XI

transport très spécifiques, qui ouvrent de nombreuses perspectives pour comprendre la physique de l'interaction électron-phonon et qui fournissent également de nouveaux paradigmes pour des applications en photonique, électronique, ou en récupération de l'énergie solaire.

Résumé

I

Ré sumé

L'intégration de semi-conducteurs III-V sur silicium (Si) combinant les avantages du Si (abondance terrestre, faible coût et prévalence dans les industries électroniques et photovoltaïques) avec les excellentes propriétés optoélectroniques des III-V, est d'un grand intérêt dans les applications en photonique et récupération de l'énergie solaire [1,2].

Comme exposé dans le Chapitre 1, L'approche monolithique comme méthode d'intégration III-V/Si offre potentiellement d'énormes avantages, notamment des fonctionnalités polyvalentes, des coûts de production réduits sur de grandes surfaces de substrat et une intégration dense. Cependant, les différences fondamentales des propriétés physiques entre les matériaux III-V et Si, comme le paramètre de maille, le coefficient de dilatation thermique ou encore la polarité, introduisent une densité élevée de divers défauts lors de la croissance de semi-conducteurs III-V sur Si.

Parmi ces défauts, les parois d'antiphase (APB), constituées de liaisons chimiques accidentelles III-III ou V-V, sont particulièrement emblématiques de la croissance III-V/Si. Comme le montrent les Fig. 1a et b, une APB dite « stoechiométrique » a un nombre égal de liaisons III-III et V-V (Fig. 1a), alors qu'une APB « non stoechiométrique» n'est composée que de liaisons III-III ou de liaisons V-V (Fig. 1b). Des recherches récentes sur les mécanismes de génération, de propagation et d'enfouissement des APB [3,4] donnent des clés pour contrôler à volonté leur distribution dans les échantillons III-V/Si. Cependant, peu de travaux se sont concentrés spécifiquement sur l'étude de leurs propriétés physiques [5]. Cette thèse se concentre sur l'étude des propriétés optoélectroniques, vibrationnelles, et de transport des APB dans les matériaux III-V/Si, et leur utilisation pour les dispositifs de récupération d'énergie solaire, d'un point de vue théorique et expérimental. Les techniques de caractérisations structurales, optiques et de transport, ainsi que les méthodes de calcul par la théorie de la fonctionnelle de la densité (DFT) utilisées dans cette thèse sont exposées dans le Chapitre 2. En outre, nous avons effectué des calculs de structure de bande et de phonons pour les APB stoechiométriques basés sur la théorie de la fonctionnelle de la densité (DFT).

Nous constatons que les atomes au niveau des APB introduisent non seulement plusieurs états électroniques localisés sur le dessus de la bande de valence (Fig. 3a), réalisant un confinement quantique des porteurs, mais introduisent également 5 modes de phonons localisés et plats dans la bande interdite de dispersion des phonons (Fig. 3b), permettant le confinement des phonons. L'impact de l'interaction électron-phonon sur les processus de photoluminescence est ensuite clarifié en combinant les résultats du calcul DFT avec des expériences de spectroscopie Raman et de photoluminescence. Le spectre de Photoluminescence (PL) contient plusieurs répliques de phonons et un facteur Huang-Rhys de 8 est finalement déterminé pour la raie d'émission liée aux APB (Fig. 4), soulignant qu'un couplage électron-phonon important et inhabituellement fort peut être obtenu en raison du confinement simultané des porteurs de charge et des phonons au voisinage des structures Résumé IV APB stoechiométriques 2D (Fig. 5). Au Chapitre 4, une nouvelle association de matériaux en tandem GaPSb/Si (Fig. 6) a été étudiée pour les applications en craquage photo-électro-chimique (PEC) de l'eau à partir de l'énergie solaire [7,8]. L'évolution de la structure de bande des alliages GaPSb est déterminée pour l'ensemble des compositions en Sb en combinant des mesures d'absorption avec des calculs de liaisons fortes, comme le montre la Figure 7a. Le croisement indirect (bande L) vers direct (bande Γ) se produit pour une composition en Résumé V Sb de 30%. En particulier, à 32% de Sb, l'alliage GaP1-xSbx atteint une configuration de gap direct souhaitée de 1.7 eV, permettant une absorption efficace de la lumière du soleil.

En outre, l'alignement de bande de GaP1-xSbx/Si par rapport aux niveaux de potentiel redox de l'eau a été analysé (Fig. 7b), ce qui montre que l'association 1.7eV/1.1eV

GaPSb/Si est une combinaison intéressante à la fois pour les réactions de production d'hydrogène et de production d'oxygène.

VI

En outre, une photoanode GaP0.67Sb0.33/Si avec une combinaison de bande interdite très proche de l'optimum (1.7eV/1.1eV) a été étudiée pour le craquage PEC de l'eau par l'énergie solaire. Elle montre de bonnes performances en raison de la forte absorption de la lumière, notamment avec des couches de protection (TiO2) et le co-catalyseur (Ni) (Figure 8), ce qui indique que ce matériau en configuration tandem est très prometteur pour le craquage de l'eau à haut rendement et à faible coût. En conclusion, les résultats de recherche présentés dans cette thèse démontrent que les parois d'antiphase III-V/Si ont des propriétés optoélectroniques, vibrationnelles et de 1

General introduction

Integration of III-V semiconductors on silicon (Si), combining the excellent optoelectronic properties of III-V semiconductor with the benefits of silicon (such as low cost, abundance, maturity of manufacturing technologies), is of great interest for integrated photonics and energy harvesting (including solar cells and solar water splitting)

applications [1,2]. Especially the monolithic integration of III-V semiconductors on Si offers great potential advantages, including versatile functionalities, low-cost and largearea production, and dense integration. However, the material dissimilarities between III-V and Si, such as lattice constant, thermal expansion coefficient, and polarity, introduces a high density of various crystal irregularities during the growth of III-V on Si, that were considered for years as detrimental defects, without a clear assessment of their physical properties. However, for the development of highly-integrated devices, the specific impacts of these crystal irregularities on the optoelectronic properties of III-V/Si have to be clarified.

Especially, very specific crystal defects, known as AntiPhase Boundaries (APBs), are very specific to the III-V/Si monolithic integration and have been considered as "device killers" for a long time. The recent researches on the generation, propagation and burying mechanisms of APBs [3,4] gave keys to tune at will their distribution in III-V/Si samples. Nevertheless, only few works studied specifically the optoelectronic properties of APBs. This thesis not only demonstrates that APBs are nanostructures with unique physical properties, but also shows how it can be used for energy harvesting applications.

Therefore, this work, combining experiments and theory, focuses on the study of the specific optoelectronic, vibrational and transport properties of APBs in III-V/Si materials, and their use for energy harvesting devices.

Dissertation overview:

Chapter 1 introduces the research background of the monolithic integration of III-V semiconductors on Si. A brief state-of-the-art of the different III-V/Si integration approaches for photonics and solar energy harvesting is presented. Then, a focus is given on the issues raised during the crystal growth of III-V semiconductors on silicon substrate.

Finally, a summary of the recent research developments about the generation, propagation and burying mechanisms of APBs, and their known properties is proposed, where it can be seen that their physical properties were almost not explored in previous works.

Chapter 2 presents the experimental and theoretical supports of this work. The different experimental setups (structural, optical or photo-electro-chemical (PEC) characterizations) are described, with the related mechanisms, techniques and parameters settings used in this thesis. For the theoretical calculations, we firstly give a brief introduction to the basic concepts of the DFT, then the calculation steps and parameter settings for the determination of electronic band structure, and of the vibrational and electron-phonon coupling properties are presented.

Chapter 3 focuses on the study of the electron-phonon coupling around stoichiometric APBs in an InGaP/SiGe/Si sample. 2D vertical In-rich stoichiometric APB singularities are investigated firstly with X-Ray Diffraction (XRD), scanning transmission electron microscopy (STEM) and energy-dispersive x-ray spectroscopy (EDX) structural characterizations. The impact of the electron-phonon interaction on the photoluminescence processes is then clarified by combining Raman and photoluminescence experiments with electronic band and phonon DFT calculations. The strength and origin of the observed coupling is then deeply discussed, with regards to the APB symmetry properties.

Chapter 4 aims to assess the potential of a novel GaPSb/Si tandem materials association for high-efficient low-cost solar water splitting, based on bandgap engineering.

Thus, the quality of samples is investigated firstly based on X-ray diffraction (XRD), scanning electron microscopy (SEM), and atomic force microscopy (AFM) structural characterizations. Then, the bandgap of GaPSb alloys is determined on the whole Sb range by absorption measurements and tight-binding calculations. Based on the bandgap and band alignments determined, the potential of GaPSb/Si material for water splitting is then discussed. Finally, the PEC performances of a GaP0.67Sb0.33/Si photoanode are investigated and discussed in terms of solar spectrum absorption, while transport properties in these photoanodes are questioned.

Chapter 5 demonstrates that epitaxial III-V/Si samples with vertical nonstoichiometric APBs are hybrid materials composed of bulk semiconductors and 2D topological semi metallic inclusions. Based on structural, transport, PEC experiments and theoretical calculations, the photo-activity, ambipolar transport and PEC performances are investigated for the epitaxial III-V/Si samples with non-stoichiometric APBs.

Silicon properties

Silicon was named by Thomas Thomson in 1817 [5] and was first isolated by Swedish chemist Berzelius in 1825 [6]. Pure silicon is a dark gray solid with the same crystalline structure as diamond, which consists of two face-centered cubic (fcc) sublattices with one sub-lattice shifted one quarter of the body-diagonal with respect to the other sub-lattice (as shown in Fig. 1.1), called "two interpenetrating face-centered cubic" primitive lattices by Ashcroft and Mermin [7]. Silicon has several great advantages. It is the seventh-most abundant element in the universe and the second-most abundant element on earth, after the oxygen. Silicon is naturally found in the form of compounds with oxygen, aluminum and magnesium but can easily be synthesized with a high purity at a low cost. Silicon has also a high thermal 6 conductivity, which is important for electronic devices, because as the current density and working time increase, more heat is generated, causing device operating temperature to rise and thus leading potentially to device degradation. Moreover, the electrical properties of silicon can be finely tuned by implantation processes, widely used during the doping steps within the fabrication of MOS (metal-oxide-semiconductor) transistors. For these reasons, silicon is today the ruling material in micro-electronics and photovoltaics (PV).

On the other hand, optical properties of silicon remain fundamentally limited. Indeed, Si has an indirect bandgap, with the minimum of the conduction band in the X valley of the reciprocal space and the maximum of the valence band in the Γ valley (see Fig. 1.2) [8]. Thus, a radiative recombination between an electron and a hole requires a phononassisted transition in order to satisfy the momentum conservation. The poor light emission efficiency of Si is thus directly related to the low probability of an electron to meet a phonon before it meets a non-radiative center. Besides, for photovoltaic applications, crystalline silicon is also not the most efficient material due to its indirect bandgap and consequently to slow recombinations, long charge carrier lifetimes and low absorption properties. It is usually compensated by using thick Si absorbers. The crystal structure of most III-V semiconductors consists in two face-centered cubic (fcc) sub-lattices, composed of two different atoms: one sub-lattice is occupied by Group-III atoms, while the other one is occupied by Group-V atoms, leading to the socalled Zinc-Blende structure (as shown in Fig. 1.3). Consequently, III-V materials are polar semiconductors, while Si is a non-polar semiconductor [START_REF] Holt | Extended Defects in Semiconductors: Electronic Properties, Device Effects and Structures[END_REF]. Most of the III-V semiconductors present a direct bandgap (Fig. 1.2 left) and are thus naturally suited for applications in photonics or energy harvesting [START_REF] Tanabe | A Review of Ultrahigh Efficiency III-V Semiconductor Compound Solar Cells: Multijunction Tandem, Lower Dimensional, Photonic Up/Down Conversion and Plasmonic Nanometallic Structures[END_REF]. They typically provide efficient light emission and high photon absorption because of their direct bandgaps (e.g. for InP, GaAs, InAs or GaSb), high electron mobilities (e.g. for InAs, InP and GaAs). They can also be alloyed, which gives the opportunities to tune at will their bandgaps and bandlineups, from the infrared range to the ultraviolet one (this will be discussed in the next section). On the other hand, the largest disadvantage of III-V semiconductors is their high cost. Unlike silicon wafers, which are produced at a large scale worldwide, with extraordinary purity over very large wafers, III-V crystals are produced in much smaller volumes and are far more expensive. 200 mm III-V wafers can cost as much as $5,000 (This also depends on the type of III-V used) while 300 mm silicon wafers cost less than $100 [START_REF]Why GaAs and III-Vs Probably Won't Replace Silicon | by Tapabrata Ghosh | Medium[END_REF]. Historically, cost has been one of the major barriers against the massive development of III-Vs.

Bandgap engineering through alloying

Bandgap engineering is the process of controlling or altering the band gap of a material. This is typically done with semiconductors by controlling the composition of alloys or constructing layered materials with alternated compositions [START_REF]Band-gap engineering[END_REF]. Bandgap engineering of III-V semiconductors materials started 30 years ago and today many existing commercial devices (including LEDs, lasers and so on) use this approach. Based on bandgap engineering, both the bandgap and band type of III-V semiconductors can be adjusted by adding other elements (one or two typically) to achieve ternary or quaternary alloys (pentanary alloys are most of the time too difficult to control to be implemented in devices). In many semiconductor alloys, a linear interpolation between the respective lattice parameters at the two end points may be used to determine the lattice parameter of the alloy. It is a function of the composition, often referred as the Vegard's law. However, the bandgap of most of the semiconductors does not vary linearly with the composition and a deviation is observed, which is known as the bandgap bowing [START_REF] Kour | Review-Recent Advances and Challenges in Indium Gallium Nitride (InxGa1-xN) Materials for Solid State Lighting[END_REF]. To account for this deviation, a correction term known as bowing parameter b is added. This bowing parameter b is needed to determine the bandgap of semiconductor alloys [START_REF] Vegard | Die Konstitution der Mischkristalle und die Raumfüllung der Atome[END_REF]. The composition dependence of bandgaps of ternary alloys AxB1-xC with two kinds of cations (or CAxB1-x with two kinds of anions) at room temperature can be determined by using the standard formula [START_REF] Mahala | p-GaN/i-InxGa1-xN/n-GaN solar cell with indium compositional grading[END_REF]:

𝐸 𝑔 (𝑥) = 𝐸 𝑔 (𝐴𝐶) * 𝑥 + 𝐸 𝑔 (𝐵𝐶) * (1 -𝑥) -𝑏 * 𝑥(1 -𝑥) (1-1)
Where 𝐸 𝑔 is bandgap of the ternary alloy AxB1-xC (or CAxB1-x), 𝐸 𝑔 (AC) is the bandgap of the binary alloy AC, 𝐸 𝑔 (BC) is the bandgap of the binary alloy BC, b is the bowing parameter, x is the concentration of the A element and (1-x) is the concentration of the B element.

Bandgap engineering was investigated for many ternary or quarternary III-V alloys.

Figure 1. 4 shows the bandgap of some important binary and ternary III-V semiconductors as a function of the lattice parameter. The bandgaps of the ternaries are given at their minimum, which can correspond to the gaps at the , X or L valleys, from the respective band gap bowing parameters provided in ref. [START_REF] Vurgaftman | Band parameters for III-V compound semiconductors and their alloys[END_REF][START_REF] Vurgaftman | Band parameters for nitrogen-containing semiconductors[END_REF], based on Eq.1-1.

General interest of III-V and Si co-integration 1.1.3.1 Integrated photonics

The Si-based complementary metal-oxide-semiconductor (CMOS) technology is the fundamental building block of the contemporary electronics [START_REF] Thompson | s law: the future of Si microelectronics[END_REF][START_REF] Park | Heteroepitaxial Growth of III-V Semiconductors on Silicon[END_REF]. However, the electrons travelling through copper wires is now a bandwidth limitation for the most advanced MOS transistors. In contrast, photons move at light speed and avoid the bottleneck of metallic connections. Besides, electronic current generates non-negligible heat through the Joule effect, which is harmful to electronic devices. Photons can transmit great amounts of information, releasing only a fraction of the energy they carry [START_REF]What is Integrated Photonics? | Knowledge and Innovation for Manufacturing Initiative[END_REF].

10

Therefore, Si photonics integrating optical systems at the inter-chip or even intra-chip scale, has emerged as a highly attractive research area for ultra-fast and low-cost optical data transmission on the Si-on-insulator (SOI) platform compatible both with the CMOS process [START_REF] Park | Heteroepitaxial Growth of III-V Semiconductors on Silicon[END_REF] and photonic architectures requirements.

However, as mentioned above, the bulk Si material has only limited optical properties. Although Si-based light-emitting devices, like Si Raman laser [START_REF] Rong | A continuous-wave Raman silicon laser[END_REF] and p-n junction emitter [START_REF] Ng | An efficient room-temperature silicon-based light-emitting diode[END_REF][START_REF] Pillai | Enhanced emission from Si-based light-emitting diodes using surface plasmons[END_REF] were developed, their performances are still far from the one of III-V devices [2,[START_REF] Park | Heteroepitaxial Growth of III-V Semiconductors on Silicon[END_REF]. Indeed, III-V semiconductors have been widely used in lightemitting/absorbing devices, including light-emitting diodes (LEDs), lasers, detectors due to their superior optoelectronic properties. Thus, integration of III-V on Si leveraging the benefits of both Si and III-V semiconductors has brought great deal of interest and excitement for integrated photonics (Fig. 1.5). power applications [START_REF] Jain | III-V Multijunction Solar Cell Integration with Silicon: Present Status, Challenges and Future Outlook[END_REF]. III-V solar cells have steadily shown performance improvement 11 at ~1% increase in efficiency every year. The recent world record efficiency of 44.7% at 297 suns for a four-junction III-V solar cell was achieved [28]. Nevertheless, their expensive cost is the biggest impediment in their large-scale deployment for terrestrial applications. On the other hand, Si has been one of the dominant prime materials for photovoltaic technology since the 1950s [START_REF] Shah | Photovoltaic Technology: The Case for Thin-Film Solar Cells[END_REF]. But the performance of single-junction Si solar cells has almost saturated at around 25%. The domination of silicon solar cells is now facing the strong competition of other efficient technologies, such as perovskites solar cells [START_REF] Tsai | High-efficiency two-dimensional Ruddlesden-Popper perovskite solar cells[END_REF].

The cost of the substrate significantly contributes to the high cost of III-V solar cells.

Typically, GaAs or Ge substrates used for III-V multijunction solar cell growth are not only smaller in diameter but are also greatly more expensive than Si [START_REF] Jain | III-V Multijunction Solar Cell Integration with Silicon: Present Status, Challenges and Future Outlook[END_REF]. Besides, Si also offers higher thermal conductivity and superior mechanical strength in comparison to GaAs or Ge substrates [START_REF] Jain | III-V Multijunction Solar Cell Integration with Silicon: Present Status, Challenges and Future Outlook[END_REF]. Integration of III-V multijunction solar cell on Si substrate, where the Si can potentially be used as an active bottom sub-cell or just as an inactive low cost starting template, provides a promising path for low-cost and high-efficiency photovoltaic devices [1].

(b) solar water splitting

Solar water splitting converting solar energy to hydrogen through photoelectrochemical (PEC) cell, provides a promising path for scalable and sustainable carbon-free hydrogen production [START_REF] Bard | Artificial photosynthesis: solar splitting of water to hydrogen and oxygen[END_REF][START_REF] Lewis | Powering the planet: Chemical challenges in solar energy utilization[END_REF]. It is one significant milestone on the road to a sustainable energy future, especially under the current background of global energy and environmental crisis. The development of the PEC water splitting technology (as shown in Fig. 1.6), where the sunlight falling on an immersed semiconductor photoelectrode, turns the liquid water into gaseous storable hydrogen that can be reused on demand for heat or electricity production, has driven many researches in the past years.

Within the same way as for photovoltaic applications, the development of efficient 12 semiconductor photoelectrodes for solar water splitting not only requires excellent optical properties to enable high light absorption but also the use of low-cost materials for largescale deployment and development. Besides, with the flexibility offered by bandgap engineering strategies with III-V semiconductors, not only the bandgap can be adjusted to obtain high light absorption, but also the band alignment with water redox potential levels can be optimized. Therefore, the III-V/Si heterogeneous association of materials is also highly attractive and interesting for low-cost high efficiency solar water splitting applications. At the beginning of this work (2017), the only existing III-V/Si photoelectrodes were made of III-V Nano-wires, with complex designs and/or limited performances. We will see later (Chapter 4) that the two first monolithic III-V/Si photoelectrodes were demonstrated in 2019 simultaneously by our laboratory (with Pbased photoelectrodes) [START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF] and the one of Pr. Nötzel (with N-based photoelectrodes) [START_REF] Kumar | Quantum dot activated indium gallium nitride on silicon as photoanode for solar hydrogen generation[END_REF].

III-V/Si integration methods

Plenty of strategies have been tried to co-integrate III-V semiconductors and silicon for optoelectronic applications [START_REF] Yamaguchi | Analysis of strained-layer superlattice effects on dislocation density reduction in GaAs on Si substrates[END_REF]. Two main integration methods can be considered: hybrid integration and monolithic integration, as shown in Fig. 1.7. The hybrid one includes "sticking" or "bonding" of optical layers or devices directly on the silicon chip.

The bonding process can be achieved by direct bonding (molecular-assisted), adhesive 13 bonding (polymer-assisted) or metal-assisted bonding [START_REF] Lucci | Surface and interface contributions to III-V/Si hetero-epitaxial growth: Theory and Experiments[END_REF]. The monolithic integration consists in growing the III-V semiconductors directly on the Si substrate. This approach has different advantages with respect to the hybrid one, such as low fabrication costs and time. It should also ease the very large-scale integration of III-V devices on Si, which can hardly be achieved through the hybrid integration. In the following sections, more information about the monolithic integration method will be provided.

Monolithic integration of III-V on Si

As discussed above, the monolithic integration approach offers many promises.

Nevertheless, its primary challenge is to reach a high III-V/Si interface crystal quality and manage the crystal defects which may propagate in the volume of the device and degrade its performances and stability.

The lattice mismatch between the III-V and Si is one main parameter determining the quality of the epitaxial III-V/Si materials. Fig. 1.4 displays the bandgap energies of different III-Vs as a function of their lattice parameter. On the one hand, it reveals that some III-V semiconductors (such as GaAs, GaSb or InP) have a direct bandgap but a large lattice mismatch with the silicon. This will cause the formation of defects (e.g. misfit dislocations) and decrease the photonic devices efficiency. On the other hand, III-Vs such as AlP or GaP are quasi-lattice-matched to the silicon but, they have an indirect bandgap which can limit as well the device performances. When growing lattice mismatched III-

14

Vs on Si substrate, the integration is called metamorphic, otherwise it is called pseudomorphic. The approaches and the strategies, which were used for overcoming the related growth issues, will be presented in the following parts.

Metamorphic integration

For most applications, the main advantages of III-V semiconductors are their direct bandgap, such as GaAs or InP. However, they also present a large lattice mismatch to the silicon. As a result, the misfit dislocations formation is unavoidable. Crystal defects are usually considered as non-radiative recombination centers in the materials (we will see afterwards that it is not always the case). Depending on the targeted application, the density of these defects has to be reduced at a given value to ensure both efficient transport and adapted optical properties. For instance, considering laser devices, it is commonly accepted that the density of emerging dislocations should lie below 10 6 cm -2 to guarantee reasonable laser performances and stability, and should even reach 10 4 cm -2 for high performances III-V lasers. Fig. 1.8 shows the propagation of dislocations in a GaAs on Si epilayer [START_REF] Luxmoore | III-V quantum light source and cavity-QED on Silicon[END_REF]. A large number of approaches have been tried to grow high quality III-V epitaxial layers on Si substrate with low density of dislocations to improve the device performances. One can for instance benefit from the specific crystal relaxation processes of Sbbased alloys which have a very large lattice mismatch with Si (e.g. 12% for GaSb and 13%

for AlSb). The strain energy due to the large mismatch could be relieved into the III-V epilayer through the formation of self-organized interfacial misfit dislocations (IMD) arrays [START_REF] Huang | Simultaneous interfacial misfit array formation and antiphase domain suppression on miscut silicon substrate[END_REF][START_REF] Jallipalli | Atomistic modeling of strain distribution in self-assembled interfacial misfit dislocation (IMF) arrays in highly mismatched III-V semiconductor materials[END_REF]. As shown in Fig. 1.9, the formation of planar array of 90° misfit dislocations which are partly confined at the interface, could be induced when AlSb is grown on Si substrate. The low threading dislocation density obtained with this approach highlights the potential of this strategy for the monolithic integration. Kim et al. [START_REF] Kim | Growth mode and structural characterization of GaSb on Si (001) substrate: A transmission electron microscopy study[END_REF] and Akahane et al. [START_REF] Akahane | Initial growth stage of GaSb on Si(001) substrates with AlSb initiation layers[END_REF] improved the quality of the GaSb/Si crystal interface via an AlSb nucleation layer. GaSb-based devices have also been realized in the literatures [START_REF] Reboul | Continuous-wave operation above room temperature of GaSb-based laser diodes grown on Si[END_REF][START_REF] Monge-Bartolome | Etched-cavity GaSb laser diodes on a MOVPE GaSb-on-Si template[END_REF][START_REF] Tournet | GaSb-based solar cells for multi-junction integration on Si substrates[END_REF][START_REF] Tournié | Antimonide-based optoelectronic devices grown on Si substrates[END_REF][START_REF] Tournie | Epitaxial Integration of GaSb-based mid-IR devices on Silicon[END_REF]. Another effective strategy to reduce the threading dislocations is the use of strained layer superlattice (SLS) and short period superlattice (SPL) [START_REF] Yamaguchi | Analysis of strained-layer superlattice effects on dislocation density reduction in GaAs on Si substrates[END_REF]. This approach has been presented by growing high quality InGaAs quantum-dots (QDs) on Si substrate and by growing InP-on-Si heterostructures [START_REF] Luxmoore | III-V quantum light source and cavity-QED on Silicon[END_REF][START_REF] Samonji | Reduction of threading dislocation density in InP-on-Si heteroepitaxy with strained short-period superlattices[END_REF]. An example is shown in Fig. 1.10.

Figure 1.10 Cross-sectional TEM images, displaying the suppression of threading dislocations through the growth of strained superlattice (SPS) and short period superlattices (SPL). Courtesy of [START_REF] Luxmoore | III-V quantum light source and cavity-QED on Silicon[END_REF] The adaptation of the lattice parameter by compositionally graded buffers is also a widely investigated approach, for instance SixGe(1-x)/Si pseudo substrate [START_REF] Fitzgerald | Totally relaxed GexSi1-x layers with low threading dislocation densities grown on Si substrates[END_REF] to reduce the dislocation density at the GaAs/Si interface. High performance GaAs-based solar cells [START_REF] Andre | Investigations of high-performance GaAs solar cells grown on Ge-Si/sub 1-x/Ge/sub x/-Si substrates[END_REF] and laser structures [START_REF] Groenert | Monolithic integration of room-temperature cw GaAs/AlGaAs lasers on Si substrates via relaxed graded GeSi buffer layers[END_REF] have been realized. TEM cross-section micrograph of the GaAs/AlGaAs laser structure on Si substrate with relaxed graded SiGe buffer layers is shown in Fig. 1.11, where the dislocations are suppressed effectively. This strategy was also used to grow lattice mismatched GaAs0.7P0.3 solar cells on Si substrate [START_REF] Geisz | Lattice-mismatched GaAsP Solar Cells Grown on Silicon by OMVPE[END_REF]. Besides, the growth of III-V nanowires on Si is another widely used strategy to mitigate the appearance of defects [START_REF] Svensson | Monolithic GaAs/InGaP nanowire light emitting diodes on silicon[END_REF]. It includes growing III-V core-shell nanowires through catalyst-assisted (for instance with the help of Au nanoparticles) or catalyst-free processes [START_REF] Koblmüller | Growth and properties of InGaAs nanowires on silicon[END_REF][START_REF] Mårtensson | Epitaxial III-V Nanowires on Silicon[END_REF]. Fig. 1.12 shows monolithic GaAs/InGaP nanowires light emitting diodes on silicon with Au nanoparticle catalyst.

