with 5 minutes of stirring. Again, to facilitate the precursor dissolution, 2.5 mmol of ethanolamine were added as dispersant. By varying the precursor concentration of the copper precursor solution, the element ratio in the films was tuned, as shown in Table V-2. The EDS analyses of the films were conducted with 10 kV voltage and the results are

The morphology of some of the samples is presented in Figure

Chapter VI In situ XRD characterization of formation mechanism of AgCuO 2 phase in hydrothermal reaction

Most of the previously reported synthesis techniques for AgCuO2 phase were through solution reaction routes. The starting materials were mostly stoichiometric mixture of AgNO3 and Cu(NO3)2 with K2S2O8 as oxidizer, in basic media. The starting materials were added into NaOH solution and reacted at 90 °C. [1][2][3][4] Under these reaction conditions, the coprecipitation between oxidized Ag 2+ and Cu 2+ were widely believed to be the formation mechanism of AgCuO2. Meanwhile, from the work of D.Muñoz-Rojas et al., 5 AgO and CuSO4 were also reported as starting material. With a hydrothermal reaction at 120 °C in 1 M KOH solution, the desired AgCuO2 phase was also obtained. Since AgO remains a solid phase during the reaction, this result implied that the reaction mechanism tacking place during the hydrothermal reaction might be different from the mechanism of previous coprecipitation approaches.

In Once the capillary reached the set point temperature (less than 10 s, temperature is calibrated previously by placing a thermocouple inside the capillary), the samples were exposed to the X-ray and the diffracted signals were collected by an Oxford Diffraction Titan CCD detector.

In order to eliminate the possible effect of light on the reaction process (Ag compounds being photosensitive), the whole preparation and characterization process was done in the dark.

The reflections from the sapphire capillary, shadow of the beam stopper and bad pixels were masked out from the raw data. Then the corrected 2D data were integrated into 1D by Fit2D software. 7 All initial calibrations were performed using a NIST 660a LaB6 standard.

VI.2 Effects of solution pH and reaction mechanism

The first factor we studied during the reaction was the pH, by modifying the amount of KOH Comparing with the standard 1 M KOH concentration, an increase of the background signal

Then, since AgCuO2 is more stable and thus less soluble in basic solution, 2,14 the coprecipitation between [Ag(OH)2] 2-and [Cu(OH)2] 2-happens, as shown in Equation VI 7.

Therefore, the silver and copper hydroxide complexes in solution are consumed by AgCuO2 phase precipitation, displacing the equilibrium between metal oxides and metal hydroxide complex towards the latter. Thus, AgO and CuO powders were consumed through such hydroxide complex co-precipitation and the AgCuO2 phase was formed. The rather constant background signal obtained in all patterns supports this view since a change in the total concentration of dissolved species would lead to a change in background.

The mechanism of the process is also illustrated in Figure VI-7. With this model, we can predict that with higher reaction temperatures, metal oxides can dissolve faster in the alkaline solution due to the equilibrium with metal hydroxide complex in the solution. Thus, the AgCuO2 precipitation speed can be largely enhanced. This was indeed the case, as it is shown in the next section.

Chapter 6. In-situ XRD characterization of formation mechanism of AgCuO2 phase in hydrothermal reaction -204 -

VI.3 Effects of temperature and calculation of the reaction activation energy

In addition to studying the effects of pH on the reaction outcome and mechanism, the activation energy was also estimated by recording the reaction duration at different temperatures. The solution pH was maintained at 1 M with 250 bar pressure and the reaction temperature was varied from 50 °C to 120 °C, as shown in Table VI-3. Since the slurry used had the same concentration, the quantity of injected reactants in the capillary tube were roughly the same. In all reaction temperatures, AgCuO2 was obtained as the final product. The full patterns of the in situ XRD characterizations are presented at the end of this chapter as appendix. In all cases a similar reaction evolution was observed, with the intensity of reflections corresponding to AgO, CuO decreasing and the reflections for AgCuO2 increasing.

To

VI.4 Chapter VI conclusions

In this chapter, the hydrothermal reactions of AgCuO2 phase synthesis were studied.

Choosing AgO, CuSO4 and KOH solution as precursors, the reactions were firstly conducted in the laboratory under 1 bar pressure with various KOH concentrations. AgCuO2 could only be obtained when using 1 M KOH solution, while in the other cases only Cu4(SO4)(OH)6 was obtained. To further investigate the effects of solution pH, similar samples were synthesized and followed by in situ XRD in synchrotron. By reviewing the evolution of the reflection intensity for each phase, the reasons for the absence of AgO phase were explained, and why only with enough high KOH concentration the AgCuO2 phase can be formed.