Pseudomorphic integration

The pseudomorphic integration consists in growing III-V semiconductors lattice matched to the Si substrate, to avoid the formation of misfit dislocations. The low lattice mismatch allows avoiding the formation of misfit dislocations between GaP or AlP and the silicon surface until a critical thickness of around 100 nm. The GaP is one of the best candidates for this approach. Indeed, the lattice mismatch of GaP and Si is only 0.37% at room temperature. Thus, the epitaxial growth of quasi-lattice matched GaP on silicon is considered to be an efficient platform for the subsequent integration of low defect or defect-free III-V based heterostructures. Fig. 1. [START_REF]Band-gap engineering[END_REF] shows good atoms alignments [START_REF] Wang | Abrupt GaP/Si hetero-interface using bistepped Si buffer[END_REF] due to its very low lattice mismatch to the Si. It is possible to avoid defects formation such as microtwins (MTs) and to annihilate antiphase domains [START_REF] Wang | Quantitative evaluation of microtwins and antiphase defects in GaP/Si nanolayers for a III-V photonics platform on silicon using a laboratory X-ray diffraction setup[END_REF] by choosing the good growth conditions. As shown in Fig. 1.14, most of the antiphase boundaries (APBs) could be annihilated within 10 nm from the GaP/Si interface. Takagi et al. [START_REF] Takagi | Generation and suppression process of crystalline defects in GaP layers grown on misoriented Si(100) substrates[END_REF] reported that it is possible to reduce the defects formation and improve the surface roughness through migration enhanced epitaxy (MEE). Volz et al. [START_REF] Volz | GaP-nucleation on exact Si (001) substrates for III/V device integration[END_REF] were able to grow a defect-free GaP/Si(001) after almost 50 nm GaP layer thickness. A two-step growth process was developed by Grassman et al. by using low temperature MEE, and then followed by molecular beam epitaxy (MBE) [START_REF] Grassman | Control and elimination of nucleation-related defects in GaP/Si(001) heteroepitaxy[END_REF]. In this way, a very smooth and defect-free GaP/Si interface was realized. However, the indirect nature of its bandgap has imposed severe limitations on the performances of GaP-based laser devices. Based on bandgap engineering, lasing on silicon was even demonstrated with GaAsPN quantum wells [START_REF] Liebich | Laser operation of Ga(NAsP) lattice-matched to (001) silicon substrate[END_REF].

With the incorporation of Nitrogen (N), the GaP indirect bandgap can be converted into a pseudo direct one for overcoming the indirect band structure issue. Also, some GaP/Sibased photovoltaic devices were achieved, either considering GaP/Si heterostructures or by developing tandem-like architectures [START_REF] Cornet | Chapter 28 -GaP/Si-Based Photovoltaic Devices Grown by Molecular Beam Epitaxy[END_REF]. More recently, our group developed high quality phase-controlled GaP/Si micro-resonator for non-linear integrated photonics [START_REF] Saleem-Urothodi | Loss assessment in random crystal polarity gallium phosphide microdisks grown on silicon[END_REF].

Issues of monolithic integration of III-V semiconductors on Si

The quality of III-V/Si materials is crucial for the performances of the developed devices. In the monolithic integration context, in order to obtain a high crystal quality III-V/Si interface and III-V layers, different materials growth challenges have to be taken into account, including the Si surface preparation, the lattice mismatch and thermal expansion coefficient, the 3D-growth or the issues related to the polar on non-polar epitaxy [START_REF] Lucci | Surface and interface contributions to III-V/Si hetero-epitaxial growth: Theory and Experiments[END_REF]. In the following, main issues related to the III-V/Si monolithic integration directly connected to this work, are briefly exposed.

Lattice mismatch and thermal expansion coefficient

The lattice mismatch and thermal expansion coefficient strongly influence the quality of the monolithic integration of III-V semiconductors on Si. The lattice mismatch is often used to quantify the difference of lattice parameters between the epitaxial layers deposited and the substrate. The lattice mismatch f is defined by the following equation [START_REF] Ayers | Heteroepitaxy of Semiconductors: Theory, Growth, and Characterization[END_REF]:

𝑓 = 𝑎 𝑠 -𝑎 𝑓 𝑎 𝑠 (1-2)
Where 𝑎 𝑓 and 𝑎 𝑠 are the lattice parameters of the semiconductor epitaxial film and substrate, respectively.

The generated strain can be accommodated either through elasticity or through plasticity. In the first case (usually associated to small lattice mismatch, and thin epilayers), the elastic energy is mostly stored through a biaxial strain of the epilayer to match the in-plane lattice constant of the substrate, as shown in Fig. 1.15a. The deposited material is coherently strained to the substrate which means that the lattices planes are continuous at the interface. In this case, the epitaxial film is compressively strained if its lattice constant is larger than that of the substrate, otherwise it is tensily strained. On the other hand, the plastic relaxation occurs for larger mismatches or thicker epitaxial films 20 (Fig. 1.15b). In this case, if the epitaxial film thickness is above a given "critical thickness", elastic energy will be released through the creation of the so-called misfit dislocations. The critical thickness was investigated in many materials systems by scientists. In the case of GaP/Si, Soga et al. [START_REF] Soga | Dislocation generation mechanisms for GaP on Si grown by metalorganic chemical vapor deposition[END_REF] reported a critical thickness less than 90 nm, and Takagi et al. [START_REF] Takagi | Lattice relaxation process and crystallographic tilt in GaP layers grown on misoriented Si(001) substrates by metalorganic vapor phase epitaxy[END_REF] and Skibitzki et al. [START_REF] Skibitzki | GaP collector development for SiGe heterojunction bipolar transistor performance increase: A heterostructure growth study[END_REF] gave further estimation around 45-95 nm and 64 nm, respectively. In the case of GaSb/Si, the mismatch is so high that dislocations appear almost instantaneously (the critical thickness is very small). In this case, the formation of a regular network of misfit dislocations have been observed with Sb-based materials [3,[START_REF] Huang | Simultaneous interfacial misfit array formation and antiphase domain suppression on miscut silicon substrate[END_REF].

In addition, the lattice parameter also depends on the temperature. Therefore, the thermal expansion coefficient has a main impact on lattice mismatch issues. This effect can be very different from one semiconductor to another [START_REF] Pohl | Epitaxy of Semiconductors: Introduction to Physical Principles[END_REF]. When the temperature varies from room temperature to growth temperature (in the range of 350°C to 1000°C), microcracks can be induced, due to the different thermal expansion coefficients of the substrate and the epilayer. In the case of III-V/Si epitaxy, the III-V epilayer may undergo a large tensile or compressive stress. For instance, the thermal expansion coefficient of GaP is almost twice the one of the silicon (αGaP=4.65*10 -6 /K and αSi=2.4*10 -6 /K), which means that the GaP layer will be more compressively strained at growth temperature. In fact, the lattice mismatch is 0.52% at typical GaP/Si MBE growth temperature (580°C). However, it will become 0.36% at room temperature.

Polar on non-polar epitaxy

While the management of misfit dislocations and strain is a general concept in epitaxy, other defects are more specifically related to III-V/Si heterogeneous epitaxy. This is the case of stacking faults (SFs), microtwins (MTs) and antiphase domains (APDs).

These defects are attributed to the difference of the two crystal symmetries and surface energies, and are specifically observed when a polar semiconductor (e.g. a III-V compound) is heteroepitaxially deposited on a non-polar one (e.g. the silicon).

Stacking faults and microtwins

The stacking faults are planar defects which alter the crystallographic plane order. In addition, a stacking fault is bounded by two partial dislocations given by the dissociation of a perfect dislocation (which has a higher energy). If the crystallographic order is changed so that the atoms arrange in a mirror symmetric-like order, the stacking fault formed is called microtwin. The microtwins can be described as a 60° or 180° crystal rotation with respect to the <1 1 1> (Fig. 1.17). This is commonly observed in the case of Zinc-Blende III-V/Si epitaxy. It should be noted that to our knowledge, the precise origin and description of their formation (chemical preparation of the silicon, influence of silicon steps, growth temperature, surface energy) was not yet clarified in the literature up to now.

Antiphase domains and antiphase boundaries

As mentioned previously, the III-V Zinc-Blende material structure is a material with polar bonds where each group-III atom is bounded to four group-V atoms. The Si diamond crystal structure is however non-polar, and each Si atom is bounded to other four Si atoms. During the epitaxial growth of III-V semiconductors on Si, antiphase domains can be formed, which are bounded by the so-called antiphase boundaries (APBs). APBs are highlighted by the reversal of the III-V bond polarity. A sketch of a vertical stoichiometric APB at the atomic scale is shown in Fig. 1. [START_REF] Vurgaftman | Band parameters for III-V compound semiconductors and their alloys[END_REF]. More details about APBs will be given in the following.

Generation mechanism and properties of antiphase boundaries

For years, APBs were considered as the most important issue to be understood during III-V/Si epitaxy, as their presence strongly modify electrical and optical properties of the III-V/Si devices, as well as the surface roughness of samples. Its generation mechanism and properties were thus widely investigated since the 80's for the different applications of III-V/Si materials.

Classification and formation energy of antiphase boundaries

An AntiPhase Domain (APD) is a domain composed of bulk semiconductor, bounded by APBs. APBs are planar defects consisting of anti-site ("wrong") bonds, i.e.

III-III and/or V-V bonds. The existence of APBs was first proposed by Aminoff & Broome in 1931 [START_REF] Aminoff | Strukturtheoretische Studien über Zwillinge[END_REF]. APBs can be classified as stoichiometric ones or non-stoichiometric ones, depending on the total density of each type of wrong bonds. The stoichiometric ones have equal numbers of III-III and V-V bonds within the same APB as shown in Fig. 1.19a. On the other hand, non-stoichiometric ones do not have the same number of III-III and V-V

bonds. An extreme case can be found when considering APBs composed only of III-III bonds or V-V bonds, which are supposed to be highly charged, as shown in Fig. 1.19b. In the different chapters of this thesis, we will sometimes consider only this extreme case as representative of non-stoichiometric APBs for the sake of simplicity. The formation energies of stoichiometric and nonstoichiometric APBs lying along different crystallographic plane (including {110}, {111}, {112}, and {113}) of GaAs and GaP were studied theoretically using a full-potential linearized augmented plane-wave densityfunctional approach and described in the pioneering works of Rubel et al. [START_REF] Rubel | Formation Energies of Antiphase Boundaries in GaAs and GaP: An ab Initio Study[END_REF]. The corresponding results are shown in Fig. 1.20. They found that the formation energy of APB's in GaP is about 20% higher than that in GaAs. Besides, it reveals that the nonstoichiometric {111} APB has the highest formation energy, while the stoichiometric {110} APB has the lowest energy due to charge compensation effect. Both {112} and {113} APB's have an intermediate value of the formation energy. On this basis, many papers in the literature considered that the generation of stoichiometric APBs should be the dominant process during III-V/Si heteroepitaxy. We will show in next chapters that it is not the case.

Formation mechanism of antiphase boundaries 1.4.2.1 Historical approach

At the beginning of my thesis, the formation mechanism of APBs, and the role of the starting Si surface, was mostly understood through the historical models proposed in the 80's. In 1987, two different mechanisms for the generation of APBs were firstly proposed by Kroemer [START_REF] Kroemer | Polar-on-nonpolar epitaxy[END_REF]. In this seminal work, it was considered that if neither group-III nor group-V elements have a preferential probability of bonding to the substrate surface, APBs could be induced at the interface between the III-V semiconductors and the Si substrate, as shown in Fig. 1.21a. In order to simplify the process, it has been widely accepted that the deposition of III-V materials on Si (001) always starts with the same kind of atom in the literature, as shown in Fig. 1.21b-d.

In addition, Kroemer also proposed a strong correlation between the APBs and the kind of steps present on the Si (001) surface. If the III-V layer deposition nucleates always with the same atom on the Si, a Si surface populated only with D-steps does not lead to the formation of APBs, as shown in Fig. 1.21b. However, when the step height corresponds to an odd number of atomic planes, the III-V material grows with inverse polarities on each side of the step generating an APB, as shown in Fig. 1.21c and Fig. 1.21d [START_REF] Kroemer | Polar-on-nonpolar epitaxy[END_REF][START_REF] Calvo | GaSb sur substrat Si(100) "on-axis" pour l'optoélectronique intégrée[END_REF].

The idea above rapidly became the main motivation for using substrates with a miscut, in order to form biatomic steps (Fig. 1.21b) at the Si surface and avoid theoretically the generation of APBs in the subsequent III-V growth process. Based on this model, K. Volz et al. explained their remarkable achievements about high quality epitaxial GaP/Si by considering a 2D III-V growth mode on the substrate [START_REF] Kroemer | Polar-on-nonpolar epitaxy[END_REF]. The recent studies in our group and in this thesis strongly challenge this view, which will be discussed in the following parts. Besides, based on the index planes and types of APBs, researchers discussed the propagation or annihilation of APBs [START_REF] Kawabe | Self-Annihilation of Antiphase Boundary in GaAs on Si(100) Grown by Molecular Beam Epitaxy[END_REF]. It was considered that the stoichiometric APBs with lower formation energies usually form perpendicularly to the interface. In this case, the stoichiometric APB will propagate easily and cross the whole III-V layer, emerging at the surface. For the non-stoichiometric APBs with higher formation energy, they were considered to form in higher index planes with an inclined angle with respect to the substrate. In this case, it may lead to cross another APB, and thus allowing the so-called annihilation of these defects. In the present work, we will also challenge this view. For years, researchers have been widely adopting the above model for the description of APBs generation, and as an explanation for the needs of using a miscut to achieve good quality III-V/Si epilayers [START_REF] Calvo | GaSb sur substrat Si(100) "on-axis" pour l'optoélectronique intégrée[END_REF]. Indeed, this model provides an easy representation about APBs generation and propagation, but also faces fundamental inconsistencies with the experimental data published during these last 40 years. Together with French academic partners (IES, IPR, C2N, CEMES and CRHEA), our laboratory revisited this model in 2018 [3] and in 2020 [4], which will be explained in the next section. which shows that AlSb nucleation on Si generates 3D islands. This growth mode is also experimentally confirmed by some other studies [START_REF] Ponchet | Interface energy analysis of III-V islands on Si (001) in the Volmer-Weber growth mode[END_REF][START_REF] Ernst | Formation of planar defects in the epitaxial growth of GaP on Si substrate by metal organic chemical-vapor deposition[END_REF]. Therefore, the 2D nucleation growth mode of III-V semiconductor layer on Si substrate in the historical approach shown in Fig. 1.21 simply cannot happen. Secondly, they demonstrated that APBs are generated during the coalescence of 3D islands having different phases, which was briefly suggested in pioneering works on III-V/Ge [START_REF] Petroff | Nucleation and growth of GaAs on Ge and the structure of antiphase boundaries[END_REF]. They also proved that the size of individual monodomain islands can be much larger than the distance between steps (the terrace width), suggesting that the step itself cannot be the main cause for the APDs generation, which also contradicts the historical Kroemer's model. Thirdly, they showed that the elastic energy does not have a significant impact on the island morphology at the coalescence growth step for most III-V/Si materials systems (and therefore does not impact the APD distribution), as most of the islands are already plastically relaxed (even for GaAs/Si [START_REF] Hull | Nucleation of GaAs on Si: Experimental evidence for a three-dimensional critical transition[END_REF]) at this stage (or quasi-unstrained in the particular case of the quasilattice-matched GaP/Si system). In addition, they made the hypothesis that the epitaxial relationship is defined locally at the nucleation site which further governs the phase distribution.

In 2020, this work was further developed, generalized and completed by our group [4]. We investigated comprehensively the effect of the Si substrate miscut on III-V epitaxy on Si and proposed two regimes that determine the initial APB distribution. At very low miscut angles, several islands of the same domain can nucleate on a terrace. The APB distribution is then expected to follow the step distribution on the surface. This is the socalled terrace-driven regime (Fig. 1.23a). At higher miscut angles, monodomain islands nucleate randomly, and as explained above, the size of these islands is larger than the terraces width. Therefore, the APB distribution is no longer correlated to the step distribution. This is the so-called nucleation-driven regime (Figure 1.23b). We also demonstrated that for the majority of III-V semiconductors grown on Group-IV substrates, the critical miscut angle at which the regime evolves from terrace-driven to nucleationdriven is in the range of [0.1-1]°. In this thesis, most of the Si substrates used for the deposition of III-V layers have an intentional or residual miscut larger than 0.1° [START_REF] Calvo | GaSb sur substrat Si(100) "on-axis" pour l'optoélectronique intégrée[END_REF].

Therefore, in most cases the formation of the two different polar-domains will be considered as governed by the nucleation-driven regime. Besides, through introducing the growth rate imbalance coefficient Cα/β (which is a simple ratio between the growth rate of two adjacent phases α and β), we discussed the growth conditions requirements needed for the propagation or burying of APBs. When the Cα/β is lower or larger than unity, the β phase or α phase will grow faster than the other one. Like this, only one phase will become dominant, and therefore the APDs will be buried. When the Cα/β equals to one, the growth rates are the same and APBs will propagate to the surface. Moreover, we demonstrated how the growth rate imbalance is closely related to the growth temperature and V/III ratio chosen during the III-V/Si growth. As a general conclusion of these recent experimental and theoretical developments, one can keep in mind that these works give the keys to ultimately control the lateral and vertical extents of APBs and APDs [START_REF] Saleem-Urothodi | Loss assessment in random crystal polarity gallium phosphide microdisks grown on silicon[END_REF][START_REF] Guillemé | Antiphase domain tailoring for combination of modal and 4¯ -quasiphase matching in gallium phosphide microdisks[END_REF].

29

Optoelectronic properties of antiphase boundaries

Many works in the broad literature on the subject of III-V/Si monolithic integration made reference to APBs as detrimental defects for photonic or energy harvesting devices [START_REF] Galiana | Characterization of antiphase domains on GaAs grown on Ge substrates by conductive atomic force microscopy for photovoltaic applications[END_REF][START_REF] Paladugu | Site Selective Integration of III-V Materials on Si for Nanoscale Logic and Photonic Devices[END_REF], without providing explanation of this observation. Of course, this statement is based on the common observation of scientists that any laser or photovoltaic device simply cannot work if APBs are not buried in the structure. To explain this behavior, scientists have claimed for decades that APBs are non-radiative recombination centers.

But, as we are going to explain in this thesis, this is far from the real situation. Indeed, when looking carefully at the atomic scale structure of one antiphase boundary, the topology observed is fundamentally different as compared with e.g. grain boundaries in other polycrystalline materials, as they do not contain any dangling bonds, and respect the lattice ordering. Surprisingly, at the beginning of this thesis, only a few works focused specifically on the study of antiphase boundaries optoelectronic properties.

Conclusions and objectives of the thesis

In summary, the integration of III-V semiconductors on Si combining the advantages of Si (earth abundance, low cost and prevalence in the electronics and PV industries) with the excellent optoelectronic properties of III-V, is of great interest in photonics and energy harvesting applications. The different III-V/Si integration approaches and the specific issues related to the monolithic integration were presented in this chapter. Among the different crystal defects generated at the III-V/Si interface and propagating in the materials and devices, the recent researches about the antiphase domain and boundaries generation, propagation and burying were reviewed, giving keys to tune at will their distribution in III-V/Si samples.

At the beginning of my thesis, the main objective was to clarify and assess optoelectronic properties of III-V (GaP-based)/Si samples and devices. Rapidly, from the results obtained, we jointly decided to refocus the thesis objectives on the theoretical and experimental study (Chapter 2) of the specific optoelectronic properties of APBs, and its use for energy harvesting devices. This led to the following main results:

-First demonstration of the strong electron-phonon coupling arising in stoichiometric III-V/Si APBs and clarification of optoelectronic properties (Chapter 3).

-First demonstration of a bulk III-V/Si photoelectrode for solar water splitting applications, with assessment of III-V/Si bandgaps and bandlineups, and promising PEC performances (Chapter 4).

-First evidence of the hybrid 2D-semimetal/semiconductor ambipolar and photoactive properties of epitaxial III-V/Si samples with non-stoichiometric APBs (Chapter 5).

Chapter 2: Experiments and DFT calculations

In this chapter, the experimental and theoretical tools supporting this work are presented. All the samples used in this work were grown by Molecular Beam Epitaxy (MBE), except when bare substrates were used as a reference. The basic principle of MBE sample growth in our lab is shown in appendices and the specific growth parameters will be given when needed in the following chapters. Here, the different experimental characterizations setups are first described, with the related physical mechanisms, and parameters settings used in this thesis. The DFT calculations are then introduced, with a brief introduction to the basic concepts of the DFT, followed by the calculation steps and parameter settings for the determination of electronic band structure, and of the vibrational and electron-phonon coupling properties.

Structural investigations and surface processes

Atomic force microscopy

Atomic Force Microscopy (AFM) is a characterization technique used to obtain the topography of the surface of a sample. This technique is easy to use because no special treatment of the sample is needed before measurement, and it is a technique that can be used for a large variety of materials, including soft or non-metallic materials. The AFM used in this thesis is a Veeco di Innova setup, presented in Fig. 2.1, allowing producing 100×100 μm 2 images with a lateral resolution of 15 nm and a vertical resolution of 0.2 nm in standard conditions. The operating principle of the AFM is to use a fixed tip (typically made by Si3N4 or Si) on a deformable cantilever to probe the surface of the sample through the tip/surface interaction. Depending on the type of sample analyzed, the interaction can occur through different types of forces: Van Der Waals force, capillarity force, electrostatic force or magnetic force. The deviation of the cantilever Z is proportional to the interaction force F: F = -kZ. The radius of the tip used is typically around few nanometers, allowing an accurate surface characterization of the sample at a nanometer scale. A red laser source is focalized on the top of the cantilever and reflected to a four-quadrant photodiode to record the spatial variation of laser intensity due to the cantilever's deviation (X, Y and Z directions), and the cantilever's torque. Meanwhile, the substrate is moved to keep a constant interaction between the tip and substrate surface. The movement is realized by a cylinder-shaped piezoelectric (PZT) tube, whose shrinkage and expansion will lead to a movement in three dimensions. Through a loop system, the Z-movement of the piezoelectric tube is indirectly controlled by the signal received from the photodiode. AFM measurements can be performed in two basic working modes (see Fig. 2.2):

contact mode and non-contact mode ("tapping" mode). In the contact mode, the type of force is repulsive and the tip scans over the surface maintaining the cantilever vertical deviation at a constant level. Before starting the measurement, the laser is adjusted at the end of the cantilever and the reflected spot is positioned as a reference at the center of the photodiode. During the scanning, the deflection is recorded by the optical system and is compared by a DC feedback loop to the set-point level. The resulting error signal from the amplifier is used to actuate the PZT, moving the sample in order to maintain constant the tip/sample interaction by applying the required voltage. The tapping mode alternatively uses both the attractive and repulsive forces. The cantilever vibrates at high frequency. Therefore, the detected signal on the photodiode alternates. The amplitude of the vibration is sensitive to the force between tip and sample surface, which changes with the distance between them. In this situation, the feedback loop aims to maintain a constant oscillation amplitude keeping constant the average distance between the cantilever and the surface. As compared to the contact mode, tapping mode shows some advantages:

samples with soft surfaces will be less damaged and the tip will be less contaminated. In this thesis, all the AFM measurements were performed in tapping mode with the cantilever tuned around 293 kHz.

Conductive atomic force microscopy

In order to characterize the transport properties of APBs of III-V/Si samples, Conductive Atomic Force Microscopy (C-AFM) measurements were performed, in close collaboration with IPR (Institut de Physique de Rennes).

C-AFM is a measurement technique that derives from atomic force microscopy, where a conductive tip scans the sample surface and a voltage is applied between the tip and the sample to generate an image related to the current measured. A topographic image is also generated simultaneously. Both the current and the topographic images are taken from the same area of the sample, which allows the identification of features on the surface conducting more or less the current, as shown in Fig. 2.3, and allowing the identification of possible current artefacts that could simply be due to the morphology.

The main advantage of C-AFM over standard electrical measurement techniques is the high spatial resolution. On the other hand, a meaningful quantitative evaluation of the local current is hard to achieve, and most of the time, qualitative analysis is only proposed.

In this work, we will use only qualitative comparisons between the different parts of the samples. In this thesis, the C-AFM measurements are carried out on a Bruker Multimode AFM with a Pt/Ir coated Sb doped silicon cantilever (spring constant of 0.1 N/m, and resonant frequency of 10 kHz). Before measurements, the sample surface was exposed to a HF-5% solution for 1 minute (to remove the surface oxide), then the samples were rinsed in deionized water and then quickly transferred to the C-AFM set-up. Measurements were performed under dry nitrogen gas to limit humidity. Then different biases were applied to the III-V/Si samples at 0.3 Hz in contact mode. All current mapping images were obtained with the cantilever moving across the thin film surface. In this thesis, the compositions of different III-V/Si samples are determined with the XRD technique. For example, Figure 2.5 shows the XRD patterns for one GaP1-xSbx/Si sample by using a ω/2θ scan. Assuming that the GaP1-xSbx epitaxial layer is fully plastically relaxed on the Si substrate, the free-of-strain cubic lattice constant of GaP1-xSbx (𝑎 𝑓𝑜𝑠 ) can be obtained based on the equation:

X-ray diffraction

𝑎 𝑓𝑜𝑠 = 𝑛 2 sin 𝜃 √ℎ 2 + 𝑘 2 + 𝑙 2 (2-1)
Where n is equal to 1 for the first order diffraction maximum,  is the wavelength of the X-ray source (Cu target was used in this thesis),  can be obtained in the ω/2θ scan, and the (h, k, l) are the Miller indices.

Then the Sb content in GaP1-xSbx can be calculated from the Vegard's law:

x=

𝑎 𝑓𝑜𝑠 -𝑎 𝐺𝑎𝑃 𝑎 𝐺𝑎𝑆𝑏 -𝑎 𝐺𝑎𝑃 (2-2)
Where 𝑎 𝐺𝑎𝑃 and 𝑎 𝐺𝑎𝑆𝑏 are the lattice constant of GaP and GaSb. Based on the equations above, the Sb content for the studied sample is 33% here.

However, if the relaxation and strain state of the epitaxial layer is unknown or not certain, then reciprocal space mappings (RSM) will be needed to determine the relaxation degree (R) and calculate the composition. Lattice constants in the in-plane and out-of-plane directions can be directly deduced from the RSM [START_REF] Nguyen | Sillicon photonics based on monolithic integration of III-V nanostructures on silicon[END_REF]:

𝑎 𝐿 // = 𝑎 𝑆 𝑆 𝑠𝑢𝑏 // 𝑆 𝐿 //
(2-3)

𝑎 𝐿 ⊥ = 𝑎 𝑆 𝑆 𝑠𝑢𝑏 ⊥ 𝑆 𝐿 ⊥ (2-4)
Where 𝑎 𝐿 // and 𝑎 𝐿 ⊥ are the lattice constant of the epitaxial layer in the directions parallel and perpendicular to the surface (in-plane and out-of-plane), 𝑎 𝑆 is the lattice constant of the substrate, 𝑆 𝐿 // , 𝑆 𝐿 ⊥ , 𝑆 𝑠𝑢𝑏 // and 𝑆 𝑠𝑢𝑏 ⊥ a re the positions of the layer and the substrate diffraction spots in the RSM in the directions parallel (corresponding to Sx) and perpendicular(corresponding to Sz) to the surface. The free of strain cubic lattice constant of the GaPSb layer (𝑎 𝐿 𝑓𝑜𝑠 ) is then given by:

𝑎 𝐿 𝑓𝑜𝑠 = 𝑎 𝐿 ⊥ -𝛼𝑎 𝐿 // 1-𝛼 = 𝑎 𝐿 ⊥ +2 𝐶 12 𝐶 11 𝑎 𝐿 // 1+2 𝐶 12 𝐶 11 (2-5)
Where 𝛼 = The relaxation degree (R) can also be determined by using:

𝑅 = 𝑎 𝐿 // -𝑎 𝑆 𝑎 𝐿 𝑓𝑜𝑠 -𝑎 𝑆 (2-7)
In Fig. 2.6, the black dotted line represents the full relaxation line, which goes through the center of Si RSM signal and the origin point (Sz=Sx=0). In this case, the signal of GaPSb appears around the full relaxation line, indicating that the GaPSb alloy is almost fully plastically relaxed. In this thesis, the full relaxation lines will be given in all the RSM for the samples studied. If the peak of the epitaxial layer is vertically aligned on the substrate peak, then the layer will be considered as fully biaxially strained (R=0). If the peak of the layer is located around the full relaxation line, then the layer will be considered as fully relaxed (R=1). In between, the samples will be considered as partially relaxed.

40

The XRD characterizations in this thesis are performed on a 4 circles Brucker D8

Diffractometer. A Bartels asymmetric Ge (220) monochromator was used for both line scans and reciprocal space maps. The detection is ensured by a Lynxeye TM , 1 dimensional position sensitive detector (PSD) allowing a collection angle of 2.6° over 2θ.

Scanning transmission electron microscopy

The scanning transmission electron microscope (STEM) has become one of the prominent structural analysis tools for high spatial resolution imaging and spectroscopy of materials, most notably at atomic resolution [START_REF] Nellist | Scanning Transmission Electron Microscopy[END_REF]. A highly-focused electron probe is raster-scanned across the material, and the various types of electron scattered are then collected and analyzed as a function of the position. The sample needs to be thinned to guarantee transmission of electrons, and limit the volume of the probed layer. The transmitted electrons at high scattering angle can be collected to form high-resolution, chemically sensitive, atomic number contrast images. The X-rays generated can be collected using an energy-dispersive X-ray spectroscopy (EDX) detector and further used to form high spatial resolution compositional maps.

The STEM measurements in this thesis were performed by our German partner (IHP-microelectronics). The STEM high-resolution imaging and Energy Dispersive X-Spectroscopy (EDX) (presented in Chapter 3) were obtained by a FEI Tecnai Osiris TEM tool, operating at 200 kV with implemented energy-dispersive x-ray (EDX) detector.

TEM lamellas were prepared by milling and undercutting processes using a Zeiss NVision 40 focused ion beam operating at up to 30 kV. In this thesis, to improve the surface quality, III-V/Si samples were polished applying CMP on the surface with 1% H3PO4.

Chemical mechanical polishing

Raman spectroscopy

Raman spectroscopy is a spectroscopic technique typically used to determine vibrational modes of molecules. Raman spectroscopy is based on inelastic scattering of photons, known as Raman scattering. Typically, for Raman measurement, the samples are illuminated by a monochromatic laser source and the scattered light is then collected in a spectrometer.

When the samples are illuminated by a monochromatic laser source, most of the scattering that takes place is elastic, with no loss of energy, and therefore no frequency change. This is known as Rayleigh scattering (Fig. In this thesis, the Raman measurements were performed in collaboration with both IPR (Institut de Physique de Rennes) and IHP microelectronics (Germany). At IPR, the Raman image was obtained with a HR800 micro-Raman spectrometer (HORIBA Scientific). The HR800 system is a high-resolution spectrometer fitted with a confocal microscopy system coupled to different lasers. Our experiments were performed in "confocal mode" with a 532 nm laser diode for excitation and, to avoid sample heating, an optical density filter was used.

Ellipsometry

Ellipsometry is traditionally used to determine the complex refractive index of thin film materials (n '= n + ik) and their thickness. It is a non-destructive technique, currently used for determining the dielectric functions of thin films and substrates [START_REF] Aspnes | Spectroscopic ellipsometry-A perspective[END_REF]. In this thesis, we mainly focused on the determination of the complex part (k) of the refractive index to deduce the absorption coefficient (α) of the III-V layers grown on Si substrate by applying the well-known formula:

𝛼 = 4𝑘 𝜆 (2-8)
Ellipsometry principle is based on the measurement of a change in polarization as light reflects or transmits from/through a material (see schematic image in Fig. 2.12). The polarization change is represented as an amplitude ratio, Ψ, and the phase difference, Δ, according to the following relation:

𝑟 𝑝 𝑟 𝑠 = tan Ψ. 𝑒𝑥𝑝{𝑖∆} (2-9)
Where rp and rs are respectively the complex reflection coefficient for p and s polarization. Our experimental setup (Horiba UVSEL 2) comprises two detectors for a spectral range from UV light (6 eV) until infrared wavelengths (0.5 eV). This equipment mainly consists in two optical arms and a sample-holder (as shown in Fig. 2.12). The first arm comprises a Polarization State Generator (PSG) coupled to a light source. In every case, the PSG includes a linear polarizer set at an azimuth P with respect to the plane of 46 incidence. The second arms is used to determine the polarization and intensity of the outcoming beam. It is composed of a Polarization State Analyser (PSA) and a detector.

Figure2.12 Real and schematic [START_REF] Garcia-Caurel | Application of Spectroscopic Ellipsometry and Mueller Ellipsometry to Optical Characterization[END_REF] images of the ellipsometry setup.

From the two values Ψ and , the ratio () between rs and rp can be deduced. For a bulk material (a substrate), by taking into account the angle of incidence (), the so-called "pseudo" dielectric function can be easily obtained with the equation below:

〈ε〉 = 𝜀 𝑟 + i𝜀 𝑖 = 𝑠𝑖𝑛 2 θ[1 + 𝑡𝑎𝑛 2 𝜃( 1-𝜌 1+𝜌 ) 2 ] (2-10)
However, for thin film deposited on a substrate, no direct formula can be applied, but the use of a model allows the measurement to be fitted. Thus, extraction of the absorption coefficient of a given single layer, or even of a layer in a multiple-stack structure can be done.

Photoelectrochemical characterizations

Photoelectrochemical (PEC) water splitting allows turning the liquid water into hydrogen and oxygen gases, based on semiconducting photoelectrodes under sunlight and a given voltage [START_REF] Kang | Electrochemical Synthesis of Photoelectrodes and Catalysts for Use in Solar Water Splitting[END_REF].

Linear sweep voltammetry

Linear sweep voltammetry (LSV) is a voltammetric method where the current at a working electrode is measured while the potential between the working electrode and a reference electrode is swept linearly in time [START_REF] Nahir | Linear-Sweep Voltammetry of Irreversible Electron Transfer in Surface-Confined Species Using the Marcus Theory[END_REF]. Normally, the LSV can be measured under three different conditions: in the dark, under light and in chopped light mode (dark and light states shown alternatively).

Incident photon-to-current efficiency spectra

The incident photon-to-current efficiency (IPCE) is a measure of the ratio of the photocurrent (converted to an electron transfer rate) versus the rate of incident photons (converted from the calibrated power of a light source) as a function of wavelength [START_REF] Chen | Incident Photon-to-Current Efficiency and Photocurrent Spectroscopy[END_REF].

The IPCE measurements were performed with a CIMPS-QE IPCE 3 workstation (Zahner) comprising a TLS03 tunable light source controlled by a PP211 potentiostat in the same cell as the one used for classical electrochemical experiments. The applied potentials were set to -0.74 V and 1.26 V vs RHE for the IPCE results presented in Chapter 5. The light modulation frequency was 1 Hz, the settling time was 5 s and the number of counts was 25. The Thales software provides the spectra in photocurrent efficiency (A W -1 ) or in IPCE (%). In order to check the validity of the IPCE measurements, the photocurrent integration is performed from the photocurrent density spectrum (in A W -1 ) and using the AM 1.5G reference solar spectrum obtained from the ASTM (American Society for Testing Materials) [START_REF] Oh | Elucidating the performance and unexpected stability of partially coated watersplitting silicon photoanodes[END_REF]. The result is compared to the one determined by LSV.

Electrochemical impedance spectroscopy

Electrochemical impedance spectroscopy (EIS) uses the application of a small, sinusoidal, alternating current (AC) or voltage to the sample under investigation. The AC perturbation (either voltage or current) is applied over a range of frequencies, and the electrical response of the system is measured and analyzed over this range of frequencies.

In this case, the impedance can be represented as a complex function [START_REF] Barsoukov | Impedance Spectroscopy: Theory, Experiment, and Applications[END_REF]:

𝑍 = 𝑉 𝑡 𝐼 𝑡 = 𝑉 0 𝑒 𝑗𝜔𝑡 𝐼 0 𝑒 𝑗𝜔𝑡-𝑗𝜑 = 𝑍 0 𝑒 𝑗𝜑 (2-13)
Where Z is the impedance, V is the voltage, I is the current. V and I have a phase difference of 𝜑. j is the imaginary component and 𝜔 is the angular frequency.

Then the impedance can be expressed as:

50 𝑍 = 𝑍 0 (cos(𝜑) + 𝑗𝑠𝑖𝑛(𝜑)) (2-14)
real and imaginary parts of the impedance can thus be written:

𝑍 𝑟𝑒𝑎𝑙 = 𝑍 ′ = 𝑍 0 cos(𝜑) (2-15) 𝑍 𝑖𝑚𝑔 = 𝑍 ′′ = 𝑍 0 sin (𝜑) (2-16)
With a phase angle

𝜑 = tan -1 𝑍 ′′ 𝑍 ′ (2-17)
And a modulus

|𝑍| = [(𝑍′) 2 + (𝑍") 2 ] 1/2 (2-18)
Based on the formulas above, different parameters can be obtained by EIS measurements. Therefore, the data can be presented in different ways. For instance, real and imaginary parts of the impedance can be plotted against one another, which is called Nyquist plots (Fig. semiconductor photoelectrode.

The Mott-Schottky curve is often used to determine the flat band potential or doping density of the system. The Mott-Schottky equation makes the link between the capacitance C and the applied potential E across a semiconductor-electrolyte junction following:

1 𝐶 2 = 2 𝑒𝑁 𝐷𝑜𝑝𝑎𝑛𝑡 𝐴 2 Ꜫ 0 Ꜫ 𝑟 (𝐸 -𝐸 𝑓𝑏 - 𝑘𝑇 𝑒 ) (2-19)
Where, 𝜀 𝑟 is the relative permittivity of the semiconductor, 𝜀 0 is the permittivity in vacuum, A is the surface area, e is the charge of an electron, 𝑁 𝐷𝑜𝑝𝑎𝑛𝑡 is the free carrier density, k is Boltzmann constant and T is the temperature. 𝐸 𝑓𝑏 corresponds to the flatband potential which can be extracted from the x-intercept of the linear portion of the MS curve. For the p-doped and n-doped semiconductors, the MS curve will show negative or positive slop respectively, as shown in Fig. 2.16. 52

DFT calculations

Introduction to density functional theory

Density functional theory (DFT) is one of the most widely used ab initio methods for the determination of electronic properties of atoms, molecules, crystals, surfaces, and their interactions [START_REF] Argaman | Density functional theory: An introduction[END_REF]. It is considered as the most successful quantum mechanical approach for describing the electronic structure of matter. In the quantum mechanics theory, the system's wave function contains all the information of a given system, which can be obtained by solving the Schrödinger equation:

𝐻 ̂𝛹 = 𝐸𝛹 (2-20)
Where 𝐻 ̂ is the Hamiltonian operator, 𝛹 is the wave function and 𝐸 is the energy. The exact solution of this equation for a many-body system (which contains 3N variables; N being the number of electrons) is analytically impossible. Therefore, a series of approximations and numerical methods are introduced to solve this problem.

D. Hartree introduced one of the first procedures to calculate approximately wave functions and energies [START_REF] Hartree | The wave mechanics of an atom with a non-Coulomb central field. Part I. Theory and methods[END_REF][START_REF] Hartree | The wave mechanics of an atom with a non-coulomb central field. Part II. Some results and discussion[END_REF] to various atoms and ions (He, R + , Na + and Cl -), in which the electrons are considered independent and interact only by the average static field with the other electrons. Therefore, the 3N dimensional many body Schrödinger equation is transformed to a set of single particle equations. However, due to the mandatory antisymmetric property of the wavefunction, this method is incompatible with the Pauli Exclusion Principle [START_REF] Slater | Note on Hartree's Method[END_REF][START_REF] Fock | Näherungsmethode zur Lösung des quantenmechanischen Mehrkörperproblems[END_REF]. In order to overcome this problem, the Slater determinant [START_REF] Slater | Note on Hartree's Method[END_REF][START_REF] Slater | The Theory of Complex Spectra[END_REF] was considered by Hartree and Fock in 1930 [START_REF] Fock | Näherungsmethode zur Lösung des quantenmechanischen Mehrkörperproblems[END_REF].

In parallel of the Hartree method, Thomas [START_REF] Thomas | The calculation of atomic fields[END_REF] and Fermi [START_REF] Fermi | Statistical method to determine some properties of atoms[END_REF] proposed another approach, which was further improved by Dirac [START_REF] Dirac | Note on Exchange Phenomena in the Thomas Atom[END_REF]. In this pioneer approach, a manybody system is not described by its wave function but by its electronic density. This is the cornerstone of the development of density functional theory, simplifying the N-body system problem by decreasing the degrees of freedom from 3N to only 3 variables. The 53 question was then to know if the electronic density was sufficient or not to describe the electronic structure of a system. In 1964, Hohenberg and Kohn tackled this problem and solved it by proposing 2 new theorems [START_REF] Hohenberg | Inhomogeneous Electron Gas[END_REF]. One year later, a practical method to calculate the electronic density was proposed by Kohn and Sham [START_REF] Kohn | Self-Consistent Equations Including Exchange and Correlation Effects[END_REF]. At this stage, DFT simulations became thus consistent and feasible. Despite all the approximations proposed to simplify the calculation method, Hohenberg-Kohn and Kohn-Sham theorems are still considered as the major breakthroughs [START_REF] Lejaeghere | Reproducibility in density functional theory calculations of solids[END_REF], which led to the formulation of the density functional theory used today.

Hohenberg-Kohn theorems

In 1964, Hohenberg and Kohn demonstrated that the ground state of a many particles system is a unique functional of the electron density 𝜌(𝒓). They proposed to replace all the various interactions in the system by an external potential and formulate the ground state energy of the system as a functional of the electronic density for further simplification. Therefore, the total energy of a system can be expressed as:

𝐸[𝜌] = 𝑇[𝜌] + 𝑉 𝐸𝐸 [𝜌] + 𝑉 𝑁-𝑒 [𝜌] = 𝐹 𝐻𝐾 [𝜌] + ∫ 𝑑𝑟 3 𝜌(𝒓) 𝑉 𝑁-𝑒 (𝒓) (2-21)
Where, 𝐹 𝐻𝐾 [𝜌] is a universal functional related to the functional of the kinetic energy coming from the interacting electrons 𝑇[𝜌] and the functional of the electron-electron interactions 𝑉 𝐸𝐸 [𝜌]. The potential 𝑉 𝑁-𝑒 [𝜌] corresponds to the total external potential due to the nuclei which are considered as clamped.

Kohn-Sham approach

The energy functional above contains a kinetic energy term. There is no known expression for this term, and thus the functional cannot be evaluated as it stands in the Hohenberg-Kohn formulation.

In order to overcome this, Kohn and Sham replaced the system with a fictitious noninteracting system that generates the same ground state electron density than the real system of interacting electrons. The 𝐹 𝐻𝐾 [𝜌] was redefined by introducing the following functional 𝐸 𝑋𝐶 [𝜌]:

𝐹 𝐾𝑆 [𝜌] = 𝑇 𝑆 [𝜌] + 𝑉 𝐻 [𝜌] + 𝐸 𝑋𝐶 [𝜌] (2-22)
Where 𝑇 𝑆 [𝜌] is the kinetic energy functional of non-interacting electrons, 𝑉 𝐻 [𝜌] is the Hartree potential functional and 𝐸 𝑋𝐶 [𝜌] is the exchange and correlation energy. The associated potential to the latter is:

𝑉 𝑥𝑐 (𝒓) = 𝛿𝐸 𝑥𝑐 [𝜌] 𝛿𝜌 | 𝜌(𝑟) (2-23)
So, the equation changes to:

[ 1 2 ∇ 2 + 𝑉 𝑁-𝑒 (𝒓) + 𝑉 𝐻 (𝒓) + 𝑉 𝑥𝑐 (𝒓)] 𝜑 𝑖 (𝑟) = 𝜖 𝑖 𝜑 𝑖 (𝒓) (2-24)
This is known as the Kohn-Sham (KS) equation, where all the terms are known except 𝑉 𝑥𝑐 (𝒓). We thus need to introduce an approximate exchange-correlation functional. The accuracy of the electronic density of the system investigated will depend on the choice of this functional and how close it comes to the exact value [START_REF] Perdew | Prescription for the design and selection of density functional approximations: More constraint satisfaction with fewer fits[END_REF]. The most common approximations are the local-density approximation (LDA) [START_REF] Kohn | Self-Consistent Equations Including Exchange and Correlation Effects[END_REF] and the generalized gradient approximation (GGA) [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF][START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF], which are described below.

Local-Density approximation (LDA)

The local density approximation (LDA) is one of the simplest approximations. In the local-density approximation, the density is considered locally uniform [START_REF] Perdew | Self-interaction correction to density-functional approximations for many-electron systems[END_REF]. The exchange-correlation energy at every position is considered to be the same as it would be for a homogeneous electron gas having the same density. Thus, the total exchange correlation energy of the system can be written as:

𝐸 𝑥𝑐 𝐿𝐷𝐴 [𝜌] = ∫ 𝜌(𝒓)𝜖 𝑥𝑐 [𝜌(𝒓)]𝑑𝒓 (2-25)
Where 𝜖 𝑥𝑐 [𝜌(𝒓)] is the exchange-correlation energy per particle for a uniform gas density 𝜌(𝒓). The Generalized Gradient Approximation (GGA) is an improved form based on the LDA functional. In the generalized gradient approximation, the non-homogeneity of the local density is taken into account by including the gradient of the electron density through a general function f depending on 𝜌(𝒓) and ∇𝜌(𝒓) [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF].

𝐸 𝑋𝐶 𝐺𝐺𝐴 [𝜌] = ∫ 𝑓(𝜌(𝒓), ∇𝜌(𝒓))𝑑𝒓 (2-26)
In this work, the GGA in the Perdue-Burke-Ernzerhof (PBE) approach [START_REF] Perdew | Generalized Gradient Approximation Made Simple[END_REF] was used for the structure optimization and phonon calculations. For electronic band structure calculations, the conventional functionals such as LDA and GGA always underestimate electronic band gaps. In order to overcome this issue and at a lower cost of simulation (compare to GW "many-body" corrections), the Heyd-Scuseria-Ernzerhof (HSE) hybrid functional and TB-mBJ (Tran-Blaha modified Becke-Johnson) potential were chosen for the band structure calculations.

Heyd-Scuseria-Ernzerhof (HSE) hybrid functional

GGA functional makes DFT a useful tool in computational physics and chemistry, and hybrid density functionals improve the GGA results by taking into account a given fraction of Hartree-Fock (HF) exchange. However, the expensive computation cost remains an issue. In order to improve computational efficiency, Heyd, Scuseria, and Ernzerhof (HSE) proposed an accurate and efficient screened hybrid functional-HSE hybrid functional based on the PBE functional, which uses an error-function-screened Coulomb potential to calculate the exchange portion of the energy [START_REF] Heyd | Hybrid functionals based on a screened Coulomb potential[END_REF].

As the Coulomb operator can be separated into long-range (LR) and short-range (SR) parts, all the terms of the PBE exchange energy can be splitted into SR and LR components. The LR components of the HF exchange have only very small contributions to this functional and it can be compensated at some point by the description of the exchange energy in the PBE approach [START_REF] Heyd | Hybrid functionals based on a screened Coulomb potential[END_REF]. Therefore, the HSE functional is finally expressed as:

𝐸 𝑥𝑐 𝜔𝑃𝐵𝐸ℎ = 𝑎𝐸 𝑥 𝐻𝐹,𝑆𝑅 (𝜔) + (1 -𝑎)𝐸 𝑥 𝑃𝐵𝐸,𝑆𝑅 (𝜔) + 𝐸 𝑥 𝑃𝐵𝐸,𝐿𝑅 (𝜔) + 𝐸 𝑐 𝑃𝐵𝐸 (2-27)
Where 𝑎 is the mixing parameter, and 𝜔 is an adjustable parameter for controlling the short-range interaction. 𝐸 𝑥 𝐻𝐹,𝑆𝑅 (𝜔) is the short-range Hartree-Fock exact exchange functional, 𝐸 𝑥 𝑃𝐵𝐸,𝑆𝑅 (𝜔) and 𝐸 𝑥 𝑃𝐵𝐸,𝐿𝑅 (𝜔) are the short-and long-range components of the PBE exchange functional, and 𝐸 𝑐 𝑃𝐵𝐸 (𝜔) is the PBE correlation functional.

In this work, we used HSE functional to study the electronic band structure of III-V materials. Firstly, the band structure of simple zinc-blende III-V binary semiconductors unit cells (including InP, GaP, GaSb and GaAs) were calculated by HSE with different 𝜔.

From the comparison with bandgaps determined experimentally, the 𝜔 was fixed, and then used for further band structure calculations of more complex structures.

TB-mBJ potential

Even if the HSE functional gives similar results that the traditional HF calculations, the calculation is much faster. But still, for large structures with dense k points, the HSE calculation remains quite expensive (in terms of computational cost) and time-consuming.

Therefore, for the calculations of large structures with dense k points, HSE functional is not a good choice. On the other hand, as mentioned above, the calculations with GGA and LDA functionals are fast and economical but greatly underestimate the bandgaps for many semiconductors. In order to improve the calculation accuracy of the bandgaps in DFT, an alternative system was proposed by Tran and Blaha [START_REF] Tran | Accurate Band Gaps of Semiconductors and Insulators with a Semilocal Exchange-Correlation Potential[END_REF] by modifying the Becke-Johnson (BJ) potential. The Becke-Johnson potential is expressed as [START_REF] Becke | A simple effective potential for exchange[END_REF]:

𝑉 𝑥 𝐵𝐽 (𝒓) = 𝑉 𝑥 𝐵𝑅 (𝒓) + 1 𝜋 √ 5 6 √ 𝑡(𝒓) 𝜌(𝒓) (2-28)
Where 𝑡(𝒓) is the density of the kinetic energy:

𝑡(𝒓) = 1 2 ∑ ∇∅ 𝑖 * (𝒓)∇∅ 𝑖 (𝒓) 𝑁 𝑖=1
(2-29)

57

𝑉 𝑥 𝐵𝑅 (𝒓) is the Becke-Roussel (BR) potential [START_REF] Becke | Exchange holes in inhomogeneous systems: A coordinatespace model[END_REF]:

𝑉 𝑥 𝐵𝑅 (𝒓) = - 1 𝑏(𝒓) (1 -𝑒 -𝑥(𝒓) - 1 2 𝑥(𝒓)𝑒 -𝑥(𝒓) ) (2-30) 𝑏(𝒓) = [ 𝑥 3 (𝒓)𝑒 -𝑥(𝒓) 8𝜋𝜌(𝒓) ] 1 3
(2-31)

Tran and Blaha introduced a parameter "c" to modify the relative weights of the two terms of the BJ potential and the modified potential (TB-mBJ) are shown as [START_REF] Tran | Accurate Band Gaps of Semiconductors and Insulators with a Semilocal Exchange-Correlation Potential[END_REF]:

𝑉 𝑥 𝑇𝐵-𝑚𝐵𝐽 (𝒓) = 𝑐𝑉 𝑥 𝐵𝑅 (𝒓) + (3𝑐 -2) 1 𝜋 √ 5 6 √ 𝑡(𝒓) 𝜌(𝒓) (2-32)
TB-mBJ potential was proved to reach similar accuracy than computationally more expensive functionals, such as HSE. This approach is only slightly more expensive than conventional LDA and GGA calculations. Therefore, it is a suitable method for the DFT calculations performed over large systems or with very dense k points.

Self-Consistency and structure optimization

The Self-Consistency loop is shown in the Fig. 2.17 (right part). Firstly, an initial guess for electron density 𝜌 𝑖𝑛 (𝒓) is needed to calculate Hartree potential 𝑉 𝐻 (𝒓) and exchange-correlation potential 𝑉 𝑥𝑐 (𝒓), which will be further used to solve KS Equation.

Once a new set of wave functions 𝜑 𝑖 𝑜𝑢𝑡 (𝒓) is obtained by solving the non-interaction electron problem in the external potential 𝑉 𝑒𝑥𝑡 , a new electron density can be constructed

as 𝜌 𝑜𝑢𝑡 (𝒓) = ∑ |𝜑 𝑖 𝑜𝑢𝑡 (𝒓)| 2 𝑁 𝑖=1
. The self-consistency is reached if 𝜌 𝑜𝑢𝑡 (𝒓) = 𝜌 𝑖𝑛 (𝒓). If this is not the case, the procedure will be iterated with a new guess for electron density constructed from 𝜌 𝑜𝑢𝑡 (𝒓). For the structure optimization, the self-consistent calculation is done for each new atomic position. The algorithm of relaxation will stop at specific force criteria (as shown in the left part of Fig. 2.17). This method is used for system such as bulk material or slabs for surfaces and interfaces studies. In this thesis, for structure optimization, the structures were considered as relaxed until the Hellmann-Feynman forces on the atoms were less than 10 -4 eV/Å before any other subsequent calculations.

58 Figure 2.17 Flow chart of the iteration scheme. The scheme on the left corresponds to DFT selfconsistent calculation for relaxation and structure optimization. The scheme on the right, in the red frame, is the self-consistent total energy calculation [START_REF] Lucci | Surface and interface contributions to III-V/Si hetero-epitaxial growth: Theory and Experiments[END_REF].

Computational codes

The density functional theory has been implemented in various computational codes.

In this thesis, most DFT calculations (including electronic band structure and phonon calculations) were performed with the Vienna Ab-initio Simulation Package (VASP) code [START_REF] Kresse | Efficiency of ab-initio total energy calculations for metals and semiconductors using a plane-wave basis set[END_REF]. The VASP code is by far one of the most popular computational codes for calculations of infinite systems such as solid-state materials and condensed matter.

The self-consistency cycles and choices of numerical methods used to calculate the electronic Kohn-Sham ground state in VASP are quite efficient and robust. VASP code uses plane wave basis set rather than localized basis set. The disadvantage of using plane waves is that it requires rapid oscillation of the wave function near the core region. For the core electrons, it is necessary to use a large number of plane waves and a large plane wave cut off. For reducing the computational workload, in VASP code, pseudopotentials or the Projector Augmented-Wave (PAW) method [START_REF] Blöchl | Projector augmented-wave method[END_REF] are used to describe the 59 interaction between ions and non-valence electrons so that the size of the basis set can be reduced. The main idea of the pseudopotential method is that the core electrons are not considered and therefore the rapid oscillations of the valence wave functions around the nucleus, which causes extremely large basis set, are avoided. Indeed, the core electrons are often considered as chemically inert and chemical bondings mostly involve the valence electrons. For the PAW method, the potentials can be defined as the frozen core all-electron potentials, which unifies all-electron and pseudopotential approaches. In order to improve the accuracy, the PAW method also takes the core wave functions into account. It divides the wavefunction in two parts, with a partial wave in a sphere around the atom called the augmentation region and outside the sphere called the interstitial region. Inside the augmentation region, the all-electron wave function can be expanded.

These two terms are then added and the overlapping part is subtracted to obtain a wave function close to the all-electron wave function.

Periodic Boundary Conditions (PBCs), in all spatial directions, are imposed in the VASP code to treat infinite number of atoms. The Bloch functions are used to take into account the system periodicity:

𝜑 𝒌 (𝒓) = 𝑒 𝑖𝒌.𝒓 𝑢 𝒌 (𝒓) (2-33)
Where 𝑢 𝒌 (𝒓) is a function with the lattice periodicity 𝑢 𝒌 (𝒓) = 𝑢 𝒌 (𝒓 + 𝑹), so that:

𝜑 𝒌 (𝒓 + 𝑹) = 𝑒 𝑖𝒌.𝑹 𝑢 𝒌 (𝒓) (2-34)
Where 𝒌 is the reciprocal lattice vector in the Brillouin zone and 𝑹 is the translation vector of the lattice. During the computation, only the atoms within one periodic cell are optimized and computed taking into account their interactions with atoms in the neighboring cells.

The package also includes the Monkhorst-Pack [START_REF] Monkhorst | Special points for Brillouin-zone integrations[END_REF] to sample the Brillouin zone.

The VASP offers support for a large variety of platforms. Besides, high performance computation (parallel computing) was used in this work.

In addition, the band symmetry properties and electron-phonon interaction were studied by using Quantum ESPRESSO (QE). QE is an open-source software [START_REF] Giannozzi | QUANTUM ESPRESSO: a modular and open-source software project for quantum simulations of materials[END_REF].

Similar to VASP, QE also uses planewave basis sets combined with pseudopotentials or the PAW method to achieve DFT electronic band structure calculations. Periodic boundary conditions are also imposed in this code to treat infinite number of atoms. As compared to VASP, the calculations by QE need longer time generally. Besides, implementation of the hybrid functionals (HSE, in particular) is easier with VASP than with QE. But QE also has its specific advantages. For instance, detailed symmetry information of electronic states can be extracted from the band structure calculations and it also has advanced modules for electron-phonon interaction calculations.

Electronic band structure calculation

The band structure calculations in this work were performed with the PAW method.

The calculations contain three main steps. Firstly, self-consistent field (SCF) calculation (as shown in Fig. 2.17) is needed to obtain a self-consistent electron density using a

61

Besides, the band folding effect will occur in band structures for the supercell [START_REF] Yang | Symmetry demanded topological nodal-line materials[END_REF].

Phonon calculation

There are two main methods for phonon dispersion calculations: Finite displacement method and Linear Response method. In this work, we chose the Finite displacement method to do phonon calculations (the results are shown in Chapter 3) combining VASP and PHONOPY software [START_REF] Togo | First principles phonon calculations in materials science[END_REF].

One way to implement the theory of crystal vibrations for realistic structures is to explicitly calculate the forces between every atom in the crystal and construct the force constant matrix of the crystal. This force constant matrix then allows us to calculate the 62 normal modes at any specific wavevector, q. In order to calculate the forces caused by the displacement of an atom i, one can shift this atom i, and then use DFT to calculate the forces on every other atoms using the Hellman-Feynman theorem [START_REF] Feynman | Forces in Molecules[END_REF]. This method of calculating the force constant matrix by explicitly displacing atoms is called the Finite displacement method. This method is generally quicker and computationally cheaper than the Linear Response method [START_REF] Giannozzi | An introduction to linear response and to phonon calculations[END_REF], which uses density functional perturbation theory to calculate forces.