Besides the in situ XRD characterizations with different KOH concentration, the effects of reaction temperature were also studied with 1 M KOH solution and 250 bar pressure. The morphology of the grains synthesized at different temperatures was also studied. With a lower reaction temperature, averagely smaller grain size and anisotropic crystal growth are found. With increasing reaction temperature from 50 °C to 120 °C, the reactions were completed in shorter times. By extracting the reaction durations at each temperature, the Arrhenius equation was used and thus the reaction activation energy was estimated about 67.7 kJ/mol.

VI.5 Appendix V:

In situ XRD of hydrothermal reactions at different temperatures.

Conclusions and perspectives

In this thesis work, we have achieved four main results, including silver and its oxide thin film coatings, tuning the Cu2O films into (111) texture in AA-MOCVD, obtaining dense, continuous AgCuO2 thin films with unprecedentedly reported electrical and optical characterizations and studying the solution mechanism of AgCuO2 by in-situ XRD in ESRF synchrotron .

Deposition of silver and oxidation to silver oxide films.

Since our initial attempt was depositing AgCuO2 films in AA-MOCVD with the oxygen plasma coupling, the work began with silver metallic coatings. During this section of work, with the assistance of ethanolamine as dispersant, low cost silver precursor AgAc was firstly reported for silver coatings in MOCVD. However the silver coatings were not homogeneous due to the low volatility of AgAc. The commercialized silver metal organic precursors are rather expensive concerning the quantity of precursor consumption in AA-MOCVD.

Therefore, to obtain affordable silver metalorganic precursors with better volatility and deposition quality, a collaboration with Prof.Graziella Malandrino from University of Catania, Italy was stablished. Through this, two brand new silver precursors were reported, which were Ag(hfac)phenanthroline (AgP) and Ag(hfac)triglyme (AgT). Both precursors were able to delivery good quality silver coatings. Especially with AgT precursor, carbon free and homogeneous Ag coatings can be achieved. In addition, the two newly designed silver precursors have much lower ligands than commercial products.

After the Ag films were able to be deposited and optimized, the next step was trying to obtain silver oxide films. However, direct deposition of silver oxide films in MOCVD was always very challenging due to the incompatibility of the temperature between the precursors decomposition and the silver oxide phase. Therefore, an after-treatment approach was adopted, i.e. the first step was depositing Ag metallic films and then oxidizing them. To oxidize the silver films, we tested electrochemical oxidation and oxygen plasma treatment.

From both methods, AgO films were obtained. With oxygen plasma treatment, the oxidization duration (about 2 min) was much shorter than the electrochemical oxidation (at least 20 min). In addition, the plasma oxidized silver oxide films had less impurities and more homogeneous surfaces.

In addition to the bandgap estimation, the AgCuO2 films were also characterized by Hall Effect measurement. From the Hall Effect measurement, p type conductivity was confirmed.

The lowest resistivity obtained was 0.2 Ω.cm, which is even lower than other CuMO2 (M= Al, Cr, Sr, etc) delafossite compounds. In addition, the general low carrier concentration in those delafossite compounds was not observed in the SILAR coated AgCuO2 film, which is possibly due to the mixed valence band and charge delocalization in this compound. The maximum carrier mobility obtained was 24 cm 2 /V.s and the carrier concentration was in the order of 10 17 cm -3 ~ 10 18 cm -3 . Therefore, AgCuO2 is an excellent p type metal oxide with low resistivity. With its high carrier mobility and low bandgap, it presents the possibility of applications in both oxide transistor and photovoltaic devices. given and thus the activation energy was estimated about 67.7 kJ/mol.

Perspectives in future work

1. Co-deposition of AgCuO2 in MOCVD/ALD is still doable provided low decomposition temperature precursors. In addition, to avoid the corrosive radicals, ozone generator is suggested for providing oxidizing species instead of oxygen plasma system.

2. Furthermore works are required to study the electronic structure of the AgCuO2 phase.

Especially in order to form a p-n junction for device applications, it's important to know the electron affinity thus to choose a proper n type metal oxide material.

3. XPS is also needed to compare the charge states of silver atom in this deposited thin films and compare it with the AgCuO2 phase synthesized through other approaches in literatures.

4. When all the electronic structure information are obtained for AgCuO2 thin films, they should be applied into transistors or photovoltaic devices.