The phonon calculations by Finite displacement method use the following steps:

1. Creation of supercell structures with (or without) displacements from a unit cell.

2. Calculation of sets of forces.

3. Calculation of Force constants from the sets of forces.

4. Building of Dynamical matrix from the force constants.

5. Determination of phonon frequencies and eigenvectors from the dynamical matrices with specified q-points.

Calculation of the electron-phonon interaction (EPI) matrix elements

In the classic Fröhlich analytical model, the electron phonon coupling matrix element is evaluated in the case of an isotropic interaction of an electron in a parabolic band with a single dispersionless LO phonon mode [140]:

𝑔 𝑞 = 𝑖 |𝑞| [ 𝑒 2 4𝜋𝜀 0 4𝜋 𝑁Ω ℏ𝜔 𝐿𝑂 2 ( 1 𝜖 ∞ - 1 𝜖 0 )] 1/2 (2-35)
Where q is the phonon wave vector, Ω the unit cell volume, N the number of unit cells in the Born-von Kármán supercell, and e, 𝜖 0 , 𝜖 ∞ and ħ are the electron charge, static and high-frequency dielectric constants, and reduced Planck constant, respectively. As shown by Eq.2-35, the Fröhlich coupling constant 𝑔 𝑞 diverges at long wavelengths (q → 0) following a simple 1/q variation.

63

In the context of DFT, EPI matrix elements (𝑔 𝑚𝑛𝑣 (𝑘, 𝑞)) can be calculated using a state-of-the-art numerical implementation of the general microscopic theory [START_REF]School on Electron-Phonon Physics from First Principles[END_REF][START_REF] Verdi | Electron-Phonon Vertex from First Principles[END_REF][START_REF] Giustino | Electron-phonon interactions from first principles[END_REF].

This theory generalizes the classic Fröhlich approach to anisotropic materials and multiple phonon branches by separating the short-(S) and long-range (L) contributions to the matrix elements [START_REF] Verdi | Electron-Phonon Vertex from First Principles[END_REF]:

𝑔 𝑚𝑛𝑣 (𝑘, 𝑞) = 〈𝛹 𝑚𝑘+𝑞 |∆ 𝑞𝑣 𝑉|𝛹 𝑛𝑘 〉 = 𝑔 𝑚𝑛𝑣 𝑆 (k,q) + 𝑔 𝑚𝑛𝑣 𝐿 (𝑘, 𝑞) (2-36)
Where the two electronic states 𝛹 𝑛𝑘 and 𝛹 𝑚𝑘+𝑞 are defined by the m and n indices for the electronic dispersion curves, as well as the wavevectors k+q and k. The phonon branch is defined by 𝑣 and q is the phonon wavevector.

The EPI matrix element has the meaning of a probability amplitude for the scattering between the initial electronic state |𝛹 𝑛𝑘 〉 and the final state 〈𝛹 𝑚𝑘+𝑞 | via the perturbation ΔqνV. The long-range (L) matrix element 𝑔 𝑚𝑛𝑣 𝐿 (𝑘, 𝑞) represents the generalization of the Fröhlich vertex for ab initio calculations [START_REF] Verdi | Electron-Phonon Vertex from First Principles[END_REF].

𝑔 𝑚𝑛𝑣 𝐿 (𝑘, 𝑞) = 𝑖 4𝜋 Ω 𝑒 2 4𝜋𝜀 0 ∑ ( ℏ 2𝑁𝑀 𝑘 𝜔 𝑞𝑣 ) (1/2) 𝑘 × ∑ (𝑞 + 𝐺) • 𝑍 𝑘 * • 𝑒 𝑘𝑣 (𝑞) (𝑞 + 𝐺) • 𝜖 ∞ • (𝑞 + 𝐺) 𝐺≠-𝑞 × 〈𝛹 𝑚𝑘+𝑞 |𝑒 𝑖(𝑞+𝐺)•𝑟 |𝛹 𝑛𝑘 〉 (2-37)
Where the angular brackets have to be evaluated within the Born-von Kármán supercell.

This formula shows that the increase of this contribution at long wavelengths, q → 0, is more complex than in the case of the classic Fröhlich analytical model [START_REF] Verdi | Electron-Phonon Vertex from First Principles[END_REF].

In this thesis, the EPI matrix element was calculated based on QW code and the results are shown in Chapter 3. The symmetry properties of electronic states and electron-phonon interaction matrix elements were calculated and extracted by using metaGGA (TB-mBJ) potentials and LDA exchange-correlation functional, respectively, using the QE package code.

For the DFT calculations of nonstoichiometric APB structures, III-III and V-V nonstoichiometric APB structures with GaP, GaAs and GaSb materials were constructed.

In order to investigate separately the III-III and V-V APB singularities, slab structures were constructed by adding vacuum on both sides and the surfaces were passivated with fictitious hydrogen atoms to avoid contribution of localized surface states. Two kinds of nonstoichiometric APB structures with relatively short and long lengths were constructed Wannier functions within HSE potential [START_REF] Franchini | Maximally localized Wannier functions in LaMnO3 within PBE+ U, hybrid functionals and partially self-consistent GW: an efficient route to construct ab initio tight-binding parameters for eg perovskites[END_REF] as implemented in Wannier Tools package [START_REF] Wu | WannierTools: An open-source software package for novel topological materials[END_REF]. The Bloch wave functions are projected onto the s, p, d atomic orbitals of Ga and s, p atomic orbitals of P, As and Sb. The symmetry properties of the electronic states were analyzed starting from the irreducible representations provided by Quantum Espresso package codes.

Summary

In this chapter, firstly, the different experimental characterization techniques involving structural (AFM, C-AFM, XRD, STEM, CMP), optical (PL, Raman, ellipsometry) and PEC characterizations have been presented with the related physical mechanisms, experimental setups and parameter settings used in this thesis. Then DFT calculations have been described with a brief introduction to the basic concepts of the DFT, followed by the calculation steps and parameter settings for the determination of electronic, vibrational and electron-phonon coupling properties.

Chapter 3: Optoelectronic and vibrational properties of stoichiometric anti-phase boundaries

The initial objective of my thesis was to explore the optoelectronic properties of III-V/Si epilayers. The preliminary observations of unusual PL features on III-V/Si samples rapidly showed the need for a deep understanding of optoelectronic and vibrational properties of these samples, which, as we understood, could only be driven by the presence of some stoichiometric antiphase boundaries. This led to the demonstration of a strong electron-phonon coupling in these 2D nano-objects, published in 2020 in ref. [START_REF] Chen | Strong Electron-Phonon Interaction in 2D Vertical Homovalent III-V Singularities[END_REF].

The results presented hereafter have been obtained with the participation of many members of the OHM research team. In addition, group IV sample growth and STEM/EDX measurements have been performed in collaboration with IHP microelectronics (Germany). The Raman measurements were achieved in collaboration with both IHP microelectronics and IPR (Institut de Physique de Rennes).

Introduction

The electron-phonon interaction has a strong impact on electrical, optical and thermal properties of solids [START_REF] Yu | Engineering Acoustic Phonons and Electron-Phonon Coupling by the Nanoscale Interface[END_REF][START_REF] Jin | Interlayer electron-phonon coupling in WSe2/hBN heterostructures[END_REF]. Especially, strong electron-phonon interactions were recently investigated for materials and devices developments as well as for understanding of fundamental physical processes, such as phonon lasing [START_REF] Czerniuk | Lasing from active optomechanical resonators[END_REF],

superconductivity [START_REF] Lanzara | Evidence for ubiquitous strong electron-phonon coupling in hightemperature superconductors[END_REF] or colossal magnetoresistance [START_REF] Driza | Long-range transfer of electron-phonon coupling in oxide superlattices[END_REF]. The strength of this interaction and the formation of a polaron however highly depends on the polarizability of the chemical bonds composing the material [START_REF] Byrnes | 1 Basic theory and phenomenology of polarons[END_REF]. Therefore, strong electron-phonon interaction was preferentially studied in highly polar materials, such as alkali halides [START_REF] Byrnes | 1 Basic theory and phenomenology of polarons[END_REF], oxides [START_REF] Driza | Long-range transfer of electron-phonon coupling in oxide superlattices[END_REF][START_REF] Wang | Tailoring the nature and strength of electron-phonon interactions in the SrTiO3 (001) 2D electron liquid[END_REF] or perovskites [START_REF] Wright | Electron-phonon coupling in hybrid lead halide perovskites[END_REF]. The strength of this interaction is commonly quantified with the so-called Huang-Rhys factor S, whose maximum value may reach 100

68
with alkali halides [START_REF] Leung | On the luminescence quenching of F centres in alkali halides[END_REF]. On the other hand, group-II-VI, III-V or IV semiconductors, have been widely used in photoelectric device applications due to their excellent optical and/or transport properties. When these bulk materials are pure and undoped, the electronphonon interaction remains generally weak, reflected by quite low Huang-Rhys (S) factor (e.g. 0.0033 for InAs [START_REF] Bissiri | Optical evidence of polaron interaction in InAs/GaAs quantum dots[END_REF] and 0.035 for CdSe [START_REF] Gindele | Phonon interaction of single excitons and biexcitons[END_REF]), as the chemical bonds are either covalent (group IV) or slightly polar (group III-V or II-VI). Effective enhancement of the electron-phonon interaction was successfully obtained by using spatial confinement or carrier localization effects [START_REF] Heitz | Enhanced Polar Exciton-LO-Phonon Interaction in Quantum Dots[END_REF]. A. Alkauskas et al. comprehensively and deeply discussed the electron-phonon interaction in III-V semiconductors defects [START_REF] Alkauskas | Tutorial: Defects in semiconductors-Combining experiment and theory[END_REF].

Especially, punctual antisite defects were found to greatly enhance the electron-phonon interaction [START_REF] Fockele | Arsenic antisite defects in AlxGa1-xAs observed by luminescence-detected electron-spin resonance[END_REF]. M. Kang et al. discovered Holstein polaron in surface-doped layered

MoS2 semiconductor [START_REF] Kang | Holstein polaron in a valley-degenerate two-dimensional semiconductor[END_REF]. Electron-phonon coupling was also investigated in many different localized systems such as nanoscale core-shell structures [START_REF] Yu | Engineering Acoustic Phonons and Electron-Phonon Coupling by the Nanoscale Interface[END_REF], graphene and InAs double quantum dot (DQD) systems [START_REF] Roulleau | Coherent electron-phonon coupling in tailored quantum systems[END_REF], WSe2/hBN heterostructure [START_REF] Jin | Interlayer electron-phonon coupling in WSe2/hBN heterostructures[END_REF] or graphene/Monolayer h-BN structure [START_REF] Brar | Hybrid Surface-Phonon-Plasmon Polariton Modes in Graphene/Monolayer h-BN Heterostructures[END_REF].

In this chapter, we demonstrate strong electron-phonon interactions in 2D vertical

In-rich Indium gallium phosphide ((In,Ga)P) stoichiometric APB, where the homovalent bonds drastically change the local polarity and charge distribution in the crystal, leading to efficient phonon confinement and localized holes wavefunctions. It is furthermore surrounded by a Ga-rich (In,Ga)P barrier enabling the quantum confinement of carriers.

The unprecedented electron-phonon coupling strength measured in these buried and nonoxidized III-V vertical singularities opens many perspectives to understand the physics of the electron-phonon interaction in semiconductors.

Sample growth

The preparation of the studied In0.3GaP on SiGe/Si sample contains two main steps.

The first step is the growth of a relaxed SiGe buffer layer on Si substrate [START_REF] Skibitzki | Lattice-engineered Si1-xGex-buffer on Si(001) for GaP integration[END_REF], performed by our German partner (IHP-microelectronics). A 760nm Si0.433Ge0.567 layer was grown

Structural characterizations and post-growth processing

As discussed in Chapter 1, APBs are commonly formed during the epitaxy of Zinc-Blende III-V semiconductors on a group IV substrate in the diamond phase (e.g. silicon (Si) or Germanium (Ge)). It originates from the two different ways for the III-V crystal to fit the group IV substrate orientation [3]. As a consequence, two domains with different phases may grow on the same substrate, separated by wrong III-III or V-V bonds, inside the III-V matrix, as illustrated in Fig. 3.1. This leads to the insertion of a local translational symmetry breaking. III-III (V-V) bond also adds 2 free holes (2 free electrons) in the semiconductor. However, unlike doping, APBs extend in the form of vertical 2D planes (at least locally), that can emerge at the sample surface. From this description, APB can thus be considered as 2D homovalent singularities in a III-V bulk matrix. As described in Chapter 1, theoretically, the APB composition may range from a pure vertical III-III APB to a pure V-V APB, resulting in these two extreme cases to a large excess of charges released in the lattice. But stoichiometric APBs (with equal numbers of III-III and V-V bonds within the same APB), depicted in Fig. 3.1, have been mostly considered, due to their relatively low formation energy resulting from the charge compensation effect [START_REF] Rubel | Formation Energies of Antiphase Boundaries in GaAs and GaP: An ab Initio Study[END_REF].

70

We will see in Chapter 5 that it is also the only configuration that leads to the presence of a semiconducting bandgap, and thus the only possible additional contribution to the optical properties of the layer studied. in Fig. 3.2b (black circular areas) and in Fig. 3.3f, arising from demixing due to the nonequilibrium growth conditions used [START_REF] Yuan | Interplay Effect of Temperature and Excitation Intensity on the Photoluminescence Characteristics of InGaAs/GaAs Surface Quantum Dots[END_REF]. Removal of these surface QDs will be discussed hereafter. As described previously, APBs are composed of III-III or V-V bonds [START_REF] Rubel | Formation Energies of Antiphase Boundaries in GaAs and GaP: An ab Initio Study[END_REF][START_REF] Tea | Theoretical study of optical properties of anti phase domains in GaP[END_REF], which induce a lack or excess of electrons, as compared to the surrounding matrix. Therefore, it is expected to behave as acceptor or donor-like centers [START_REF] Rubel | Formation Energies of Antiphase Boundaries in GaAs and GaP: An ab Initio Study[END_REF]. One of the first consequences of the intrinsic APB structure is that inversion symmetry, which is absent in the bulk III-V zinc-blende lattice is restored locally. An InP supercell containing two stoichiometric APB planes (called hereafter APB-supercell) with 32 atoms (n=32) is shown in Fig. 3.11.

For comparison, a reference InP supercell structure (n=32) without APB (called hereafter bulk-supercell) was also constructed as shown in Fig. 3.12. In order to better understand and analyze the vibrational and electronic properties in the system, phonon and band structures need to be clarified. Electronic band structure and phonon calculations were performed using Density Functional Theory (DFT) for InP APB-supercell and bulk-supercell. The corresponding results and discussions will be shown in the following paragraphs. In order to identify the electronic states strictly related to the isolated APB, calculations were performed on APB supercells with n=32/48/64 atoms (with increasing number of atoms only along the [1-10] direction, leading to longer distance between the nearby APBs) based on metaGGA (TB-mBJ) [START_REF] Tran | Accurate Band Gaps of Semiconductors and Insulators with a Semilocal Exchange-Correlation Potential[END_REF] potential by Quantum Espresso package codes to obtain the symmetry properties of electronic states and get rid of finite size effects which lead to artificial electronic coupling between the APBs.

Vibrational properties calculations

Phonon dispersions of the APB-supercell and the reference bulk-supercell (InP) are then computed and shown in Fig. 3. [START_REF] Vurgaftman | Band parameters for III-V compound semiconductors and their alloys[END_REF] and Fig. 3.19, respectively, where R(0.5,0.5,0.5), Г(0,0,0) and U(0.5,0,0.5) are the high symmetry points of the first Brillouin zone. In both cases, the phonon dispersion also shows band folding effects due to the construction of the supercell. From the known bulk phonon dispersions, we can identify the contributions of underlying InP acoustic (low frequency) and optical (high frequency) phonon branches.

We may notice that the computed bulk phonon frequencies are a little smaller than the experimental ones [START_REF] Lindsay | Ab initio thermal transport in compound semiconductors[END_REF], which is an expected consequence of the generalized gradient approximation (GGA) that overestimates lattice constants and underestimates the energies of vibrational modes [START_REF] Haas | Construction of an optimal GGA functional for molecules and solids[END_REF].

Comparing the phonon dispersions of the bulk (Fig. 3.19) and APB-structure (Fig.

3.18

), it can be clearly observed that the APB-supercell introduces several new flat phonon branches in the phonon dispersion gap, positioned at around 5.3 THz, 5.9 THz, 7.8 THz, 8.0 THz, 12.6 THz, marked by orange, blue, magenta, green and purple dotted lines, respectively (Fig. 3.18). These phonons correspond at the Г point, respectively, to the Ag, Bu, Au, Bu, Ag irreducible representations of the layer group for a single APB. The symmetry properties were investigated for these additional phonon modes based on an APB supercell of 32 atoms at the DFT-GGA level (Table 3.1). The phonon modes are all grouped by pairs since the supercell contains 2 APBs, but the phonons can be labelled as well, according to the IR of the layer group, which are relevant for a single APB (Fig. 3.13). The small splitting for each pair of phonons is related to the mechanical coupling between APBs that should vanish for an infinite supercell. Therefore, the five new phonons characteristic of a single APB are identified with frequencies equal to about 5.30, 5.92, 7.84, 7.96, 12.6 THz and well-defined symmetries (Table 3.

Optical characterizations

Raman measurement

Spatially-and spectrally-resolved Raman measurements were performed on the sample after CMP and the results are shown in Fig. 3.23. The Raman peaks 2, 3, 4, 5, 6, 7, correspond respectively to Ge-Ge bonds [START_REF] Ravaux | Effect of rapid thermal annealing on crystallization and stress relaxation of SiGe nanoparticles deposited by ICP PECVD[END_REF], GaP transverse optical (TO) phonons [START_REF] Peternai | Investigation of graded InxGa1-xP buffer by Raman scattering method[END_REF], InP longitudinal optical (LO) phonons [START_REF] Peternai | Investigation of graded InxGa1-xP buffer by Raman scattering method[END_REF], GaP LO phonons [START_REF] Peternai | Investigation of graded InxGa1-xP buffer by Raman scattering method[END_REF], Ge-Si bonds [START_REF] Ravaux | Effect of rapid thermal annealing on crystallization and stress relaxation of SiGe nanoparticles deposited by ICP PECVD[END_REF] and Si-Si bonds [START_REF] Peternai | Investigation of graded InxGa1-xP buffer by Raman scattering method[END_REF], and the InP TO phonons [START_REF] Peternai | Investigation of graded InxGa1-xP buffer by Raman scattering method[END_REF] should give some contribution to the signal around 303 cm -1 . The Raman peak 1 is relatively broad and weak, but its

Ellipsometry Measurements

The

Photoluminescence measurements

Low temperature (15 K) PL experiments were performed for the sample after CMP, as shown in Fig. 3.25. In order to analyze the PL and absorption spectra comprehensively, the absorption spectrum (Fig. 3.24c) was shifted at higher energy by 66 meV to account for the Varshni temperature dependence of the bandgap [START_REF] Varshni | Temperature dependence of the energy gap in semiconductors[END_REF] and put together with the low temperature PL spectrum (Fig. 3.25). In the PL spectrum, one broad PL peak is observed between 1.8 and 2.1 eV. In the absorption spectrum, one peak between 3.5 and 4.0 eV can be observed, with a similar shape than the one usually reported for GaP [START_REF] Aspnes | Dielectric functions and optical parameters of Si, Ge, GaP, GaAs, GaSb, InP, InAs, and InSb from 1.5 to 6.0 eV[END_REF],

indicating that the absorption spectrum mainly comes from the In-poor phase. But the absorption curve starts from the high energy side of the PL around 2.1 eV, which is smaller than the expected bandgap of the In-poor phase In0.18GaP at low temperature (which should be around 2.32 eV) [START_REF] Zhang | InP-GaAs Photodiodes Tailored for Specific Wavelength Range[END_REF]. This absorption tail can result from the alloy fluctuation of the InGaP matrix or another part of the material with a lower bandgap (i.e. the In-rich phase with the APB singularity). The k refers to the exponent of the power law equation IPL ~Iex k , where IPL is the integrated photoluminescence intensity and Iex is the excitation laser power [START_REF] Schmidt | Excitation-power dependence of the near-bandedge photoluminescence of semiconductors[END_REF]. It is found that at low pumping power it shows a super-linear power-dependency (k=1.3). With the increase 91 of the power, the k value decreases and finally turns to linear dependency (k=0.96). This behavior is a typical characteristic of the emissions involving saturation effect of trap centers (acceptors or donors) [START_REF] Bimberg | Bound exciton luminescence in epitaxial Sn-doped gallium-arsenide[END_REF][START_REF] Hellmann | Low-temperature anti-Stokes luminescence mediated by disorder in semiconductor quantum-well structures[END_REF].

I A (T) = I 0 1+𝑎 1 e -𝐸 1 k B T ⁄ + I 0 1+𝑎 2 e -𝐸 2 k B T ⁄ + I 0 1+𝑎 3 e -𝐸 3 k B T ⁄ (3-1)
Here, I0 corresponds to the initial integrated PL intensity at the lowest temperature T=15 K, 𝑎 1 , 𝑎 2 , 𝑎 3 are dimensionless temperature-independent constants, and kB refers to the Boltzmann constant.

The first activation energy is attributed to alloy disorder [START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF], while the two others may represent the detrapping of localized carriers (electrons or holes) to other localized states or extended states. Following the results obtained by DFT, the 200 meV activation energy is considered to be related to the detrapping of the localized holes in the APB 92 states to extended states of the In-rich phase. In addition, time-resolved PL was measured. Figure 3.28 shows the PL decay curve at 15 K, excited by a 355 nm pulse laser at a 560 mW/cm 2 excitation density. The PL decay curve was fitted by a power decay model designed for type II system [START_REF] Mourad | Determination of valence-band offset at cubic CdSe/ZnTe type-II heterojunctions: A combined experimental and theoretical approach[END_REF]:

𝐼(𝑡) = 𝛾𝑛 0 2 (1+𝛾𝑛 0 𝑡) 3/2 (3-2)
Where 𝑛 0 corresponds to the carrier density 𝑛 at t = 0 and 𝛾 is a constant that gives the recombination rate  = (n)= 𝛾𝑛. The fitting curve is shown as the red curve in

Proposed mechanism for the PL with multi-phonon replicas

Based on the DFT calculations and experimental results above, the general mechanism of the PL resulting from the 2D homovalent singularity is schematically represented in Fig. 3.29a. After illumination, the excited electrons in the conduction band recombine with the holes localized in the APB plane and it leads to PL emission. The mechanisms of electron-phonon interaction leading to the phonon replicas in the PL spectrum are analyzed first following the simplified Franck-Condon approach (Fig.

3.29b).

The configuration coordinate diagram is used to show the potential well vs the normal mode coordinate (Q). Configuration coordinate diagram is the best way to describe the change of the defect geometry associated with the change of its electronic state [START_REF] Stoneham | Theory of Defects in Solids: Electronic Structure of Defects in Insulators and Semiconductors[END_REF], which in our simple picture is proportional to the bond length [START_REF] Alkauskas | Tutorial: Defects in semiconductors-Combining experiment and theory[END_REF]. The dotted horizontal lines represent oscillator states and in this illustration, the potential well is ideally considered as harmonic so that the energy difference between two nearby oscillator states are the same and equal to the phonon energy.

Fitting of the PL spectrum with multi-phonon replicas

The broad PL was fitted based on the Eq. 3-3. The phonon replicas were fitted by Gaussian functions with full width at half maximum (FWHM) of 26.5 meV and a phonon energy of around 24 meV was used. Note that Lorentzian functions did not allow fitting the phonon replicas correctly. The Gaussian broadening observed is mainly considered as the consequence of the inhomogeneous broadening effect due to the composition inhomogeneity or the different distances between the APBs [START_REF] Alkauskas | Tutorial: Defects in semiconductors-Combining experiment and theory[END_REF]. The broadening effect due to acoustic phonons shall amount only to a few meV range [START_REF] Besombes | Acoustic phonon broadening mechanism in single quantum dot emission[END_REF], which can be neglected in the present study. In addition, the energy difference between two nearby phonon lines in the low energy part of the spectrum becomes a bit smaller gradually, indicating that the potential well is slightly anharmonic. Anyway, this observation doesn't impact the conclusions given here.

Electron-phonon interaction calculation

To further confirm the strong electron-phonon coupling induced by the new localized phonon modes around the APB structure, we calculated electron-phonon interaction (EPI) matrix elements (𝑔 𝑚𝑛𝑣 (𝑘, 𝑞)) for both polar and non-polar-localized modes using state of the art DFT-based microscopic theory. to the coupling at the  point between the Ag localized non-polar phonon modes and, the bottom of the conduction band (m=n=CBM) or the top of the valence band (m=n=VBM).

We can find that the EPI matrix elements are large and consistent with the large S parameter. On the other hand, Figure 3.32 shows the variation of the EPI matrix element for the polar-localized phonon modes as a function of q coupled to the same electronic states. As described in Section 2.4.4, the matrix elements vanish at the  point (q=0).

They may exhibit an increase at long wavelengths, q→0 (Fig.

Discussion on the possible microscopic origins of the strong EPI

Finally, we shall discuss the possible microscopic origins of the observed strong coupling regime between localized carriers and phonons. Raman scattering measurements coupled with PL replica characterization and localized phonon simulations point toward a mechanism involving both 2D Ag even localized phonon modes and APB localized VB electronic states. Therefore, a specific optical deformation potential mechanism is the most probable process to explain the experimental observations [START_REF] Mourad | Determination of valence-band offset at cubic CdSe/ZnTe type-II heterojunctions: A combined experimental and theoretical approach[END_REF]. The observation of phonon replica is a signature of a strong electron-phonon interaction, which is consistent with the microscopic computation of a large electron phonon coupling matrix elements for the first localized non-polar Ag mode (Table 3.2). However, the second localized Ag mode is also predicted to exhibit a strong coupling to the same confined 99 valence state, despite the absence of corresponding discrete signatures on top of the PL envelope. Therefore, other aspects beyond our theoretical description and simplified APB model structure are likely to be important as well, including phonon-phonon interactions or lattice disorder. The Raman scattering experiment (Fig. 3.23) indeed shows that the first low-frequency localized mode already exhibits a broader signature than extended lattice modes. We may thus infer that such broadening is even more important for the other localized modes, leading only to shallow contributions to the envelope of the PL spectrum.

In addition, another classical electron-phonon coupling mechanism should be also considered for odd (polar) modes, that is the Fröhlich interaction. The theoretical study predicts that the APB polar phonon modes shall be located at a sizable energy above the lowest APB Ag even phonon mode, and therefore shall not be good candidates to explain the observed phonon replica. These localized polar modes are nevertheless likely to be coupled to localized VB electronic states. The matrix elements correspond to sizeable numerical values (Fig. 3.32) on the same order as the ones for non-polar modes at the  point (Table 3.2), thereby suggesting that the coupling to the localized polar modes is important for optoelectronic properties, and also playing a role for hole transport. This can shed some light on the recent experimental observation of APBs benefiting to carrier transport at room temperature [START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF][START_REF] Cornet | Chapter 28 -GaP/Si-Based Photovoltaic Devices Grown by Molecular Beam Epitaxy[END_REF]. Indeed, APBs are planar structures, not point defects and therefore, vertical transport in the APB planes can take place. In the present work, APBs can be assimilated to effective quantum wells forming planar conduction channels for carriers. We will however see in Chapter 5 that other mechanisms involving non-stoichiometric APBs have also to be considered to explain the observed transport properties in III-V/Si samples. Anyway, considering the recent researches that allow controlling both lateral and vertical extensions of APDs [3,[START_REF] Guillemé | Antiphase domain tailoring for combination of modal and 4¯ -quasiphase matching in gallium phosphide microdisks[END_REF], more investigations are needed to characterize the nature and the strength of the electron-phonon interaction processes in such homovalent singularities. In 2018, at the end of I. Lucci's PhD thesis, the OHM research team at the FOTON Institute realized the advantages of using III-V/Si materials for green hydrogen production, through the development of photoelectrodes [START_REF] Lucci | A Stress-Free and Textured GaP Template on Silicon for Solar Water Splitting[END_REF]. The results presented in this chapter were obtained in the framework of a collaboration initiated in 2019 with the department of Chemistry from UCL (University College London) for the characterizations of photo-electrodes. This led to the assessment of GaPSb/Si materials for photoelectrode developments, published in ref. [START_REF] Chen | Assessment of GaPSb/Si tandem material association properties for photoelectrochemical cells[END_REF], and the first demonstration of an operating bulk III-V/Si photoelectrode, published in ref. [START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF].

Introduction

As described in Chapter 1, solar water splitting technology producing clean and green hydrogen is of great interest for solving the environment and energy issues under the current global warming background.

However, numbers of challenges remain for the development of this technology, in terms of efficiency, profitability and sustainability. At the heart of the Photo-Electro-Chemical (PEC) conversion process is the choice of an appropriate photoelectrode material with both good bandgap and good band alignment to harvest the largest portion of the solar spectrum and provide sufficient voltage to accomplish the water splitting reactions [START_REF] Van De Krol | Photoelectrochemical hydrogen production[END_REF]. Due to the broadness of the solar spectrum, from infrared to ultra-violet light, the needs for combining monolithically different materials with different bandgaps and thus absorbing different wavelengths, is considered today as the main pathway to reach high solar-to-hydrogen (STH) conversion efficiency [START_REF] Rongé | Monolithic cells for solar fuels[END_REF][START_REF] Bae | Strategies for stable water splitting via protected photoelectrodes[END_REF]. In the tandem 102 association of two materials, the optimum combination of bandgaps has been widely discussed for different device configurations [START_REF] Hu | An analysis of the optimal band gaps of light absorbers in integrated tandem photoelectrochemical water-splitting systems[END_REF]. Especially, the combination of a 1.7

eV top absorber bandgap with a 1.1 eV bottom absorber is recognized as one of the best tandem materials configuration and gives rise to a theoretical maximum STH efficiency ηSTH, larger than 26% [START_REF] Young | Direct solar-to-hydrogen conversion via inverted metamorphic multijunction semiconductor architectures[END_REF]. Different device demonstrations were proposed with a III-V/III-V materials design, mostly based on the Ga(In)As/GaInP association, reaching ηSTH larger than 15% [START_REF] Young | Direct solar-to-hydrogen conversion via inverted metamorphic multijunction semiconductor architectures[END_REF][START_REF] Khaselev | A Monolithic Photovoltaic-Photoelectrochemical Device for Hydrogen Production via Water Splitting[END_REF][START_REF] May | Efficient direct solar-to-hydrogen conversion by in situ interface transformation of a tandem structure[END_REF][START_REF] Cheng | Monolithic Photoelectrochemical Device for Direct Water Splitting with 19% Efficiency[END_REF]. In these works, high quality materials are achieved, thanks to the lattice-matching of theses alloys on the expensive GaAs substrate. A way to reduce the cost of tandem materials association is to monolithically integrate a III-V top absorber

(1.7 eV bandgap) on the silicon substrate (because of its approximate 1.1 eV band gap, earth abundance, low cost and prevalence in the electronics and PV industries) [START_REF] Hu | An analysis of the optimal band gaps of light absorbers in integrated tandem photoelectrochemical water-splitting systems[END_REF]. But by now there are only very few reports on Si-based tandem systems for water splitting, mainly due to the difficulty in growing high quality III-V epilayers on Si. The recent progresses in the understanding of III-V/Si epitaxial processes and devices developments give new hopes for the development of high efficiency III-V/Si PEC devices on the lowcost Si substrate [3,[START_REF] Chen | Electrically pumped continuous-wave III-V quantum dot lasers on silicon[END_REF]. Following this approach, pioneering works were performed on the development of bipolar configured 1.6 eV AlGaAs on 1.1 eV Si tandem association.

Water splitting at a 18.3% conversion efficiency was reported [START_REF] Licht | Efficient Solar Water Splitting, Exemplified by RuO2-Catalyzed AlGaAs/Si Photoelectrolysis[END_REF]. Very recently, a single InGaN absorber photoanode monolithically integrated on silicon (where the Si [START_REF] Argaman | Density functional theory: An introduction[END_REF] substrate is used as a back contact) was proposed with an applied-bias photo-to-current efficiency of 4.1% [START_REF] Kumar | Quantum dot activated indium gallium nitride on silicon as photoanode for solar hydrogen generation[END_REF]. GaP-based materials were also proposed as an ideal tandem association with Si for water splitting [START_REF] Lucci | A Stress-Free and Textured GaP Template on Silicon for Solar Water Splitting[END_REF]. However, as discussed in Chapter 1, this approach suffers from the indirect and large bandgap (2.26 eV) nature of GaP. Therefore, alloying GaP with other group III-V atoms is needed to achieve a direct bandgap at 1.7

eV through bandgap engineering. Doscher H. et al. theoretically proposed GaPN/Si and GaPNAs/Si tandem lattice-matched materials for water splitting based on GaP/Si-based photoelectrochemistry experiments and GaPN/Si epitaxial growth results [START_REF] Döscher | Epitaxial III-V Films and Surfaces for Photoelectrocatalysis[END_REF]. But the N-including lattice-matched materials face the issue of excitons localization effects, which hamper easy charge carrier transportation in the developed devices. On the other hand, metamorphic III-V integration on Si is known to generate dislocations that are usually detrimental for solar devices. The epitaxial growth of Sb-based III-V compounds on Si is particularly interesting in this regard, as it allows relaxation of the crystal stress through a near-perfect misfit dislocation network localized at the III-V/Si interface [3] leading to efficient and stable photonic device demonstrations [START_REF] Nguyen-Van | Quantum cascade lasers grown on silicon[END_REF]. Recent work proposed the use of freestanding GaPSb as a photoelectrode for water splitting [START_REF] Martinez-Garcia | Unassisted Water Splitting Using a GaSbxP(1-x) Photoanode[END_REF].

However, the crystal growth and properties of this novel alloy was not yet deeply investigated.

In

GaPSb/Si material growth and targeted device design for solar water splitting

Three GaP1-xSbx/Si samples (referred to as GaPSb-1, GaPSb-2 and GaPSb-3, with increasing Sb amounts) were grown by Molecular Beam Epitaxy (MBE) on HFchemically prepared n-doped (around 10 17 cm -3 ) Si (001) substrates, with a 6° miscut toward the [START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] direction. The substrates were heated at 800°C for 10 minutes to remove hydrogen at the surface. 1µm-thick GaP1-xSbx layers were then grown at 500°C in a conventional continuous MBE growth mode, and at a growth rate of 0.24 ML/s, with a Beam Equivalent Pressure V/III ratio of 5. The whole epilayer was undoped, and epitaxial strategies to annihilate antiphase boundaries [START_REF] Quinci | Defects limitation in epitaxial GaP on bistepped Si surface using UHVCVD-MBE growth cluster[END_REF] were not used, leading to emerging APBs.

104

The schematic diagrams of the proposed GaP1-xSbx/Si tandem device for water splitting and the light absorption of the GaP1-xSbx/Si tandem system with around A miscut of around 6° toward the [START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] direction is inferred for the three samples from the positions of Si Bragg peaks, which is in agreement with the substrate specifications.

Structural properties of GaPSb/Si epilayers

The ω/2θ scans exhibit well-defined GaP1-xSbx Bragg peaks for the three samples.

Reciprocal space maps (RSM) carried out on either (004) or (115) (Fig. 4.3) reflections show a full plastic relaxation of the GaP1-xSbx layers for the three samples. GaP1-xSbx 105 lattice parameters were extracted from both reciprocal space maps and ω/2θ scans, leading to very similar values for each sample, confirming the full plastic relaxation rates and giving mean lattice parameters of 0.5665 nm, 0.5835 nm, 0.6093 nm, for the samples GaPSb-1, GaPSb-2, GaPSb-3, respectively. The Sb contents of 0.33, 0.60, and >0.99 are then inferred [START_REF] Nguyen Thanh | Structural and optical analyses of GaP/Si and (GaAsPN/GaPN)/GaP/Si nanolayers for integrated photonics on silicon[END_REF]. Sample GaPSb-3 is almost pure GaSb. These RSM images also exhibit an important Bragg peak broadening due to a relatively large crystal defect density, in low Sb content samples. Furthermore, the XRD analysis does not give any evidence of a phase separation that could occur between GaP and GaSb in the GaP1-xSbx alloys in theses growth conditions. Overall, from these measurements, we can conclude that the GaPSb epilayers grown on the Si substrates are monocrystalline, but defective, and are almost fully relaxed (strain effects won't be included in bandgap energy calculations). We will also consider that, when CMP is used, roughness of the samples cannot impact significantly the measured optical properties.

Bandgap evaluation of GaPSb alloys

Ellipsometry measurements

In order to evaluate the bandgap of the GaPSb/Si samples, ellipsometry measurements were performed. This technique has the advantage to extract the optical constants of the grown layer, independently of the Si substrate. The optical constants of The absorption spectra were employed to plot the Tauc's curve ((αhν) k vs hν) for the four GaPSb samples (Fig. 4.13) based on the Tauc's law:

(αhν) k =C(hν-Eg) (4-1)
Where: k=1/2 for the GaP indirect bandgap and k=2 for the direct bandgap GaP0.67Sb0.33, GaP0.40Sb0.60 and GaSb (this assumption will be further confirmed by the tight binding calculations in the following part), α the absorption coefficient, h the Planck constant, ν the photon frequency and C a constant. An absorption coefficient range between 3000 and 10000 cm -1 was used for bandgap evaluation of the direct bandgap GaP0.67Sb0.33, GaP0.40Sb0.60 and GaSb samples. While for GaP, with its indirect bandgap, the absorption coefficient corresponding to the bandgap is smaller and the absorption coefficient range was chosen between 100 and 5000 cm -1 . The optical band gap can be obtained by linear extrapolation of the straight-line portion to α=0. Finally, the bandgaps of the four samples were estimated: Eg=2.25±0.04 eV for GaP; Eg=1.70±0.06 eV for GaP0.67Sb0.33;

Eg=1.04±0.08 eV for GaP0.40Sb0.60, and Eg=0.68±0.09 eV for GaSb (as shown in Fig.

4.15)

, which are consistent with bandgaps reported in pioneering works for the metamorphic growth of GaPSb on InP substrate [START_REF] Loualiche | A new ternary material for Schottky diode fabrication on InP[END_REF]. In this study, we neglect the possible impact of alloy fluctuations and Sb ordering on the bandgap determination.

Further experiments would be needed to assess quantitatively the influence of such effects.

Tight Binding calculations

The band structures of the unstrained GaP1-xSbx alloys in the whole Sb compositional range have then been calculated by tight-binding calculations, using an extended basis sp3d5s* tight binding Hamiltonian [START_REF] Jancu | Empirical spds* tight-binding calculation for cubic semiconductors: General method and material parameters[END_REF]. This technique has been used for years to predict the bandgap evolution of III-V ternary alloys, as it provides a good accuracy for the bandgap calculations on simple alloys, while keeping a reasonable computing time. From this analysis, we deduce that the Sb content at which a 1.7 eV direct bandgap needed for tandem materials association can be reached, is obtained at 32%.

Band alignment of GaPSb/Si with water redox levels

For efficient photoelectrochemical water splitting, another important aspect is to estimate the band edge location with respect to the redox potentials of water. The band alignment (no external electric field applied) of the GaP1-xSbx/Si tandem architecture for water splitting is presented in Fig. 4.16 as a function of the Sb content. The VBM (valence band maximum) energy (Ev) of the GaP1-xSbx alloy over the whole Sb content range was obtained based on the absolute bandlineup between VBM energy of GaP and GaSb [START_REF] Van De Walle | Universal alignment of hydrogen levels in semiconductors, insulators and solutions[END_REF],

which follows the linear formula Ev(GaP1-xSbx)= xEv(GaSb)+ (1-x)Ev(GaP) [START_REF] Borg | Synthesis and properties of antimonide nanowires[END_REF]. In a first approximation, the evolution of the surface acidity with Sb content was neglected.

The CBM (conduction band minimum) energy of the GaP1-xSbx alloy was obtained by using the results from TB calculations presented in Fig. 4.15. For this alloy, the minimum of the conduction band is located in the X-valley between 0% and 11%. It then moves to the L valley for Sb contents between 11% and 30%. It finally reaches a direct bandgap configuration (minimum of the CB in the Γ valley) beyond 30% of Sb incorporation. While the precise value predicted by the calculations for the L to Γ valley crossover seems hard to confirm experimentally in this work, calculations clearly demonstrate that in any cases, an indirect to direct cross-over is expected for this alloy at around 30% Sb. The VBM energy of bulk GaP is very close to the oxidation potential of water, which has been verified by many reports [START_REF] Van De Walle | Universal alignment of hydrogen levels in semiconductors, insulators and solutions[END_REF][START_REF] Jing | Efficient solar hydrogen production by photocatalytic water splitting: From fundamental study to pilot demonstration[END_REF][START_REF] Kudo | Heterogeneous photocatalyst materials for water splitting[END_REF][START_REF] Chu | Roadmap on solar water splitting: current status and future prospects[END_REF].

So, here we position the redox levels of water from the VBM position of GaP [START_REF] Van De Walle | Universal alignment of hydrogen levels in semiconductors, insulators and solutions[END_REF] and the energy differences between the water redox levels and GaP VBM reported in previous works [START_REF] Van De Walle | Universal alignment of hydrogen levels in semiconductors, insulators and solutions[END_REF][START_REF] Jing | Efficient solar hydrogen production by photocatalytic water splitting: From fundamental study to pilot demonstration[END_REF][START_REF] Kudo | Heterogeneous photocatalyst materials for water splitting[END_REF][START_REF] Chu | Roadmap on solar water splitting: current status and future prospects[END_REF], with error bars to analyze potential water splitting reactions more accurately and comprehensively (Fig. 4.16).

Discussion and assessment of GaPSb/Si for water splitting

From the Fig. 4. [START_REF] Kour | Review-Recent Advances and Challenges in Indium Gallium Nitride (InxGa1-xN) Materials for Solid State Lighting[END_REF] and Fig. 4.16, it can be observed first that, in order to benefit from a direct GaPSb bandgap larger than the one of the silicon (to absorb higher energy solar radiations), the Sb content should lie between 30% (indirect to direct cross-over) and 54%

116

(bandgap equal to the Si one). In this Sb content range, it is also noticed that the band lineup is of type I, promoting the charge carrier extraction in the silicon under zero bias conditions. Now looking at the band lineups between GaPSb and the water redox potentials in the 30%-54% Sb content range, it is seen that GaPSb has a strong reduction ability due to its relatively higher CBM (around 0.9 eV higher than H + / H2 potential at 32% Sb content). On the other hand, the VBM is higher than the O2/H2O potential in the 30%-54% Sb content range, and due to the linear increase of VBM with the Sb content, the GaPSb with low Sb content (around 30%) is thus more suitable for oxygen evolution reactions (about 0.43 and 0.46 eV higher than O2/H2O potential at 30% and 32% Sb content). Overall, addition of Sb in GaP will therefore ease photocathode operation, although large currents densities can be obtained with a GaPSb/Si photoanode (this will be shown in the following part and the explanation for this given in Chapter 5) due to a strong direct bandgap absorption. Of course, the picture given here is a simplified one, as many other parameters should be taken into account, and especially the physics & chemistry of the Helmoltz layer. However, the band lineups calculated in this work give a basis for further devices analysis, where the precise determination of Nernstian shifts and band bending will be needed in future works.

Photo-Electro-Chemical characterizations

In order to further check the potential of GaPSb/Si tandem materials for solar water splitting applications, photoelectrochemical performances were investigated at UCL (London) for the GaP0.67Sb0.33/Si sample which is very close to the optimal bandgap combination.

117 Figure 4.17 The experimental setup using a three-electrode system for the photoelectrochemical measurements. The PEC properties of GaP0.67Sb0.33/Si photoanodes were investigated using a standard three-electrode configuration, consisting of a working electrode (GaP0.67Sb0.33/Si), a reference electrode (Ag/AgCl), and a counter electrode (platinum coil) in 1.0 M KOH electrolyte (pH=14) under one sun illumination, as shown in Fig. 4.17.

The voltammetry curve for the GaP0.67Sb0.33/Si photoanode without a protection layer is shown in Fig. 4.18. Under AM 1.5G simulated one sun illumination, it shows an onset potential (where the visible increase of current density starts in the voltammetry curve) for the GaP0.67Sb0.33/Si photoanode of around 0.4 V versus RHE. The saturated photocurrent density is 13.26 mA/cm 2 at 1.23 V vs RHE. For comparison, PEC measurements of GaP reference photoanodes were also taken and the corresponding voltammetry curve is shown as the green curve in Fig. 4.18. We can find that, as compared to the GaP photoelectrode, the GaPSb/Si photoelectrode shows much higher photocurrent which should, in a first approach, result from the better light absorption. Surprisingly, we note that a relatively high photo-current is observed, despite the absence of any intentional doping, and despite the numerous crystal defects that were evidenced previously (the full relaxation of the crystal implies a large density of dislocations generated during the metamorphic growth). This is very unusual, as conventional III-V photoelectrodes usually require n-or p-doping to ensure efficient transfer and transport of charge carriers to the electrolyte. We will demonstrate in Chapter 5 that emerging APBs strongly impact the operation of these photoelectrodes. STEM measurements performed on the GaP0.67Sb0.33/Si photoelectrode, provides evidence that APBs were not buried in this sample, and that they emerge at the sample surface (Fig. The large photo-current obtained with the III-V/Si photoelectrode and presented in Fig. 4.18 should however be considered carefully. Indeed, the voltammetry curves presented also show a significant current under dark conditions, which confirms that corrosion occurs significantly during the photoelectrode operation. In addition, one can expect that photo-corrosion also occurs during the photoelectrode operation. This is also confirmed by the very short life-time of the developed bare GaPSb/Si photo-electrode.

Indeed, under the considered pH=14 conditions, the bare photoelectrodes were able to work during less than 10 J(V) scans. The incident photon-to-current conversion efficiency (IPCE) as a function of the wavelength of the incident light for the GaP0.67Sb0.33/Si without coating, and reference GaP photoanodes were further investigated at an applied bias of 1.23 V versus RHE in a 1.0 M KOH electrolyte (pH=14). The IPCE of the GaP0.67Sb0.33/Si photoanode coated with Ni (8 nm) is equal or larger than 122 [START_REF] Fitzgerald | Totally relaxed GexSi1-x layers with low threading dislocation densities grown on Si substrates[END_REF].16% over the visible light region from 400 nm to 550 nm, due to the catalytic behavior of Ni for oxygen evolution reaction [START_REF] Standing | Efficient water reduction with gallium phosphide nanowires[END_REF][START_REF] Strandwitz | Photoelectrochemical Behavior of Planar and Microwire-Array Si|GaP Electrodes[END_REF][START_REF] Kargar | GaP/GaNP Heterojunctions for Efficient Solar-Driven Water Oxidation[END_REF]. In contrast, the GaP0.67Sb0.33/Si photoanode coated with TiO2 (20 nm) and Ni (8 nm) reaches a high IPCE of 76.4% at 400 nm and at least 67.1% between 400 nm and 650 nm, which indicates efficient and fast charge transfer to the semiconductor/electrolyte interface (Fig. 4.23). The IPCE falls and tends to zero around 750 nm in agreement with the energy bandgap (1.70±0.06eV) deduced previously from the measured GaP0.67Sb0.33 absorption spectrum. From IPCE measurements, we note here that the Si substrate does not contribute to the PEC response.

This is an important observation that will be widely discussed in Chapter 5, where we will show how one can promote photo-response of the III-V or of the Si material. The stability of III-V semiconductors is a critical challenge in aqueous solution [START_REF] Hu | Amorphous TiO2 coatings stabilize Si, GaAs, and GaP photoanodes for efficient water oxidation[END_REF].

To assess the stability of the GaP0.67Sb0.33/Si photoanodes, the photocurrent density versus the time (J-t) characteristics of electrodes were investigated at a constant potential of 1 V versus RHE in a corrosive solution of 1.0 M KOH (pH=14) as shown in Fig. 4. [START_REF] Ng | An efficient room-temperature silicon-based light-emitting diode[END_REF] and While these PEC results are very promising (both in terms of performances and stability), numbers of experimental observations cannot be explained with a simple bandgap engineering and band lineups approach. As discussed above, the GaPSb layer was not intentionally doped, and thus neither good carriers transport properties nor high photocurrents would normally be expected with such sample. Therefore, a deep understanding of the III-V/Si photoelectrode operating principle will be needed. It will be proposed in Chapter 5, which will give new and general insights of the involved photoelectro-chemical mechanisms.

Summary

In this chapter, we studied the novel GaP1-xSbx/Si tandem materials association for solar water splitting. The bandgaps and band structures of GaPSb alloys were determined over the whole Sb range, by combining experimental absorption measurements with tight binding theoretical calculations. The indirect (X-band) to direct (Γ-band) cross-over was found to occur at 30% Sb content. Especially, at a Sb content of 32%, the GaP1-xSbx alloy reached the desired 1.7 eV direct bandgap, enabling efficient sunlight absorption.

Moreover, the band alignment of GaP1-xSbx alloys and Si with respect to water redox potential levels have been analyzed, which shows the 1.7eV/1.1eV GaPSb/Si association is an interesting combination both for the hydrogen evolution and oxygen evolution reactions. At last, a GaP0.67Sb0.33/Si photoanode with a bandgap configuration very close to the optimal bandgap combination (1.7eV/1.1eV) was investigated for PEC characterizations and promising PEC performances were obtained due to the reduced direct bandgap of GaP0.67Sb0.33. Besides, it was proved that the stability of the GaP0.67Sb0.33/Si photoanode can be much improved with TiO2 and Ni coating layers.

These results further demonstrate the promises of the GaP1-xSbx/Si tandem system for low-cost high-efficiency solar water splitting. But it also raises many opened questions.

For instance, how can we get a high current in these PEC devices, despite the high density 125 of defects, and without the need of any intentional doping? Could it become possible to tune at will the properties of the photoelectrodes by using bandgap engineering of III-V semiconductors? Could the Si become photoactive in these photoelectrodes? Chapter 5

will try to answer some of the numerous questions arising from results shown in this chapter.

Introduction

Hybrid materials and devices have driven many researches recently, addressing the most advanced issues in today's energy harvesting [START_REF] Liang | Co3O4 nanocrystals on graphene as a synergistic catalyst for oxygen reduction reaction[END_REF], photonic [START_REF] Li | Assembling a lasing hybrid material with supramolecular polymers and nanocrystals[END_REF], sensing [START_REF] Zhou | Seebeck-driven transverse thermoelectric generation[END_REF] or computing [START_REF] Allain | Electrical control of the superconducting-to-insulating transition in graphene-metal hybrids[END_REF] technologies. Among the various challenges raised, the simultaneous and ultimate control of carrier transport and optical properties in the same material remains usually tricky. On the one hand, the bandgap of semiconducting materials provides a path for electron-hole generation through light absorption, or electron-hole 128 recombination through light emission, especially with direct bandgap semiconductors, such as III-V, II-VI, CuInGaSe or perovskites [START_REF] Yu | Fundamentals of Semiconductors: Physics and Materials Properties[END_REF]. The price to pay is the necessity to adjust carefully the internal electric field and the potential profile of stacked materials needed for carrier injection or extraction, thanks to the so-called p-type/n-type architecture or using electron and hole transporting layers. Only organic devices successfully manage to combine in the same material both photo-generation/light emission and carriers extraction or injection using bulk heterojunctions, but with limited yields, and lifetime issues [START_REF] Vandewal | On the origin of the open-circuit voltage of polymer-fullerene solar cells[END_REF]. On the other hand, the absence of a bandgap in metals, semimetals or Dirac materials make them highly interesting for transport, but can hardly be employed as photonic or energy harvesting materials [START_REF] Li | Graphene-On-Silicon Schottky Junction Solar Cells[END_REF].

Most of the classical semiconductor devices developed in the previous century were based on the combination of single usage unipolar layers, where the predominant conductive charges are holes for p-type and electrons for n-type, in dedicated architectures and applications. In the last past years, different researches pointed out the prospects offered by ambipolar devices, where simultaneous transport of electrons and holes can be achieved, thus promoting simultaneously n-and p-type characteristics. This led to recent ground-breaking demonstrations in the field of memory or logic devices [START_REF] Wang | The ambipolar transport behavior of WSe2 transistors and its analogue circuits[END_REF][START_REF] Zhou | Nonvolatile multilevel data storage memory device from controlled ambipolar charge trapping mechanism[END_REF], neuromorphic computing [START_REF] Chen | Polyoxometalates-Modulated Reduced Graphene Oxide Flash Memory with Ambipolar Trapping as Bidirectional Artificial Synapse[END_REF], sensing [START_REF] Kwon | Gate-tunable gas sensing behaviors in air-stable ambipolar organic thinfilm transistors[END_REF], light emitting devices [START_REF] Zaumseil | Spatial control of the recombination zone in an ambipolar light-emitting organic transistor[END_REF][START_REF] An | Well-Balanced Ambipolar Organic Single Crystals toward Highly Efficient Light-Emitting Devices[END_REF], photovoltaics [START_REF] Zhang | Synergistic Cascade Carrier Extraction via Dual Interfacial Positioning of Ambipolar Black Phosphorene for High-Efficiency Perovskite Solar Cells[END_REF] or photo-electrochemical cells [START_REF] Bourée | Robust Hierarchically Structured Biphasic Ambipolar Oxide Photoelectrodes for Light-Driven Chemical Regulation and Switchable Logic Applications[END_REF].

In this chapter, we demonstrate experimentally and theoretically that epitaxial III-

Sample growth

In this chapter, the III-V samples (GaP, GaPSb and GaAsP are considered in this study) were grown by Molecular Beam Epitaxy (MBE) on HF-chemically prepared ndoped (5x10 14-18 cm -3 ), p-doped (2.6x10 15 -1.4x10 16 cm -3 ) or high resistivity (>5000 Ω.cm) Si (001) substrates. As described in Chapter 4, the substrates were heated at 800°C for 10 minutes to remove hydrogen at the surface. 1µm-thick III-V layers were then grown at 500°C in a conventional continuous MBE growth mode, and at a growth rate of 0.24 ML/s, with a Beam Equivalent Pressure V/III ratio of 5. The whole epilayer was undoped, and epitaxial strategies to annihilate antiphase boundaries [START_REF] Quinci | Defects limitation in epitaxial GaP on bistepped Si surface using UHVCVD-MBE growth cluster[END_REF] were not used here, leading to emerging APBs.

Structural analysis of III-V/Si materials with emerging APBs

As The spatial distribution of the singularities is of great importance, as it directly defines their nature. Indeed, considering a perfect vertical propagation of APBs along the

[001] direction, APBs lying along the [START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] or [-110] directions will be perfectly stoichiometric, while APBs lying along [START_REF] Garcia-Caurel | Application of Spectroscopic Ellipsometry and Mueller Ellipsometry to Optical Characterization[END_REF] or [010] directions will be perfectly nonstoichiometric (i.e. with only III-III or V-V bonds). In all the other cases, this will lead to intermediate charge states. A schematic diagram of in-plane charge distribution is given in Fig. 5.5, where stoichiometric APBs lying along [START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] or [-110] correspond to the neutral configuration σ° in blue, and non-stoichiometric APBs lying along [START_REF] Garcia-Caurel | Application of Spectroscopic Ellipsometry and Mueller Ellipsometry to Optical Characterization[END_REF] or [010] correspond to the perfectly charged configuration σ +/-in red. The positive or negative nature of the charge cannot be determined here, but will be discussed in details later in this work. resolved observations performed in previous works [START_REF] Lucci | A Stress-Free and Textured GaP Template on Silicon for Solar Water Splitting[END_REF] and is not fully surprising as the APBs naturally result from the coalescence of monodomain islands [3]. The simplified picture given here does not take into account the possible charge compensation effects which could occur during the growth [START_REF] Beyer | Atomic structure of (110) anti-phase boundaries in GaP on Si(001)[END_REF], but nevertheless, it clearly evidences that the III-V/Si samples are not composed only of σ°-like APBs, but also of σ +/--like ones. This excess of charges is expected to impact significantly transport properties, which are studied in the following.

Transport characterizations of III-V/Si with emerging APBs

Hall measurements

To probe the electrical conductivity and charge carrier density, and identify the specific contribution of APBs to transport properties, 1 µm-thick GaP samples were grown on different Si substrates (6°-off or nominal ones, n-doped or with a high resistivity 134 (HR)) with either buried APDs [START_REF] Guillemé | Antiphase domain tailoring for combination of modal and 4¯ -quasiphase matching in gallium phosphide microdisks[END_REF] or emerging APBs (such as in Fig. 5.2). Importantly, all these III-V layers were not intentionally doped. For some measurements, the silicon substrate was etched, and standalone GaP membranes were bonded onto a thermally oxidized silicon host wafer with an insulating adhesive polymer (benzocyclobutene) layer, to better separate the transport properties of the III-V layer from the Si substrates. Hall effect measurements were carried out on these samples and the results are summarized in Table 5.1. For the GaP membranes with emerging or buried APBs on insulating layer, transport physical parameters could not be determined, because of their very large electrical resistance exceeding experimental capabilities. Note that a high resistance was also previously measured for a GaP layer with buried APBs on a n-doped Si substrate [START_REF] Cornet | Chapter 28 -GaP/Si-Based Photovoltaic Devices Grown by Molecular Beam Epitaxy[END_REF]. This result would be typically expected for the large bandgap non-intentionally doped GaP bulk material. However, charge carriers are easily injected into GaP layers with emerging APBs on Si substrates as demonstrated by Hall effect measurements. For both GaP/Si-n and GaP/Si-HR samples with emerging APBs, values of the measured mobility are about one order of magnitude higher than those reported in the literature for GaP/Si [START_REF] Wright | Molecular beam epitaxial growth of GaP on Si[END_REF] or even for bulk GaP [START_REF] Kao | Electron and hole carrier mobilities for liquid phase epitaxially grown GaP in the temperature range 200-550[END_REF][START_REF] Alberding | Direct comparison of time-resolved terahertz spectroscopy and Hall Van der Pauw methods for measurement of carrier conductivity and mobility in bulk semiconductors[END_REF]. More precisely, we observed that the measured mobility is very close to that expected for Si. This result indicates that 135 most of the in-plane current flows within the underlying Si substrate. For the Si substrate, the measured sheet resistance of 170 Ω/sq gives a resistivity of 5.95 Ω.cm and the carrier sheet density of 3.5×10 13 cm -2 leads to a bulk carrier density of 1×10 15 cm -3 . These values of resistivity and mobility are exactly those expected for Si with a doping concentration giving an electron density equal to the measured one [START_REF] Jacoboni | A review of some charge transport properties of silicon[END_REF]. All these results combined together unambiguously demonstrate that carriers transport in bi-domain III-V semiconductors occurs mostly vertically, and that the current path is related to APBs. This current path is due to a great contrast between the in-plane APB resistivity and the outof-plane APB resistivity. From our measurements, we estimate that the former is at least 10 3 times larger than the latter.

Conductive atomic force microscopy (C-AFM) measurements

To get more local information about transport, the nano-electronic properties of the

Linear sweep voltammetry (LSV)

The chopped-light linear sweep voltammetry (LSV) curve of the bidomain GaP/Si:n 138 sample is displayed in Fig. 5.9a. Interestingly, it shows both anodic and cathodic photocurrents. This unusual ambipolar behavior is of great interest for the development of fully integrated PEC cells, as it may allow the use of a single material to be used for both photoelectrodes. The measured anodic photocurrent is higher than the cathodic one, in good agreement with the C-AFM results presented above. For comparison, the same curve is given for a single GaP:n wafer (Fig. 5.9b), which shows, as expected, only a photoanodic behavior. and the latter reactions [START_REF] Butler | P-Type GaP as a Semiconducting Photoelectrode[END_REF][START_REF] Malizia | Formation of a p-n heterojunction on GaP photocathodes for H2 production providing an open-circuit voltage of 710 mV[END_REF][START_REF] Pishgar | Investigation of the photocorrosion of n-GaP photoanodes in acid with in situ UV-Vis spectroscopy[END_REF]. Despite the corrosion process, the reached maximum photocurrent is relatively high, and not far from that measured with bare III-V substrates, indicating that the photogeneration and transport are overall efficient in these devices. For stable water photoelectrolysis, the corrosion issue is expected to be tackled effectively in a short-term perspective by using catalytic and/or protection layers, as widely demonstrated on numerous unstable semiconductors [START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF][START_REF] Hu | Amorphous TiO2 coatings stabilize Si, GaAs, and GaP photoanodes for efficient water oxidation[END_REF][START_REF] Alqahtani | Gallium Phosphide photoanode coated with TiO2 and CoOx for stable photoelectrochemical water oxidation[END_REF].

Incident photon-to-current efficiency (IPCE) spectra

Figure 5.12 shows the incident photon-to-current efficiency (IPCE) spectra of the GaP/Si:n sample and GaP:n wafer, recorded at 1.26 V (anodic region) and -0.74 V (cathodic region) versus RHE. Firstly, the integrated photocurrents derived from the IPCE curve (0.2 mA.cm -2 for GaP/Si at -0.74 V; 1.0 mA.cm -2 for GaP/Si at 1.26 V; 0 mA.cm -2 for the GaP wafer at -0.74 V; 1.68 mA.cm -2 for the GaP wafer at 1.26 V) show good consistencies with those extracted from the voltammetry curves (Fig. 5.9), which validate the accuracy of the measurements. Secondly, the GaP/Si sample shows IPCE signal on both anodic and cathodic regions, which further confirms the ambipolar characteristic of the GaP/Si sample. Thirdly, the IPCE of GaP/Si mainly increases around 550nm, similarly to the GaP wafer, which corresponds well to the bandgap of the GaP semiconductor, indicating the carriers are mainly created by the absorption within the GaP matrix. This is an important point as it already confirms that carriers are photogenerated in bulk-like domains, and then transferred to APBs which ensure the transport, and that the absorption of the underlying Si substrate is not interfering in the PEC response. It suggests that the charge-transfer resistance (Rct) (correlated with the diameters of the semi-circles measured) [START_REF] Randviir | Electrochemical impedance spectroscopy: an overview of bioanalytical applications[END_REF] strongly decreased under illumination and that the relatively small semi-circles were obtained in both anodic and cathodic regimes, echoing with the ambipolar behavior. From all these PEC characterizations, we can therefore draw the following picture:

the III-V/Si materials with emerging APBs allow (i) photogeneration of carriers in bulk monodomain semiconducting grains, (ii) carrier separation and extraction from the bulk phase to the APB singularities, (iii) ambipolar vertical transport and (iv) efficient charge transfer at the solid/liquid interface.

DFT calculations for non-stoichiometric APB structures

In order to better understand the original electronic and PEC properties of the APB singularities buried into III-V semiconductors, theoretical calculations are performed for III-V-based non-stoichiometric APB structures.

Construction of non-stoichiometric APB supercells

As

Band structure and charge density distribution of nonstoichiometric APBs

Electronic band structure calculations were performed using Density Functional Theory (DFT) along X(0.5,0,0)-(0,0,0)-X(-0.5,0,0) k-path, where the axes x, y, z of the atoms) with a zinc-blende lattice (Fig. 5.17), the slab structures used to mimic bulk properties, exhibit slightly larger bandgaps due to quantum effect, and exhibit more bands resulting from band folding effect. The Fermi energy levels are determined by DFT calculations using dense k point meshes and a metaGGA (TB-mBJ) functional [START_REF] Tran | Accurate Band Gaps of Semiconductors and Insulators with a Semilocal Exchange-Correlation Potential[END_REF],

which possesses similar accuracy than computationally more expensive functionals, such as HSE, or even many-body calculations at the GW level [START_REF] Becke | Exchange holes in inhomogeneous systems: A coordinatespace model[END_REF] (see the next section).

Compared to the reference III-V slab structures, all the III-III and V-V APB structures introduce localized metallic states on the APB atoms in the gap, which indicates that they are a direct consequence of the specific 2D topology of the APB.

3D band structure and density of states

To further explore the intrinsic semimetal nature of these APB structures, firstprinciple tight-binding models were constructed by projecting the electronic states on to

Wannier orbitals [START_REF] Marzari | Maximally localized Wannier functions: Theory and applications[END_REF]. This analysis was performed on top of state of the art DFT 148

The total density of states (unit: states/eV) of the slab structures were extracted, as shown in Fig. 5.21, which corresponds essentially to the superposition of a 2D electronic density of states with 3D-like CB (Conduction Band) and VB (Valence Band) densities of states. In order to better capture the effect of this superposition in real heterostructures, it is necessary to consider that the experimental linear density of APBs is in fact smaller than the effective ones considered for the simulation due to the limited size of the supercells.

In the realistic material, the linear density of APB is 10. to get the real total density states. Finally, the total density of states is divided by the volume and we can get the real density of states with a states/eV/m 3 unit for the III-III APB structure in realistic materials, as shown in Fig. 5.22b. In the same way, the density states for the other III-III and V-V APB structures can be obtained, as shown in Fig. 5.22.

The DOS representation confirms the semimetallic nature of both III-III and V-V APBs and a large ratio of at least 10 3 between the number of states available at a given energy in the bulk region and the number of states localized in the APBs indicates that any charge transfer from the bulk semiconductor domains to the 2D APB singularities drastically changes the filling of the APBs states and thus easily shifts its Fermi level. Moreover, 3D plot of the VBM (Valence Band Maximum) and CBM (Conduction Band Minimum) for III-III and V-V APB structures corresponding to kz=0 planes for GaP, GaAs and GaSb are plotted in Fig. 5.22a,c,e,g,i,k. The crossing points are marked by black dots. We can observe that for the III-III APB structures, one pair of Weyl points [START_REF] Yan | Topological Materials: Weyl Semimetals[END_REF] and one closed nodal ring [START_REF] Yang | Symmetry demanded topological nodal-line materials[END_REF] around the high symmetry Y point exist. The pair of Weyl points is always observed for all the III-V materials with APBs. While the nodal ring varies from a large one (GaP) to small one (GaAs) and even disappears for GaSb, which may be related to the differences of their bandgaps. On the other hand, non-closed snakelike nodal lines [START_REF] Yang | Symmetry demanded topological nodal-line materials[END_REF] are evidenced for the V-V APB structures. The crossing points are further confirmed by DFT enlarged band structure calculations based on metaGGA (TB-mBJ) potential combined with band symmetry analysis (the detailed analysis will be given in the following part). The pairs of topological Weyl points observed for III-III APB structures are characterized by a Berry phase equals to , along a close k path in kz=0 plane (as shown in Fig. 5.23). The band crossings are protected by a combination of the time-reversal symmetry preserved in the III-V lattice and the spatial inversion symmetry restored at the APB [START_REF] Yang | Symmetry demanded topological nodal-line materials[END_REF][START_REF] Huang | Topological nodal-line semimetals in alkaline-earth stannides, germanides, and silicides[END_REF][START_REF] Weng | Topological node-line semimetal in three-dimensional graphene networks[END_REF]. Therefore, such topological features are predicted to be fingerprints of the intrinsic symmetry properties of charged APB planes [START_REF] Sekine | Tunable charged domain wall from topological confinement in nodal-line semimetals[END_REF]. Deeper investigations are however needed to further explore the experimental implications of such topological predictions, but non-stoichiometric and charged III-V APBs presented here appear as leading to a first practical device realization based on these concepts.

Enlarged band structure and Fermi energy calculations

As demonstrated above, the 3D bandstructures of nonstoichiometric APB structures determined from first-principle tight-binding model (Fig.

5.33d,e).

We note that this situation is very different from the one encountered with conventional inorganic semiconductors stacked structures, where the separation of photogenerated carriers is usually made vertically, with the help of the so-called p-n architecture. Under light illumination, the electrons and holes photogenerated in the bulk III-V matrix will then migrate to V-V and III-III APB singularities, respectively, thanks to the in-plane built-in electric field (as shown in Fig. 5.33d,e). The built-in electric field provides here the driving force for efficient photogenerated charge carriers separation. In addition, photogenerated carriers are rapidly captured by APBs due to very short distance (<100 nm), reducing the recombination-related losses, which shed light on the strong decrease of the charge-transfer resistance after illumination (Fig. 5.13) [START_REF] Sacco | Electrochemical impedance spectroscopy as a tool to investigate the electroreduction of carbon dioxide: A short review[END_REF][START_REF] Shaislamov | Growth of CuO/ZnO Nanobranched Photoelectrode with Enhanced Stability for Solar Hydrogen Generation[END_REF]. With the electrons (holes) reinjected in V-V (III-III) APBs, the Fermi energy level of V-V (III-III) increases (decreases) again, as shown in Fig. 5.33a. The induced photo-potential between these two quasi-Fermi levels can be large due to the small DOS of metallic states.

Finally, under negative bias applied vertically, the electrons in V-V APB will be driven to the surface to generate H2 and holes in III-III will be transferred to Si and further feed the counter-electrode (Fig. 5.33d). Under positive bias applied vertically, the holes in III-III will move to the surface to generate O2 and the electrons in V-V will transfer to Si and feed the counter-electrode (Fig. 5.33e), through conductive APB paths. This process also implies that the electrical resistance of the APBs is large enough to be able to apply a potential drop vertically along the APB. From the energy band point of view, under positive and negative bias, the Fermi energy levels of III-V APBs and bulk will shift down (Fig. 5.33b) and up (Fig. 5.33c), respectively, promoting the oxygen evolution (Fig. 5.33b) and hydrogen evolution (Fig. 5.33c) reactions respectively. Overall, the proposed mechanism allows us to explain why the bidomain III-V/Si material allows achieving efficient ambipolar water splitting.

Influence of Si substrate doping on PEC performances

Besides, the III-V samples grown on a n-doped Si substrate show an anodic photocurrent higher than the cathodic one (Fig. Besides, in this case, the PEC response of the Si substrate can be activated. On the other hand, if the Fermi energy of Si is below or above the Fermi levels of both III-III and V-V APB structures, then the same band bendings between Si and V-V APB or Si and III-III APB will occur. For example, this configuration is achieved for GaSb or GaSb-based alloys grown on Si:p substrates (the Fermi level can be lower than both III-III and V-V APB structures). In this situation, the band of Si will bend downward with both V-V and III-III APB structures (Fig. V/Si samples [3,[START_REF] Guillemé | Antiphase domain tailoring for combination of modal and 4¯ -quasiphase matching in gallium phosphide microdisks[END_REF], is promising for the development of novel multifunctional photoelectric devices using epitaxial III-V/Si materials with non-stoichiometric APBs.

Summary

In this chapter, we demonstrated experimentally and theoretically that epitaxial bidomain III-V/Si materials with non-stoichiometric APBs are hybrid structures, composed of bulk semiconductors enabling photo-activity, and antiphase boundaries that are 2D topological semi metallic inclusions with ambipolar transport properties. The demonstration of operating ambipolar III-V/Si photoelectrodes is completed by dedicated theoretical and experimental studies. It shows that the epitaxial bi-domain III-V/Si materials are indeed able within the same layer to absorb light efficiently, separate laterally the photo-generated carriers, transfer them to semimetal singularities, and extract both electrons and holes vertically. These very specific and multifunctional properties obtained within the same layer provide new paradigms for energy harvesting, photonics, electronics or computing applications and may also expand a new route for the exploration of novel topological materials.

Conclusions and perspectives

In Chapter 1, the advantages of III-V/Si co-integration, as well as the general interest of III-V/Si materials for integrated photonics and energy harvesting applications have been clarified. Then the III-V/Si monolithic integration methods and corresponding issues and strategies have been presented. Finally, focusing on antiphase boundaries, their generation, propagation, burying mechanisms and properties were reviewed.

In For III-V samples grown directly on Si, freshly prepared Si substrates are loaded into the MBE chamber, heated at 800°C during several minutes for dehydrogenation and cooled down to III-V layer growth temperature. For samples with Si buffer layer prior to III-V layer growth, the substrates after chemical preparation are introduced to the UHVCVD chamber and heated to 800°C (heater on the sample holder) under a residual atmosphere of silane at about 10 -6 torr. The homoepitaxial Si layer is deposited using silane (SiH4) injected through a mass flow/Baratron flux/pressure controller and the pressure is fixed at 6.10 -3 Torr, with 10 min annealing before cooling down. The homoepitaxial Si/Si sample is then transferred under UHV to the MBE growth chamber for III-V layer overgrowth.

Group-III elements (Ga, In, Al) are generated from conventional Knudsen effusion cells containing pure metal solid sources. The group-V elements (P, As, Sb) are provided via a thermal effusion cell where heated solid sources of these elements (released as tetramer such as P4 and As4) are cracked (by the valved-cracker cell) in P2 and As2 molecules.

There are two growth modes (as shown in Fig. A2). One is the traditional Molecular Beam Epitaxy (MBE) growth where the Group-III elements and the Group-V elements are supplied at the same time and the growth rate is controlled by the flux of Group-III elements. The other mode is called Migration Enhanced Epitaxy (MEE) including alternated deposition of Group-III elements and Group-V elements. The growth rate expressed by monolayer (ML) per second and the amount of group-III elements per cycle are two important parameters for the MEE growth mode. For III-V/Si samples, they can be grown by two-step growth process. It contains low-temperature MEE growth, which allows having a very smooth surface at the early stage of growth, then followed by the conventional molecular beam epitaxy (MBE) growth mode, as shown in Fig. A2. If you want to get more information about the sample growth of III-V/Si, please see the paper of Quinci et al. [START_REF] Quinci | Defects limitation in epitaxial GaP on bistepped Si surface using UHVCVD-MBE growth cluster[END_REF], where all the technical details are given.

Figure. A2

Figure 1

 1 Figure 1 Configuration atomique de l'APB stoechiométrique (a) ou non stoechiométrique (b).

Figure 2 a

 2 Figure 2 a, Image STEM en champ clair (BF) en coupe de la plate-forme InGaP/SiGe/Si. La ligne noire au milieu et la bosse sur la surface correspondent respectivement à l'APB et à la nanostructure de surface. b-e, Images de cartographie élémentaire EDX en coupe de l'indium (b), du gallium (c), du phosphore (d) et du silicium (e). f-h, cartographie élémentaire EDX en vue

Figure 3

 3 Figure 3 (a) Structure de bande électronique calculée d'une structure APB stoechiométrique, pondérée par la localisation spatiale des fonctions d'onde au voisinage de l'APB. (b) Structure de bande vibrationnelle (phonons) de la structure APB stoechiométrique. Les modes de phonon mis en évidence par des lignes pointillées orange, bleues, magenta, vertes et violettes correspondent à des modes localisés supplémentaires liés aux plans 2D d'APB.

Figure 4

 4 Figure 4 Données expérimentales (points gris) et ajustement théorique (ligne rouge) du large pic de PL basé sur le modèle de Franck-Condon. Les petits pics multicolores correspondent aux contributions distinctes des répliques de phonons.

Figure 5

 5 Figure 5 Diagramme schématique de la structure 2D verticale riche en indium, intégrant des plans APB avec une forte interaction électron-phonon.

Figure 6

 6 Figure 6 (a) Schéma du dispositif proposé en tandem GaP1-xSbx/Si pour le craquage PEC de l'eau. (b) Illustration montrant l'absorption de la lumière du soleil par le système tandem GaP1-xSbx/Si avec une combinaison de bande interdite d'environ 1.7/1.1eV pour le craquage de l'eau, qui permet une bonne absorption du spectre solaire.

Figure 7

 7 Figure 7 (a) Bandes interdites d'alliages GaP1-xSbx pour différentes compositions en Sb à température ambiante. Les lignes pleines vertes, rouges et violettes sont des courbes théoriques calculées par des calculs de liaisons fortes, correspondant respectivement aux vallées Γ, L et X. Les barres d'erreur séparées correspondent aux bandes interdites déterminées expérimentalement. (b) Alignement de bande des vallées Γ, L et X de GaP1-xSbx et bande interdite indirecte de Si par rapport aux potentiels d'oxydation et de réduction de l'eau.

Figure 8 Figure 9 Figure 9 Figure 10 Figure 11

 8991011 Figure 8 (a) Courbes de voltamétrie pour des photoanodes de GaP0.67Sb0.33/Si recouverte de Ni 8nm et de GaP0.67Sb0.33/Si recouverte par du TiO2 (20 nm) et du Ni (8 nm) dans un électrolyte KOH 1,0 M sous illumination AM1.5 simulée en fonction de la tension appliquée par rapport à l'électrode hydrogène de référence. (b) Courbe de voltamétrie (scan numéro 6) pour le GaP0.67Sb0.33/Si recouvert de Ni 8 nm ou recouvert de TiO2 20 nm et Ni 8 nm. La courbe de voltamétrie rouge correspond au sixième balayage de la photoanode GaP0.67Sb0.33/Si sans couche de protection.

Figure 12

 12 Figure 12 Diagramme schématique spatial du mécanisme de craquage de l'eau ambipolaire en tant que photocathode (a) et en tant que photoanode (b) à partir des singularités APB non stoechiométriques V-V et III-III. Les rectangles jaune et vert représentent respectivement les singularités III-III et V-V APB, et les lignes noires correspondent aux minima (maxima) de la bande de conduction (valence) du semi-conducteur III-V massif.

Figure 13

 13 Figure[START_REF]Band-gap engineering[END_REF] Illustration générale du matériau bi-domaine III-V/Si composé de zones de semiconducteurs photoactifs avec des inclusions semi-métalliques APB 2D, qui permettent un bon transport et une séparation de charge efficace, réalisant le craquage PEC ambipolaire de l'eau.

Figure 1 . 1

 11 Figure 1.1 Crystal structure of silicon. Courtesy of [7]

Figure 1 . 2

 12 Figure 1.2 Energy band diagram and main carrier transition processes in bulk InP and bulk Si.Courtesy of[8]

Figure 1 . 3

 13 Figure 1.3 Zinc-Blende structure of a III-V semiconductor. Courtesy of[START_REF]Structures of Element and Compound Semiconductors[END_REF]

Figure 1 . 4

 14 Figure 1.4 Bandgap energy versus lattice constant of the most common III-V semiconductors. Solid circles indicate a direct bandgap binary material while empty circles indicate an indirect bandgap binary material. The curves show the band gap for some of the ternary alloys, where solid lines indicate direct bandgaps and dashed lines indicate indirect bandgaps. Courtesy of[START_REF] Bayo | Theory of elasticity and electric polarization effects in the group-III nitrides[END_REF]

Figure 1 . 5

 15 Figure 1.5 Integration of III-V on Si chip for integrated photonics. Courtesy of[START_REF]Integrated Photonics Laboratory[END_REF]

Figure 1 . 6

 16 Figure 1.6 Working principle of a PEC cell with a n-doped semiconductor. Courtesy of[START_REF] Maeda | Photocatalytic water splitting using semiconductor particles: History and recent developments[END_REF]

Figure 1 . 7

 17 Figure 1.7 Hybrid integration (a) and monolithic integration (b) of III-V semiconductors on the Si substrate.Courtesy of[START_REF] Kostrzewa | InP dies transferred onto silicon substrate for optical interconnects application[END_REF][START_REF]Molecular Beam Epitaxy[END_REF]

Figure 1 . 8

 18 Figure 1.8 Cross-sectional transmission electron microscopy (TEM) image of dislocationspropagating through the GaAs grown on Si substrate. Courtesy of[START_REF] Luxmoore | III-V quantum light source and cavity-QED on Silicon[END_REF]

Figure 1 . 9

 19 Figure 1.9 Cross-sectional TEM image of a two-dimensional dislocation network confined at the AlSb/Si interface. Courtesy of [41]

Figure 1 .

 1 Figure 1.11 TEM cross-section micrograph of the GaAs/AlGaAs laser structure on Si substratewith relaxed graded GeSi buffer layers. Courtesy of[START_REF] Groenert | Monolithic integration of room-temperature cw GaAs/AlGaAs lasers on Si substrates via relaxed graded GeSi buffer layers[END_REF]

Figure 1 .

 1 Figure1.12 Core-shell GaAs/InGaP nanowires grown on Si[START_REF] Argaman | Density functional theory: An introduction[END_REF]. The epi-structure is shown in the inset. Courtesy of[START_REF] Svensson | Monolithic GaAs/InGaP nanowire light emitting diodes on silicon[END_REF]

Figure 1 .Figure 1 .

 11 Figure 1.13 Cross-section HRSTEM-HAADF image of GaP grown on Si, showing the good atoms alignment at the interface. Courtesy of [58]

Figure 1 .

 1 Figure 1.15 A 2D schematic representation of (a) elastic strain and (b) plastic relaxation at the interface of a heteroepitaxial thin layer grown on a lattice mismatched substrate. Courtesy of [71]

  They are found in close-packed structures where atomic layers are arranged in a certain stacking order. Most of the III-V semiconductors are formed in the zinc-blende structure, which includes two face centered cubic (fcc) sublattices. The crystallographic order in the (111) plane can be expressed as a ABCABCABC ordering atomic sequence. Stacking faults are formed when the sequence is altered and it is modified as ABCACABC or ABCAABCABC which are called respectively intrinsic or extrinsic stacking faults (as shown in Fig. 1.16).

Figure 1 . 16

 116 Figure 1.16 Representation of (a) intrinsic (b) extrinsic stacking faults. Courtesy of[START_REF] Németh | Transmission electron microscopic investigations of heteroepitaxial III/V semiconductor thin layer and quantum well structures[END_REF]

Figure 1 . 17

 117 Figure 1.17 Schematic representation of a microtwin model in a zinc-blend GaP structure.Courtesy of[START_REF] Skibitzki | GaP collector development for SiGe heterojunction bipolar transistor performance increase: A heterostructure growth study[END_REF]

Figure 1 . 18

 118 Figure 1.18 Sketch of the Antiphase boundaries with wrong V-V or III-III bonds marked by the dashed line. Courtesy of [73]

Figure 1 . 19

 119 Figure 1.19 Atomic configuration of stoichiometric (a) or non-stoichiometric APB (b).

Figure 1 . 20

 120 Figure 1.20 Specific formation energy of the {110}, {111}, {112}, and {113} APB's in GaAs (a)and GaP (b). "Relaxed" and "unrelaxed" refer to calculations that include or exclude relaxation of atomic coordinates within the supercell. Courtesy of[START_REF] Rubel | Formation Energies of Antiphase Boundaries in GaAs and GaP: An ab Initio Study[END_REF]

Figure 1 .

 1 Figure 1.21 (a) A III-V material grown on a Si (001) substrate in cross-sectional view, where the Si, Group-V and Group-III atoms are depicted in black, blue and red, respectively. APBs are highlighted by green lines. According to Kroemer's description, APBs can be generated by mixed nucleation on the same terrace (a), or it can be the consequence of silicon monoatomic steps (b, c&d). For simplicity, it is considered in (b), (c) and (d) that the nucleation of the III-V layer on Si starts always with the same type of atom. In this configuration, D-steps do not generate APBs. (c) and (d) show APBs generated by S-steps on the Si surface. (c) APB generated in the {110} plane perpendicular to the interface and (d) APBs self-annihilation model base on generation of APBs along {111}. Courtesy of[START_REF] Lin | Antiphase domain annihilation during growth of GaP on Si by molecular beam epitaxy[END_REF]

1. 4 . 2 . 2

 422 Generation, propagation and burying of antiphase domains: recent findings In 2018, I. Lucci et al. proposed a more comprehensive picture: an alternative model [3] based on Density Functional Theory calculation and the extensive in-depth atomicscale characterization of a variety of GaP/Si, AlSb/Si and AlN/Si samples. Conclusions of this study can be summarized as follows. Firstly, there is always a partial wetting between III-V semiconductors and Si, thus leading to the formation of pure 3D Volmer-Weber growth mode, irrespective of the lattice mismatch (Fig. 1.22). As an example, they provided the energy-dispersive x-ray spectroscopy (EDX) image displayed in Fig. 1.22a,

Figure 1 .Figure 1 .

 11 Figure 1.22 3D islands in various III-V on Si material systems. STEM-EDX image of GaSb/AlSb layers grown on 6° miscut Si, showing the Ga and Al concentration (a). Plan view STM imaging of a 3 nm-thick GaP deposition on 6° miscut Si. The 20 nm × 20 nm inset shows the atomically resolved morphology of the individual island marked in black. APB is highlighted in green (b).Courtesy of[3]

Figure 1 .

 1 Figure 1.24 (a) Sketch of a {110} supercell containing APBs. APBs are indicated by thick vertical lines. (b) Band structures for irreducible GaP unit cell (thick solid lines), and APB spacing supercell (dots). Shaded areas correspond to energy space occupied by bulk states. Courtesy of[START_REF] Tea | Theoretical study of optical properties of anti phase domains in GaP[END_REF]

Figure 1 . 25

 125 Figure 1.25 Imaginary part of the calculated dielectric function for direct optical transitions ɛi(q→0,ω). Results for bulk and APB structure are drawn with solid and dashed lines, respectively.The inset shows the same data on a log scale. Courtesy of[START_REF] Tea | Theoretical study of optical properties of anti phase domains in GaP[END_REF]

Figure 2 . 1

 21 Figure 2.1 Veeco di Innova setup of the FOTON institute at INSA Rennes [87].

Figure 2 . 2

 22 Figure 2.2 Force profile when the distance between the tip and the surface changes [88].

Figure 2 . 3

 23 Figure 2.3 Schematic diagram of Conductive AFM system [89].

X

  -ray diffraction (XRD) is a powerful nondestructive technique for characterizing crystalline materials. It provides information on structures, phases, preferred crystal orientations (texture), and other structural parameters, such as crystallinity, strain, and crystal defects[START_REF]Chapter 3 -Methods for Assessing Surface Cleanliness[END_REF]. X-ray diffraction is based on Bragg's law (nλ=2dsinθ). A monochromatic beam of X-rays is sent to the sample, and reflected X-rays are detected by a detector, as shown in Fig.2.4.

Figure 2 . 4

 24 Figure 2.4 Schematic representation of the X-ray diffraction technique [91].

Figure 2 . 5

 25 Figure 2.5 XRD pattern for one GaP0.67Sb0.33/Si sample.

Fig 2 . 6

 26 shows the RSM around the (224) Bragg peak for the GaP0.67Sb0.33/Si sample (Sx and Sz are the projected coordinates in the right handed Cartesian, with z axis parallel to the surface normal).

Figure 2 . 6

 26 Figure 2.6 The RSM around[START_REF] Vandewal | On the origin of the open-circuit voltage of polymer-fullerene solar cells[END_REF] for the GaP0.67Sb0.33/Si sample (Sx and Sz are the projected coordinates in the right handed Cartesian, with z axis parallel to the surface normal)[START_REF] Nguyen Thanh | Structural and optical analyses of GaP/Si and (GaAsPN/GaPN)/GaP/Si nanolayers for integrated photonics on silicon[END_REF]. The black dot line the represents the full relaxation line.

  the 𝐶 12, 𝐶 11 are the elastic constants of GaPSb. The Sb content x is then deduced based on the Vegard's law.

  Chemical mechanical polishing (CMP) is a process of smoothing surfaces with the combination of chemical and mechanical forces[94]. The process uses an abrasive and corrosive chemical slurry in conjunction with a polishing pad and retaining ring, typically of a greater diameter than the wafer. The pad and wafer are pressed together by a dynamic polishing head and held in place by a plastic retaining ring. The dynamic polishing head is rotated with different axes of rotation (not concentric). This removes material and tends to even out any irregular topography, making the wafer flat or planar. The schematic diagram of functional principle of CMP is shown in Fig.2.7.

Figure 2 . 7

 27 Figure 2.7 Functional principle of CMP [95].

Figure 2 . 8

 28 shows the SEM images of one III-V/Si sample before and after CMP, from which we can find that the surface of the sample becomes much smoother than the as-grown one (before CMP). Besides, during this thesis, this technique was further improved by J. Le Pouliquen and Y. Léger to highlight the APBs emerging at the surface of III-V/Si samples, by combining CMP and additional etching[START_REF] Saleem-Urothodi | Loss assessment in random crystal polarity gallium phosphide microdisks grown on silicon[END_REF].

Figure 2 . 8

 28 Figure 2.8 SEM image of a III-V/Si sample before (a) and after (b) CMP.

Figure 2 . 9

 29 Figure 2.9 shows the principle of the photoluminescence (PL) experiment. Under light illumination, the electrons of the semiconductor on ground states (in valence bandor on some trap states) are excited to higher energy states by absorbing photons of the pumping laser beam. The electrons photo-generated remain in the excited state for a very short time (around 10 -12 s). Then, the electrons relax to lower states (the bottom of the conduction band or some localized states) mostly by phonon emission. Finally, the electrons will go back to the ground states (recombine with holes) and at the same time give rise to photoluminescence.

Figure 2 . 9

 29 Figure 2.9 Principle of photoluminescence (PL) spectroscopy [96].

Figure 2 .

 2 Figure 2.10 FOTON experimental setup at room (lower part) and low temperature (upper part) [97].

  2.11 red line), which needs to be filtered out and discarded during the Raman measurement. Raman scattering, on the other hand, is due to inelastic scattering. The incident photons are scattered by the vibrational or rotational modes of molecules/dimers/bonds in the sample, which causes a change in the energy and frequency of the scattered light as compared to the incident light. During the scattering process, the incident photon may release energy to the lattice (or to some molecules) through vibrations/rotations, and thus the photon is scattered with a redshifted frequency; this is the so-called Stokes shift (Fig. 2.11). On the contrary, during the scattering, transfer of vibrational or rotational energy from the lattice (or from the molecule) to the scattered photon may occur. The light output will have a blue-shifted frequency, which is called as anti-Stoke shift (Fig. 2.11). In most cases, the photons collected and analyzed during the Raman scattering experiments are corresponding to the Stokes photons (referred to as Stokes lines), which corresponds to the most probable process, thus leading to larger light intensity.

Figure 2 .

 2 Figure 2.11 Energy-level diagram showing the states involved in Raman spectra [98].

Figure 2 . 13 48 (

 21348 Figure 2.13 Schematic diagram of 3-electrodes photoelectrochemical measurements.

Figure 1 . 14

 114 Figure 1.14 Chopped light LSV curves of semiconductor working as photocathode(a) [103] and photoanode(b) [104] given for illustration.

Figure 1 .

 1 Figure 1.14a and b show the chopped light LSV curve of some semiconductorsworking as photocathode and photoanode, respectively. The current under light-off state corresponds to the dark current and the current under light-on state contains both photocurrent and dark current. From these two figures, we can find that for all the curves, both the photo-and dark currents improve with the increase of applied voltage. Corrosion is one common phenomenon for many semiconductor photoelectrodes, which can be reflected in cyclic voltammetry curves. If the current of voltammetry curve evolves toward lower values with time, it may be the signature that corrosion occurs during the PEC reactions.

  2.15(left)). Another representation used is the Bode plot (Fig.2.15(right)), where the impedance and phase angle are plotted against frequency, which can be helpful to find capacitive or inductive effects of electrochemical systems.In this thesis, we will use the Nyquist representation. One has to keep in mind that, with this representation, any change of the semi-circle radius corresponds to an evolution of the impedance, for different samples or different experimental conditions.

Figure 2 . 51 2. 3 . 4

 25134 Figure 2.15 Electrochemical Impedance Spectroscopy representation by the Nyquist plot (left) and the Bode plot (right) [108].

Figure 2 .

 2 Figure 2.16 a and b show the Mott-Schottky curves of p-doped[START_REF] Koushik | Plasma-assisted atomic layer deposition of nickel oxide as hole transport layer for hybrid perovskite solar cells[END_REF] and n-doped[START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] semiconductor photoelectrodes, respectively.

Figure 2 .

 2 Figure 2.16 Mott-Schottky curve of a p-doped (a) [109] and a n-doped (b)[START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF]

55 2. 4 . 1 . 4

 55414 Generalized gradient approximation (GGA)

  uniform k-point mesh. Secondly, based on the electron density obtained above, the eigenvalues along high symmetry k-point paths are determined. Finally, the band structures (eigenvalues versus k point) can be plotted in 2D along a specific k direction (Fig. 2.18a) or in 3D in a given k plane (Fig. 2.18b).

Figure 2 . 18

 218 Figure 2.18 The electronic band structure along one 2D k-path in reciprocal space (a) and the 3D plotting in a k plane (b) [135].

Figure 2 .

 2 Figure 2.19 shows the principle of band folding effect.Figure 2.19a and b show the band structures of a polymer obtained with a supercell (two atoms per unit cell) and with a unit cell (one atom per unit cell), respectively. We can see the band folding effect in Fig.2.19a which is produced by folding the band structure in Fig.2.19b and the corresponding folding process is described in Fig.2.19c.

Figure 2 . 19

 219 Figure 2.19 Band folding in one-dimensional polymers [136]. (a) Band structure of a polymer in which there are two atoms per unit cell. The two branches intersect at k = π/(2a). (b) Band structure of a polymer in which there is one atom per unit cell. The Brillouin zone is doubled because the unit cell is one half as it is in (a). (c) Production of band structure in (a) by folding band structure in (b).

64 2. 4 . 5 3 . 5 .

 644535 Common DFT computational details In this work, the DFT calculations mainly contain two parts. The first one involves the electronic band structure, phonon and electron-phonon coupling calculations for InPbased stoichiometric APB structures and the corresponding results are shown in Chapter The other one is mainly about the band structure calculations of nonstoichiometric APB structures with GaP, GaAs and GaSb materials and the corresponding results are shown in Chapter For the DFT calculations of InP-based stoichiometric APB structures (Chapter 3), a Monkhorst-Pack 4x1x6 grid of k points was used for the geometry optimization step and the structures were relaxed until the Hellmann-Feynman forces on the atoms were less than 10 -4 eV/Å. Band structure and phonon calculations were performed with the density functional theory by combining the VASP and PHONOPY package codes. DFT calculations were performed using the plane-wave projector augmented wave (PAW) method as implemented in the VASP code. The standard generalized gradient approximation (GGA) parameterized by Perdew-Burke-Ernzerhof (PBE) was used for structure optimization and phonon calculation, and the Heyd-Scuseria-Ernzerhof (HSE) hybrid functional was chosen for the band structure calculation, with energy cutoff of 500eV. Besides, the symmetry information and simulated Raman spectrum of the localized phonons around APB were extracted through the phono3py [144] package codes.

65 for

 65 the DFT calculations, which are presented in Chapter 5. Around 20 Å and 30 Å vacuum were added on the short and long structures, respectively, which is thick enough to avoid any interaction between slab surfaces. The structures were optimized with Hellmann-Feynman forces criterion less than 10 -4 eV/Å, and based on a Monkhorst-Pack 6x6x1 grid of k points. The DFT band structure calculations were performed by using HSE hybrid functional and metaGGA (TB-mBJ) potential, based on the VASP code, with an energy cutoff of 500 eV. Besides, in order to study the topological properties, the topological related (including band structures and Berry phase) calculations were carried out by generating ab initio tight-binding Hamiltonians from the maximally localized

Figure 3 . 1

 31 Figure 3.1 Sketch representing the atoms arrangement in a stoichiometric APB.

Figure 3 .

 3 Figure 3.2a is a schematic representation of the InGaP/SiGe/Si sample's structure.Experimental scanning transmission electron microscopy (STEM) plan-view and crosssectional images are shown in Fig.3.2b and Fig.3.2c, respectively. The sample is composed of three main layers: a Si substrate, a 760 nm-thick Si0.433Ge0.567 buffer layer and a 210 nm-thick InGaP III-V layer (Fig.3.2a) with a nominal In content of 30%. The last epitaxial layer was grown at relatively high growth temperature (500°C) and low growth rate, to enhance non-equilibrium kinetic effects during the growth of the metastable (In,Ga)P alloy. Especially, In atoms are well-known for their ability to efficiently migrate at the surface of InGaP semiconductors during the growth, due to their low binding energy to the surface at high temperature. In this sample, APBs in InGaP are propagating almost vertically until the surface (boundary between dark grey and clear grey in Fig.3.3c), giving rise to 2D vertical homovalent singularities. We note here that, even if the TEM images are not presented here with a high resolution or may seem a bit blurry, our laboratory has a strong experience in studying these samples and can guarantee that the observed contrast can be attributed unambiguously to APBs. The 2D vertical

Figure 3 . 2 Fig. 3 . 5 .

 3235 Figure 3.2 Schematic image (a) as well as plan view (b) and cross-sectional (c) STEM images of the studied InGaP/SiGe/Si sample.

Figure 3 .

 3 Figure 3.3 a, Cross-sectional bright-field (BF) STEM image of the InGaP/SiGe/Si sample. The black line in the middle and the hump on the surface corresponds to the APB and surface nanostructure, respectively. b-e, Cross-sectional EDX elemental mapping images of indium (b), gallium (c), phosphorus (d) and silicon (e). f-h, Plan-view EDX mapping images of indium (f), gallium (g), phosphorus (h). Two kinds of In-rich nanostructures are observed: pure InP surface nanostructures and 2D In-rich vertical quantum wells (QWs) surrounding the APB singularities.

Figure 3 . 4

 34 Figure 3.4 Plan-view bright field image of the (In,Ga)P/SiGe/Si sample.

Figure 3 . 5 4 .

 354 Figure 3.5 Plan-view elemental mapping images of indium (a), gallium (b), phosphorus (c) and indium-gallium-phosphorus integration (d), corresponding to Fig. 3.4.

Figure 3 . 6

 36 Figure 3.6 Scanning electron microscopy (SEM) plan-view images of the sample before (a) and after (b) CMP.

Figure 3 . 7

 37 Figure 3.7 shows the Photoluminescence (PL) of the sample before(a) and after(b) CMP. After CMP, the low PL energy peak disappears which indicates that the low energy PL peak comes from the surface nanostructures, which were effectively removed by the CMP process.

Figure 3 . 7 Fig. 3 . 9 .

 3739 Figure 3.7 Photoluminescence of the sample before (a) and after (b) CMP.

Figure 3 . 8 X

 38 Figure 3.8 X-ray diffraction (XRD) patterns of the reported as-grown InGaP/SiGe/Si sample (before polishing).

Figure 3 . 9 X

 39 Figure 3.9 X-Ray diffraction reciprocal space mappings around (224) and (004) reflection on the sample before and after CMP. The black dotted vertical and inclined lines represent respectively the full strain lines and full relaxation lines.

Figure 3 . 10

 310 Figure 3.10 Schematic image of a possible APB growth mechanism explaining the preferential In incorporation.

Figure 3 . 11

 311 Figure 3.11 The 32 atoms APB-supercell (APB in an InP matrix) used for DFT band structure and phonon calculations.

Figure 3 . 12 78 An

 31278 Figure 3.12 Model 32 atoms bulk-supercell structure of InP used for DFT calculations.

Figure 3 . 13

 313 Figure 3.13 Illustration of the connection between the layer and the space groups. The wickoff positions for In and P atoms are given. The connection between the irreducible representations at the Г point in both cases is also indicated. The axis perpendicular to the layer is z and y respectively for the standard settings of layer and Pnnm space groups.

Figure 3 . 15

 315 Figure 3.15 Bandstructures of bulk InP for the supercell (48 atoms) computed by DFT using the HSE functional.

Fig. 3 . 3 . 15 )

 3315 [START_REF] Vegard | Die Konstitution der Mischkristalle und die Raumfüllung der Atome[END_REF] gives the energies of the various electronic states labelled according to the IRs of the Pnnm space group. As indicated in Fig.3.16, the energy splittings of pairs of states are expected to vanish for large APB separations and thus large n values. Therefore, according to the labelling of the layer space group, the top of the VB is related to the Ag IR. It is also clearly identified as a localized state for a single APB according to its electronic density profile in the APB supercell. The second VB state is also localized on the APB and corresponds to Bg layer group symmetry. The VB states which appear next about 0.5eV below the top of the VB are not located anymore on the APB and correspond to InP bulk VB states (Fig.3.14). From these simulations, we conclude that a single APB yields two VB states located into the band gap of the surrounding InP matrix. On the other hand, the first electronic states at the bottom of the CB appear to be located into the InP matrix and correspond to a Bu layer group symmetry. Overall, the bandgap of the APB-structure is equal to about 0.93 eV (Fig. 3.14). Therefore, we notice that the energy difference 81 between the bottom of the CB and the first bulk InP-like VB states amounts to about 1.43eV in nice agreement with the electronic band gap of bulk supercell InP (1.40 eV, Fig. at the same level of theory. To summarize, the overall picture emerging from this analysis, is that a type II band alignment between the APB and the surrounding InP matrix is shown, yielding two APB localized VB states. These 2D VB electronic states belong to Ag and Bg IR in a layer group representation, which are marked in Fig. 3.14 together with the layer group representation Bu of first electronic states at the bottom of the CB.

Figure 3 . 16

 316 Figure 3.16 The energies of the VB localized electronic states on APB labelled according to the IRs of the Pnnm space group at Г point vs the APB supercells size (n=32/48/64 atoms). The electronic states with B3g and B2g symmetries almost overlap.

Figure 3 . 17

 317 Figure 3.17 Partial charge density of the VBM and CBM at the Γ point for a 48 atoms supercell, integrated in the planes parallel to APBs. The dotted line corresponds to the APB plane.

Figure 3 . 18

 318 Figure 3.18 Phonon dispersion of the 32 atoms APB-structure. The phonon modes highlighted by orange, blue, magenta, green and purple dotted lines correspond to additional localized modes related to APB planes. They are labelled according to layer group IRs at Г point.

Figure 3 . 19

 319 Figure 3.19 Phonon dispersion of bulk InP for the supercell (32 atoms)

1 ) .

 1. Each two phonons grouped by pairs have almost the same vibration behavior. The IR of the layer group of the new phonon modes are marked in the Fig. 3.18.

Figure 3 . 20 85 Figure 3 .

 320853 Figure 3.20 Partial vibrational (phonon) density of states of the middle InP atoms (a), In-In APB atoms (b) and P-P APB atoms (c), corresponding to the atoms encircled by the black, red and green dotted lines respectively in Figure 3.11.

Figure 3 . 21

 321 Figure 3.21 Schematic images of vibration behaviors of the five APB localized phonon modes at Γ point: 5.3 THz-Ag(a); 5.9 THz-Bu(b); 7.8 THz-Au(c); 8.0 THz-Bu(d); 12.6 THz-Ag(e).

Figure 3 . 22

 322 Figure 3.22 Simulated Raman spectrum of the five localized phonons at the Г point, from DFT.

  central frequency (around 200 cm -1 , i.e. 6 THz) shows a reasonable agreement with the Ag Raman active mode with frequency around 5.3 THz localized at the APB calculated previously. Therefore, the Raman peak 1 in the experimental data is attributed to the Ag localized vibration with theoretical frequency around 5.3 THz. The Raman signal of the 87 experimental data in the range of 400-500 cm -1 should contain the contribution from the second Ag APB localized phonon with a theoretical frequency around 12.6 THz, but it is most probably hidden by the contributions related to the other parts of the sample. The hyperspectral mapping in the inset displays the spatial distribution of the Raman peak 3 of GaP TO (green area) and Raman peak 1 (red area) vibration modes. The clear observation of remaining red and green area in the spatial superposition of the two vibrational modes indicate that the two phonon vibrations do not originate from the same region. Based on the EDX images, we know that there are two main structures in the InGaP layer (after CMP): the main In-poor phase (InGaP with around 18% In) and the Inrich phase located around the 2D APB singularities. Therefore, we attribute peak 3 observed in the Raman spectrum to lattice vibrations inside the In-poor domains away from the APB, while peak 1 comes from phonons localized in the APB homovalent singularities. The signature of the localized phonons (peak 1) in Raman experiments is relatively weak because of the small volume fraction of APBs. But this does not entail a low oscillator strength. From the structural investigations, the band structure and phonon dispersion calculations and from Raman measurements, we can conclude that the 2D APB singularities simultaneously bring in both carrier and phonon localization in the same region of the sample, which are the basic ingredients to create uncommon carrier-phonon interaction effects. The existence of an unexpected low-frequency localized Ramanactive vibrational mode at about 6 THz is in particular consistent with first-principles predictions.

Figure 3 . 23

 323 Figure 3.23 Experimental Raman spectrum. The Raman peaks 2, 3, 4, 5, 6, 7, correspond respectively to Ge-Ge bonds, GaP transverse optical (TO) phonons, InP longitudinal optical (LO) phonons, GaP LO phonons, Ge-Si bonds and Si-Si bonds. The Raman peak 1 shows a good corresponding to the 5.3 THz phonon localized in the APB, calculated previously. The inset shows the plan-view spatial distribution on the sample of the Raman peaks 1 (around 200 cm -1 , i.e. 6 THz) and 3 (GaP TO).

10 .

 10 optical constants of the InGaP/SiGe/Si sample were measured by spectroscopic ellipsometry (VASE) at room temperature in the 1.5-5 eV photon energy region. A Tauc-Lorentz model with 2 oscillators was used to fit the ellipsometry data. The measurements performed by ellipsometry are in colored dots and the fits are plotted in black continuous lines in Fig. 3.24a. These curves give the real and imaginary parts of the so-called "pseudo" dielectric function which can be obtained from Eq. 2-9 and Eq. 2-From the fitting model, the refractive index (n) and extinction coefficient (k) were extracted, as shown in Fig. 3.24b. From the imaginary part, absorption spectrum for the InGaP layer has been deduced, as shown in Fig. 3.24c.

Figure 3 .

 3 Figure 3.24 a, Experimental ellipsometry spectra of Ɛr and Ɛi (red and blue lines) and comparison with theoretical curves by using a 2-oscillators Tauc-Lorentz model (black lines). b, n (optical index real part) and k (optical index imaginary part) optical constants extracted from the fitting model. c, Absorption spectrum deduced.

Figure 3 . 25

 325 Figure 3.25 Photoluminescence and absorption spectra at low temperature (15 K). The inset corresponds to the PL spectrum in high magnification.

Figure 3 .Figure 3 .

 33 Figure 3.26a displays the power-dependent PL spectra, from which it can be observed that the whole spectrum (including small peaks and the broad contribution) remains proportionally unchanged with increasing pumping power, which further indicates the broad PL peak contains and only contains multi-phonon replicas.Figure 3.26b displays the logarithm plot of the integrated PL intensity versus excitation power.

Figure 3 . 3 . 27 .

 3327 Figure 3.26 (a) Power-dependent Photoluminescence spectrum at 15 K. (b) Logarithm plot of the integrated PL intensity versus excitation power.

Figure 3 .

 3 Figure 3.27 Temperature-dependent PL. The inset shows the corresponding Arrhenius plot of the integrated PL intensity.

Figure 3 . 93 Figure 3 .

 3933 Figure 3.28 Time-resolved photoluminescence decay of the polished InGaP/SiGe/Si sample.

Figure 3 . 29

 329 Figure 3.29 Mechanism of the Photoluminescence with a spatial representation (a) or within the electron-phonon interaction configuration coordinate point of view (b).

Figure 3 .

 3 [START_REF] Tsai | High-efficiency two-dimensional Ruddlesden-Popper perovskite solar cells[END_REF] shows the experimental data and the theoretical fitting achieved by the Franck-Condon principle. Accurate fitting of the experimental curve was only obtained with a S parameter of 8 and a ZPL positioned at 2.097 eV.

Figure 3 . 30

 330 Figure 3.30 Experimental data (grey dots) and theoretical fitting (red line) of the broad PL peak based on the Franck-Condon model. The small multicolored peaks correspond to the distinct phonon replicas contributions.

Figure 3 . 31

 331 Figure 3.31 Schematic plot of the 2D vertical In-rich structure, embedding APB planes with strong electron-phonon interactions.

97

 97

3 . 98 Figure 3 .

 3983 Figure 3.32 EPI matrix elements for the phonon branches corresponding to the localized polar phonon modes at the  point as a function of q for (a),(c) the top of the valence band (m=n=VBM) and (b),(d) the bottom of the conduction band (m=n=CBM) along the X and Z in-plane directions of the APB (The -X and -Z paths corresponding to the [001] and [110] crystallographic directions, Fig. 3.11).

100 3. 10 Summary 101 Chapter 4 :

 100101014 photoluminescence processes was clarified, by combining transmission electron microscopy, X-ray diffraction, ab initio calculations, Raman spectroscopy and photoluminescence experiments. 2D localization and layer group symmetry properties of homovalent electronic states and phonons were studied by first principles methods, leading to the prediction of a type II band alignment between the APB and the surrounding semiconductor matrix. A Huang-Rhys factor of 8 was finally experimentally determined from the APB emission line, underlining that a large and unusually strong electronphonon coupling can be achieved by 2D vertical quantum confinement in an undoped III-V semiconductor. This work extends the concept of electron-phonon interaction to 2D vertically buried III-V homovalent nanoobjects and therefore provides different approaches for material designs, vertical carrier transport, heterostructure design on silicon and device applications with weakly polar semiconductors.

  the next sections, we assess the potential of a new monolithic III-V/Si tandem materials association (GaP1-xSbx/Si) for solar water splitting. The GaP1-xSbx was directly grown on Si substrate by molecular beam epitaxy (MBE), with the aim to keep the greatest merit of the tandem system, i.e. an efficient light absorption at 1.7 eV on the top of the 1.1 eV silicon bandgap. The bandgaps and band alignments of the GaP1-xSbx alloys are carefully and comprehensively studied over the whole Sb range, by combining the experimental data with tight binding (TB) theoretical calculations, and its possible use as a photoelectrode is discussed.

1. 7 / 1 .

 71 photons are absorbed and low-energy photons are transmitted and harvested by the Si substrate or bottom cell, leading potentially to very large light absorption of the sunlight (as shown in Fig. 4.1b). Then the photo-induced charges (electrons and holes) are generated in both layers, and, depending on the doping, one kind of charges flows toward the illuminated surface (the top surface of GaPSb layer) to generate H2 or O2 and the other one flow toward the back contact which is connected to the Si substrate (as shown in Fig. 4.1a) and is further extracted to feed the counter-electrode.

Figure 4 . 1

 41 Figure 4.1 (a) Schematic of the proposed GaP1-xSbx/Si tandem device for PEC water splitting. (b) Sunlight absorption illustration of GaP1-xSbx/Si tandem system with a 1.7/1.1eV bandgap combination for water splitting which guarantees large light absorption.

Figure 4 . 2

 42 Figure 4.2 shows the X-ray diffraction (XRD) ω/2θ profiles for the three samples.

Figure 4 . 2 X

 42 Figure 4.2 X-Ray Diffraction profiles for the three MBE-grown GaP1-xSbx/Si samples with different Sb contents: GaPSb-1 (i.e. GaP0.67Sb0.33) (a), GaPSb-2 (i.e. GaP0.40Sb0.60) (b) and GaPSb-3 (i.e. GaSb) (c).

Figure 4 . 3 X

 43 Figure 4.3 X-Ray diffraction reciprocal space maps around (004) and (115) for the three samples, correspondingly (Sx and Sz are the projected coordinates in the right handed Cartesian, with z axis parallel to the surface normal) [92]. The red lines represent the full relaxation lines.

Figure 4 . 4

 44 Figure 4.4 shows the scanning electron microscopy (SEM) images of the three GaP0.67Sb0.33/Si, GaP0.40Sb0.60/Si, GaSb/Si samples on plan-view (Fig.4.4 a, c, e) and cross-sectional view (Fig.4.4 b, d, f). From these images, it can be observed that the samples GaP0.40Sb0.60/Si and GaSb/Si exhibit relatively smooth surfaces as compared with the other GaP0.67Sb0.33/Si sample. The roughness observed is attributed to both emergence of some crystal defects such as residual dislocations or antiphase boundaries and residual stress in the sample. The corresponding atomic force microscopy (AFM) images are given in the Fig. 4.5 and the RMS (root-mean-square) roughnesses of the surfaces were found to be 22.80 nm for GaP0.67Sb0.33, 9.32 nm for GaP0.40Sb0.60 and 7.91 nm for GaSb. It suggests an improvement of the crystal quality with increasing Sb content. This is also confirmed by a very significant Bragg peak sharpening observed on reciprocal space

Figure 4 . 4

 44 Figure 4.4 Scanning electron microscopy (SEM) images of the GaP0.67Sb0.33, GaP0.40Sb0.60, GaSb samples before CMP on plan view (a),(c),(e) and cross-sectional view (b),(d),(f).

Figure 4 . 5

 45 Figure 4.5 Atomic force microscopy (AFM) images of the GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b), and GaSb (c) samples before CMP.

Figure 4 . 6

 46 Figure 4.6 Scanning electron microscopy (SEM) plan view images of the GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b) and GaSb (c) samples after CMP.

Figure 4 . 7

 47 Figure 4.7 Atomic force microscopy (AFM) images of the GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b)and GaSb (c) samples after CMP.

109 the

 109 Fig. 4.9 and Fig.4.10 (red curves).

Figure 4 . 8

 48 Figure 4.8 Experimental ellipsometry spectra of Is and Ic (red and blue lines) for two incidence angles and comparison with theoretical curves by using a 2-oscillators Tauc-Lorentz model (black lines) for the GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b) and GaSb (c) samples, respectively.

Figure 4 .

 4 Figure 4.9 n (optical index real part) and k (optical index imaginary part) optical constants of GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b) and GaSb (c) extracted from the fitting of Fig.4.8 (a), (b), (c), respectively.

Figure 4 . 10

 410 Figure 4.10 Optical absorption spectra of the GaP0.67Sb0.33 (a), GaP0.40Sb0.60 (b) and GaSb (c) semiconductors. The red lines were deduced from ellipsometry measurement and the black lines were obtained from TB calculations.

4. 12 )

 12 , and show good agreement.

Figure 4 .

 4 Figure 4.11 (a) Experimental ellipsometry spectra of Is and Ic (red and blue lines) and comparison with fitting curves obtained with the Tauc-Lorentz model (black lines) for GaP. (b) n (optical index real part) and k (optical index imaginary part) optical constants of GaP extracted from the fitting of (a).

Figure 4 . 12

 412 Figure 4.12 Optical absorption spectra of GaP (a) and GaSb (b). The red lines were deduced from ellipsometry measurements performed on our GaP/Si and GaSb/Si samples. The experimental data (green dots) and the theoretical ones (black lines) are obtained from the literature[START_REF] Adachi | Optical dispersion relations for GaP, GaAs, GaSb, InP, InAs, InSb, AlxGa1-xAs, and In1-xGaxAsyP1-y[END_REF].

Figure 4 . 13

 413 Figure 4.13 Tauc plot of (αhν) k versus photon energy (hν) for GaP (a), GaP0.67Sb0.33 (b), GaP0.40Sb0.60 (c) and GaSb (d) epilayers grown on silicon.

113 Figure 4 . 14

 113414 Figure 4.14 Room temperature band structures of the bulk unstrained GaP (a), GaP0.67Sb0.33 (b), GaP0.40Sb0.60 (c) and GaSb (d) obtained by TB calculations.

Figure 4 . 15

 415 Figure 4.15 Room temperature bandgaps of GaP1-xSbx alloys with different Sb contents. The green, red and purple solid lines are theoretical curves computed by tight-binding calculations, corresponding to Γ, L and X valleys, respectively. The separated error bar lines correspond to the bandgaps determined experimentally. The black square dots correspond to the bandgaps of GaP and GaSb given in ref.[START_REF] Panish | Temperature Dependence of the Energy Gap in GaAs and GaP[END_REF][START_REF] Wu | Photoluminescence of high-quality GaSb grown from Ga-and Sb-rich solutions by liquid-phase epitaxy[END_REF].

Figure 4 . 16

 416 Figure 4.16 Band alignments of Γ, L and X valleys of GaP1-xSbx and indirect bandgap of Si with respect to water oxidation and reduction potentials. The relative position of redox potentials of water are shown as green bars at pH=0.

Figure 4 . 18

 418 Figure 4.18 Voltammetry curve of GaP0.67Sb0.33/Si photoanode without coating (red line) and GaP wafer (green line) in 1.0 M KOH (pH=14) electrolyte under simulated AM1.5 illumination versus RHE.

  4.19a). Besides, Fig. 4.19bshows the temperature-dependent PL of the GaP0.67Sb0.33/Si sample. As shown in the previous part, the bandgap of this sample is around 1.7eV, while its maximum PL peak is around 1.35 eV. There is therefore a strong Stokes-shift, which is necessarily a consequence of deep carrier localization in the GaP0.67Sb0.33 layer. Three activation energies 8 meV, 30 meV and 130 meV are extracted. The first one is attributed to a small alloy disorder[START_REF] Alqahtani | Photoelectrochemical water oxidation of GaP1-xSbx with a direct band gap of 1.65 eV for full spectrum solar energy harvesting[END_REF], while the two others represent the detrapping of localized carriers to extended states. Even if the detailed analysis of local PL signatures or phonon replica was not performed on this sample, these PL properties are very similar to the ones of In-rich APB structures discussed in details in Chapter 3, linked to the strong electron-phonon coupling, further confirming the presence of APBs in the GaP0.67Sb0.33/Si photoelectrode.

Figure 4 .

 4 Figure 4.19 a. The cross-sectional high-resolution STEM images of the GaP0.67Sb0.33/Si photoelectrode. b. Temperature-dependent photoluminescence spectra between 15 K (purple) and 150 K (dark red). The inset corresponds to the Arrhenius plot for the two transitions, being Low

  photoanode as a protection layer from (photo-) corrosion. Then, an 8 nm-thick metallic nickel (Ni) layer was deposited onto the surface of the TiO2/GaP0.67Sb0.33/Si photoanode as a catalyst by a thermal evaporator, as shown schematically in Fig.4.20.

Figure 4 . 20

 420 Figure 4.20 Structure of the coated GaP0.67Sb0.33/Si photoanode with protection layer (TiO2) and catalyst (Ni).

Figure 4 .

 4 Figure 4.21a presents the voltammetry curves for GaP0.67Sb0.33/Si coated with TiO2 and Ni and GaP0.67Sb0.33/Si coated with Ni photoanodes in 1.0 M KOH (pH=14) electrolyte under one sun illumination. The GaP0.67Sb0.33/Si coated with Ni (8 nm) showsan onset potential of 0.37 V and a saturated current density of 3.84 mA/cm 2 at 1.23 V versus RHE. In contrast, the GaP0.67Sb0.33/Si coated with both TiO2 (20 nm) and Ni(8 nm)

Figure 4 .

 4 Figure 4.21 (a) Voltammetry curves of GaP0.67Sb0.33/Si coated with Ni 8nm and of GaP0.67Sb0.33/Si coated with TiO2 20 nm and Ni 8 nm photoanodes in 1.0 M KOH electrolyte under simulated AM1.5 illumination versus RHE. (b) Voltammetry curve (scan number 6) for the GaP0.67Sb0.33/Si coated with Ni 8nm and coated with TiO2 20 nm and Ni 8 nm photoanodes. The red voltammetry curve corresponds to the sixth scan of the GaP0.67Sb0.33/Si photoanode without coating.

Figure 4 .

 4 photoanode (see Fig.4.22) was ~53.2% at 400 nm and decreases towards longer wavelengths (> 550 nm). This is due to the large and indirect bandgap of GaP[START_REF] Hu | Amorphous TiO2 coatings stabilize Si, GaAs, and GaP photoanodes for efficient water oxidation[END_REF].

Figure 4 . 22

 422 Figure 4.22 Incident photon-to-current conversion efficiency (IPCE) of the GaP wafer and GaP0.67Sb0.33/Si without coating in 1.0 M KOH electrolyte at 1.23 V versus RHE.

Figure 4 . 23

 423 Figure 4.23 Incident photon-to-current conversion efficiency (IPCE) of the GaP0.67Sb0.33/Si without coating, GaP0.67Sb0.33/Si coating with Ni 8 nm, and GaP0.67Sb0.33/Si coating with TiO2 20 nm and Ni 8 nm in 1.0 M KOH electrolyte at 1.23 V versus RHE.

Fig. 4 . 25 . 4 . 25 .

 425425 Fig. 4.25.The photocurrent of the GaP0.67Sb0.33 photoelectrode without coating decreases very rapidly, which is attributed to (photo-) corrosion. However, after the modification of the GaP0.67Sb0.33 surface by the deposition of Ni (8 nm) as a catalyst, the current was maintained for around 1.30 h with a small decay in the current density and decreased

Figure 4 . 24

 424 Figure 4.24 Photocurrent density versus time (J-t) for GaP0.67Sb0.33/Si without coating and GaP0.67Sb0.33/Si coated with Ni 8 nm photoanodes in a 1.0 KOH (pH=14) aqueous solution under one sun illumination.

Figure 4 . 25

 425 Figure 4.25 Photocurrent density versus time (J-t) for GaP0.67Sb0.33/Si coated with TiO2 20 nm and Ni 8 nm photoanode in a 1.0 KOH (pH=14) aqueous solution under one sun illumination.

127 Chapter 5 :

 1275 photoelectrodes, our research efforts focused on the understanding of the possible involved mechanisms to explain their performances. The presence of non-stoichiometric Antiphase Boundaries in III-V/Si samples was rapidly considered as the most credible explanation to interpret the unusual transport properties observed. This chapter gives a detailed analysis of the novel physical properties of non-stoichiometric APBs, describes how it can be used for the development of III-V/Si photoelectrodes, but also for other applications in photonics, electronics/computing or energy harvesting. Contrary to Chapter 4, I was actively involved in the photoelectrochemical (PEC) measurements in collaboration with ISCR. The Hall effect and C-AFM experiments were performed in collaboration with Institut de Physique de Rennes (IPR). The results presented here were just submitted for publication.

V/

  photoelectrodes. The unusual topological properties of the 2D semimetallic inclusions are also discussed. The unique physical properties of these heterostructures open new paradigms for energy harvesting, photonics, electronics or computing devices, combining high efficiency and predictable properties of inorganic materials and nano-scale charge carriers transport management flexibility usually offered by organic materials.

129

 129

  photoluminescence properties. Therefore, APBs should no longer be considered as nonradiative recombination centers, but as a 2D homovalent singularity in a bulk semiconducting matrix. Due to the insertion of a local translational symmetry breaking, III-III (V-V) bond also adds 2 free holes (2 free electrons) in the semiconductor. While different experimental works highlighted the strong influence of APBs on transport

Figure 5 . 1

 51 Figure 5.1 Schematic of non-stoichiometric APB atomic configuration. The APBs result from the coalescence of two domains marked by A and B (referred to as bi-domain) with different phases.

Figure 5 . 2 131 Figure 5 . 2

 5213152 Figure 5.2 Top-view SEM images of one GaP/Si sample with emerging antiphase domains. The inset corresponds to the same sample after APB developing treatment on the surface.

Figure 5 . 3

 53 Figure 5.3 Plan-view SEM image of one GaPSb/Si sample with emerging antiphase domains after APB developing treatment on surface. The white lines correspond to APBs, due to the faster chemical etching of APBs.

Figure 5 . 4

 54 Figure 5.4 Plan-view SEM image of n-doped GaP wafer.

Figure 5 . 5 133 Figure 5 . 6

 5513356 Figure 5.5 Schematic diagram of in-plane charge distribution. Red (blue) color corresponds to directions, along which vertical APBs are overall charged (neutral).

Figure 5 . 6

 56 Figure 5.6 AFM image of the GaP/Si sample after the APB developing treatment on the surface revealing that APBs follow different local crystallographic directions.

Table 5 . 1 :

 51 Schematic illustration of Hall effect measurements and the relevant results (resistivity, doping concentration and mobility) of four different samples: bi-domain GaP membrane on insulating layer with emerging APBs; mono-domain GaP membrane on insulating layer with buried APBs; bi-domain GaP layer with emerging APBs on a n-doped Si substrate; bi-domain GaP layer with emerging APBs on a highly resistive Si substrate. The green color represents the APBs in the schematic diagrams of the different samples. The results demonstrate that the current flows vertically through the III-V APBs and laterally through the silicon.

136 Figure 5 . 7

 13657 Figure 5.7 Topography (a) and current distribution C-AFM images of the bi-domain GaP/Si sample under negative voltage (-8.5V) (b) and positive voltage (+9.0V) (c).

Figure 5 . 8 137 5. 5

 581375 Figure 5.8 Topography (a) and current distribution (b) C-AFM images (DC bias = -7.5V) of the bi-domain GaPSb/Si sample.

Figure 5 . 9

 59 Figure 5.9 Chopped-light linear sweep voltammetry curves (scan rate = 50 mV/s) of the GaP/Si (a) and GaP wafer (b).

Figure 5 .

 5 Figure 5.10 Chopped-light linear sweep voltammetry curves (scan rate = 50 mV/s) of GaPSb/Si (a) and GaPAs/Si (b) samples with emerging APBs.

Figure 5 . 11

 511 Figure 5.11 Voltammetry curve of GaPSb/Si:p sample with emerging APBs in 1.0M H2SO4 (pH=0.3) electrolyte under chopped simulated sunlight.

Figure 5 .

 5 Figure 5.12 IPCE spectra recorded at -0.74 V vs RHE (cathodic region (CR)-green and black lines) and 1.26 V vs RHE (anodic region (AR)-red and purple lines) for the GaP/Si (red and green lines) and GaP wafer (purple and black lines). The dotted lines correspond to integrated photocurrent and the inset are the enlarged IPCE spectra around 350-800nm.

Figure 5 . 13

 513 Figure 5.13 Nyquist plots (between 100 kHz and 1 Hz) of the GaP/Si sample in the dark and under illumination at different applied potentials (OCP point, anodic region (AR), cathodic region (CR)).

Figure 5 .

 5 Figure 5.14a shows the Mott-Schottky (MS) plots recorded at 500 Hz for the GaP wafer, which demonstrates the typical positive slope of the reciprocal of the square of capacitance versus the applied potential expected for an n-doped semiconductor. From the slope, the donor density is evaluated to 2x10 17 cm -3 , in good agreement with the GaP wafer specifications. The MS plot for the GaP/Si sample recorded at 1 kHz is displayed in Fig.5.14b, and reveals both the positive and negative slopes, echoing with the previously discussed ambipolar behavior and suggesting attractive simultaneous n-and p-type-like electronic nature of III-V/Si materials[START_REF] Meng | Solar Hydrogen Generation by Nanoscale p-n Junction of p-type Molybdenum Disulfide/n-type Nitrogen-Doped Reduced Graphene Oxide[END_REF][START_REF] Li | Photoelectrochemical performance enhanced by a nickel oxide-hematite p-n junction photoanode[END_REF][START_REF] Zhu | Synthesis of the 0D/3D CuO/ZnO Heterojunction with Enhanced Photocatalytic Activity[END_REF].

Figure 5 .

 5 Figure 5.14 Mott-Schottky plot for the GaP-wafer (a) and bi-domain GaP/Si sample (b) in the dark.

discussed in the Chapter 3 , 144 Figure 5 . 15

 3144515 Figure 5.15 Zinc-blende (a), III-III APB (b) and V-V APB (c) slab structures with Pmma space group for theoretical calculations.

Figure 5 .

 5 Figure 5.16 Longer III-III (a) and V-V (b) APB slab structures.

  reciprocal space correspond to the[START_REF] Kang | Electrochemical Synthesis of Photoelectrodes and Catalysts for Use in Solar Water Splitting[END_REF],, [010] crystallographic directions, based on Heyd-Scuseria-Ernzerhof (HSE) hybrid functional. The band structure plots of the III-III and V-V APB structures for GaP, GaAs and GaSb are shown in Fig.5.18b,c,e,f,h,iwhere the color scale indicates the spatial localization of the electronic state in the APB plane. Note that the common k-paths chosen in all the cases for the sake of comparison 145 correspond to APB in-plane directions (Fig.5.15b,c) to illustrate the effects of the 2D electronic dispersion. Figure 5.18a,d,g show the band structure of the respective reference III-V zinc-blende slab structures for GaP (Fig. 5.18a), GaAs (Fig. 5.18d) and GaSb (Fig. 5.18g), respectively. Compared to the normal bulk structures (unit cell with 2

Additionally, Fig. 5 .

 5 19a,b show the planar averages of the charge densities integrated parallel to the APB plane. The charge density increases on III-III APB and drops on V-V APB, revealing that more electrons (holes) are localized on the III-III (V-V) APB planes, and suggesting that the non-stoichiometric APB singularities behave as charged domain walls.

Figure 5 .

 5 Figure 5.17 DFT band structures based on HSE potential for zinc-blende unit cell of GaP (a), GaAs (b) and GaSb (c).

Figure 5 . 18

 518 Figure 5.18 Calculated electronic band structures of zinc-blende, III-III APB and V-V APB structures for GaP (a-c), GaAs (d-f) and GaSb (g-i). The APB band structures (b,c,e,f,h,i) are weighted by the spatial localization of the states at the APB interface and the color map from blue to red underlines the increase of localization effect of the states at the APB.

Figure 5 . 19

 519 Figure 5.19 Charge density of III-III APB (a) and V-V APB (b) structures, integrated in the planes parallel to APBs.

147

 147

  calculations at the HSE level to have a precise insight into the topological properties of APBs[START_REF] Franchini | Maximally localized Wannier functions in LaMnO3 within PBE+ U, hybrid functionals and partially self-consistent GW: an efficient route to construct ab initio tight-binding parameters for eg perovskites[END_REF]. The band structures along the X--X k-path of III-III and V-V APB structures for GaP, GaAs and GaSb were extracted, as shown in Fig.5.20, with a very good correspondence to the initial DFT calculations (Fig.5.18).

Figure 5 . 20

 520 Figure 5.20 Band structures associated to III-III and V-V APBs along X--X based on firstprinciple tight-binding model for GaP (a,b) GaAs (c,d) and GaSb (e,f).

Figure 5 . 21

 521 Figure 5.21 Total density states of the III-III and V-V APB slab structures for GaP (a,d), GaAs (b,e) and GaSb (c,f) based on first-principle tight-binding model.

Figure 5 .

 5 Figure 5.22 3D plot of CBM and VBM of the III-III and V-V APB slab structures, extracted based on first-principle tight-binding model (a,c,e,g,i,k) and realistic density of states in log scale for the III-III and V-V APB structures (b,d,f,h,j,l).

5 . 22 )Fig. 5 . 5 . 25 )

 5225525 Fig. 5.23 displays the selected k-paths in the kz=0 plane for electronic band calculations of the III-III APB structure centering around  and Y points. The corresponding band structures and enlarged band structures centering around  points are shown in Fig. 5.24 and Fig. 5.25, where the irreducible representations (IRs) of the CBM and VBM are indicated. We can observe the bands open gaps around 40-60 meV (Fig.5.25) due to the hybridization of the same irreducible representation of the CBM and VBM [136]. The band structures and enlarged band structures around crossing point centering around Y points are shown in Fig. 5.26 and Fig. 5.27, together with the irreducible representations (IRs) of the CBM and VBM around crossing points. The different irreducible representations (IRs) of the CBM and VBM avoid hybridization effects and protect the crossing points, which are evidenced by the enlarged band structures (Fig. 5.27). The crossing points are marked by green dots in the k-path illustration (Fig. 5.23), which form one pair of Weyl points and one nodal ring centering around Y point, showing good agreement with the 3D band plot based on first-principle tight-binding model (Fig. 5.22).

Figure 5 . 23

 523 Figure 5.23 Illustration of selected k-paths in the kz=0 plane for electronic band calculations based on III-III APB structure and the crossing points marked by green dots.

Figure 5 . 24

 524 Figure 5.24 Electronic bands with band symmetry information of III-III APB structure along kpath X--X, M--M, S--S and N--N.

Figure 5 . 25

 525 Figure 5.25 Enlarged band structures corresponding to Fig. 5.24.

Figure 5 . 26

 526 Figure 5.26 Electronic bands with band symmetry information of III-III APB structure along the different k-paths centering Y high symmetry k point.

Figure 5 . 27

 527 Figure 5.27 Enlarged band structures corresponding to Fig. 5.26.

Figure 5 . 28 154 Figure 5 . 29

 528154529 Figure 5.28 Illustration of selected k-paths in the kz=0 plane for electronic band calculations based on V-V APB structure and the crossing points marked by purple dots.

Figure 5 . 30 29 . 5 . 31 ,

 53029531 Figure 5.30 Enlarged band structures corresponding to Fig. 5.29.

Figure 5 . 31 156 5. 7

 5311567 Figure 5.31 Fermi energy alignments of III-V zinc-blende, III-III and V-V APB structures with n-doped and p-doped Si corresponding to GaP/Si (a), GaAs/Si (b) and GaSb/Si (c). All the energy levels are shifted with the energy of vacuum set to 0 eV.

Figure 5 . 32

 532 Figure 5.32 Electronic band structures of the stoichiometric APB structures of GaP (a) and GaSb (b). The white dotted lines corresponding to the CBM and VBM of the bulk supercell.

5. 8

 8 General description of III-V/Si photoelectrodes operation5.8.1 General picture for the ambipolar III-V/Si photoelectrodeTo describe the ambipolar PEC operation of the hybrid semimetal/semiconductor 157 bidomain III-V/Si material, from the carrier photo-generation to the interfacial Faradaic reactions (represented in this scheme as OER and HER, although photocorrosion certainly takes place), we propose the mechanism depicted in terms of band lineups for GaP (Fig.5.33a-c), together with its spatial representation (Fig.5.33d,e). As described above, the Fermi levels of V-V APBs and III-III APBs are respectively close to the CBM and VBM of the bulk (Fig.5.31a-c). When the III-III and V-V APBs couple together with bulk materials, charge redistribution occurs and align the Fermi levels to reach equilibrium state. Especially, V-V APB structures are expected to lose electrons, to lower the V-V Fermi level. Meanwhile, III-III APB structures on the contrary will tend to capture electrons, with an associated increase of their Fermi level. Indeed, even though the APBs introduce metallic states, the states are much less numerous than the bulk part as shown previously by the DOS calculation (Fig.5.22b,d,f,h,g,l), which indicates that the Fermi level of the APB structures can easily be shifted by accepting or donating electrons without significant change of the bulk III-V Fermi level. As a consequence, the bands of the bulk III-V will bend downward at the vicinity of V-V APB and bend upward at the vicinity of III-III APB (as shown in Fig.5.33d,e), and lead to more holes in V-V APB and more electrons in III-III APB structures and forming strong in plane built-in electric field with a direction from V-V APB pointing to III-III APB structures (white arrows in Fig.

Figure 5 . 33 :

 533 Figure 5.33: Ambipolar water splitting with bidomain III-V/Si materials. a-c, Fermi energy band diagram of III-V bulk, III-III APB, V-V APB and electrolyte for GaP in three cases: (a) under illumination, (b) under illumination and positive bias and (c) under illumination and negative bias. The yellow and green strips correspond to the occupied states of III-III and V-V APBs. d,e, Spatial schematic diagram of the ambipolar water splitting mechanism as a photocathode (d) and as a photoanode (e) based on V-V and III-III APB singularities, respectively. The yellow and green rectangles represent the III-III and V-V APB singularities, respectively, and the black lines correspond to the CBM and VBM of bulk III-V semiconductor. The black arrows represent the directions of carriers (electrons or holes) movement, and the red arrows represent the directions of current.

5 .9a and Fig. 5 . 10 )

 5510 photocurrent higher than the cathodic one (Fig.5.9a and Fig.5.10) while the opposite situation is encountered for the sample grown on p-doped Si with higher cathodic photocurrent (Fig.5.11). This indicates that the local band bendings of Si with III-III and V-V APBs at the III-V/Si interfaces play an important role in the process.

Figure 5 . 34

 534 Figure 5.34 Schematic diagram of bi-domain III-V/Si sample with III-III and V-V APBs. The

Figure 5 . 35 Fig. 5 .

 5355 Figure 5.35 Local band bendings between Si with III-III and V-V APBs of the GaP/Si:n samples in three conditions: without bias (a), under positive bias (b) and under negative bias (c), proposed to explain the higher anodic photocurrent and lower cathodic photocurrent as shown in the voltammetry curve (Fig. 5.9a). The small subgraphs with green and orange enclosed boxes correspond to the interfaces of Si with V-V APB and Si with III-III APB, respectively.

Figure 5 . 36

 536 Figure 5.36 Local band bendings between Si with III-III and V-V APBs of the GaPSb/Si:p sample, proposed to explain the higher cathodic photocurrent and lower anodic photocurrent as shown in the voltammetry curve (Fig. 5.11). The Fermi level of Si is either between the Fermi levels of III-III and V-V APBs and closer to the III-III's one (a) or even below both III-III and V-V Fermi energies (b). The small subgraphs with green and orange enclosed boxes correspond to the interfaces of Si with V-V APB and Si with III-III APB, respectively.

162 5. 8 . 3

 16283 photo-generated electrons close to the Si/V-V interface will easily access to the GaP layer through V-V APB leaving the holes in Si. The photo-generated holes close to the Si/III-III interface will easily flow to the GaP layer through III-III APB structure leaving the electrons in Si. This will lead to electrons/holes recombinations, as shown in Fig.5.37a.Therefore, the light absorption of n-doped Si (of the GaP/Si:n sample) doesn't lead to an obvious PEC response.

Figure 5 . 37

 537 Figure 5.37 Scheme explaining the light-absorption contribution of the different Si substrates to the PEC response for the III-V sample (a) with n-doped Si substrate (i.e. the Si Fermi level lies between the Fermi levels of III-III and V-V APB structures) and (b) with a p-doped Si substrate (where the Si Fermi level lies below both the III-III and V-V Fermi levels).

Figure 5 .

 5 Figure 5.38 IPCE spectrum recorded at -0.74 V (cathodic region) for the bi-domain GaPSb/Si:p sample, which shows a large signal between 600 and 1000 nm (corresponding to Si absorption).

38 .Figure 5 .

 385 Figure 5.39 a. General illustration of the bi-domain III-V/Si material composed of bulk photoactive semiconductors with 2D APB semi metallic inclusions, realizing ambipolar PEC water splitting. b. Schematic plot of carriers' transport along the 2D APB planes at the atomic scale.

Chapter 2 ,In Chapter 3 ,In Chapter 5 ,Finally, in Chapter 5 ,

 2355 the experimental and theoretical supports of this work have been presented. The related mechanisms, techniques and parameters settings of different structural, optical and photo-electro-chemical characterizations used in this thesis have been described. Besides, after a brief introduction to the basic concepts of the DFT, the calculation steps and parameters settings of DFT calculations for band structures, phonons and electron-phonon coupling have been presented. antiphase boundaries in InGaP were found to enhance In incorporation around the APBs in the sample, thus behaving as 2D In-rich vertical homovalent singularities. Based on Raman, PL experiments and DFT calculations, simultaneous confinement of charge carriers and phonons at the vicinity of the 2D homovalent singularity (In-rich stoichiometric APB) has been evidenced. This situation led to a strong enhancement of the electron-phonon coupling, which has been indicated by a large experimental Huang-Rhys factor and discussed in the framework of the optical deformation potential.In Chapter 4, a novel III-V/Si (GaPSb/Si) tandem materials association was investigated for solar water splitting applications. The band evolution of the GaPSb alloy has been determined over the whole Sb range by absorption measurements and tightbinding calculations. The indirect (X-band) to direct (Γ-band) cross-over was found to occur at 30% Sb content. Especially, at a Sb content of 32%, the GaP1-xSbx alloy reached the desired 1.7 eV direct bandgap, enabling efficient sunlight absorption. Besides, the 168 band alignment of GaP1-xSbx/Si with respect to water redox potential levels has been analyzed, which shows the 1.7eV/1.1eV GaPSb/Si association is an interesting combination both for the hydrogen and oxygen evolution reactions. Finally, a GaP0.67Sb0.33/Si photoanode with a very close to the optimal bandgap combination (1.7eV/1.1eV) was investigated for photoelectrochemical (PEC) characterizations and it shows promising PEC performances due to the high light absorption but the transport mechanism in these photoelectrodes is questioned. we proved that the two kinds of III-III and V-V non-stoichiometric APBs in the III-V layers are 2D semi-metallic inclusions by combining structural, transport characterizations with first-principle calculations. Besides, The III-III and V-V non-stoichiometric APBs have been evidenced to have quite high and low Fermi energy levels, respectively, which introduces a strong built-in electric field in nano-scale regions when they couple together with the bulk III-V. On this basis, we demonstrate ambipolar operation of various III-V/Si photoelectrodes (GaP/, GaAsP/ and GaPSb/ on silicon). The built-in electric field provides efficient driving force for efficient photogenerated charge carriers separation. Overall, we demonstrate that the hybrid bidomain III-V/Si materials are able within the same layer to absorb light efficiently, separate laterally the photogenerated carriers, transfer them to semimetal singularities, and extract both electrons and holes vertically, leading to efficient carrier collection and ambipolar performances.PerspectivesIn this thesis, we clarified for the first time the unusual physical properties of antiphase boundaries, which fully explain why APBs cannot be used in standard p-n architectures for photonics (e.g. lasers) or for energy harvesting (e.g. photovoltaic solar cells). But these new clarified physical properties also open many prospects to design novel devices for photonics, electronics, computing or energy harvesting. Even if a first device demonstration was performed with the development of operating photoanodes with these 2D nano-objects, many different devices architectures could be imagined from these findings. 169 More specifically, from the water splitting point of view, this work only gave the first demonstrations of operating photoanodes, and the PEC experiments were performed in acid and alkaline electrolytes with only two pH values. Systematic research of the influence of pH on band alignment between III-V and water levels could be done. The efficiency and spectral dependency of PEC performances should also be considered in the future. Besides, as demonstrated in Chapter 4 and 5, a careful optimization of Fermi levels lineups through both III-V alloys composition or Si substrate doping can optimize charge carriers transfers and activate the Si substrate PEC response. Therefore, more III-V/Si MBE samples with different composition or doping would surely be of great interest to reach high-efficiency water splitting. Finally, with these materials, a whole PEC cell combining a photo-anode and a photo-cathode should rapidly be developed, in order to demonstrate autonomous operation of the system. we demonstrated preliminary results on the topological properties of the non-stoichiometric APB structures. These investigations are of high interest, as it could lead to a new topological nano-material to be studied experimentally and theoretically in fundamental physics with expected non-trivial topology-induced transport properties.

  69at 550°C and 80 Torr reactor pressure on one clean Si (001) wafer misoriented by 4 o towards the[START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] direction with a reduced pressure chemical vapor deposition (RPCVD) system. Then SiGe was annealed in-situ at 1000°C for 5 mins under H2 flow for relaxation.The second step is the regrowth of In0.3GaP on SiGe/Si pseudo-substrate by our lab, which contains two growth modes: Migration Enhanced Epitaxy (MEE) and traditional Molecular Beam Epitaxy (MBE). The first 10nm InGaP layer was grown with the MEE growth mode. One MEE cycle consists in a first 0.63 ML (MonoLayer) Ga deposition at a growth rate of 0.1 ML/s, followed by a 0.27 ML In deposition at the same growth rate (0.1 ML/s), and then by a 4s growth interruption. Finally, a 200nm-thick In0.3GaP layer is grown at 500°C in the conventional continuous MBE growth mode and at a growth rate of 0.14 ML/s on the top of the MEE layer.

Table 3 . 1 .

 31 Space group IR and layer group IR of the five new phonon modes (all grouped by pairs) in order to evidence the phonon modes characteristic of a single APB.

	Phonon mode order	Frequency(THz)	Pnnm space group	p2/m11 layer
			IR (G)	group IR ()
	47 /48	5.269 / 5.337	Ag /B1g	Ag
	49 /50	5.891 /5.952	B2u/B3u	Bu
	51/52	7.836/7.837	B1u/Au	Au
	53/54	7.881/8.037	B2u/B3u	Bu
	95/96	12.610/12.616	B1g/Ag	Ag

Table 3 . 2

 32 Electron phonon interaction matrix element |𝑔 𝑚,𝑛,𝑣 (𝑘, 𝑔)| related to the Ag phonon non-polar modes for the bottom of the conduction band (m=n=CBM) and the top of the valence band (m=n=VBM) at the  point

		5.4THz Ag Localized	13.4THz Ag Localized
		phonon	phonon
	CBM, CBM	5.212669	4.219345
	VBM, VBM	2.327812	8.652394

Table 3 .

 3

2 lists the electron phonon interaction matrix elements |𝑔 𝑚𝑛𝑣 (𝑘, 𝑞)| related

  47 µm -1 , which can be roughly extracted from SEM images, i.e. around 1 APB each 100 nm of bulk. While, the linear density of APB of the supercell for calculation, taking III-III APB slab structure for example, is 0.5826 nm -1 i.e. around 1 APB each 1.7163 nm of bulk. Therefore, the APB part of the total density states of III-III APB supercell (Fig.5.21a) keep unchanged (i.e. is multiplied by 1) and the bulk part needs to be multiplied by 58.26 (100 nm/1.71631 nm)

  Two steps III-V on Si growth process: Silicon substrate degassing followed by MEE and MBE two growth modes[START_REF] Yanping | Structural analyses by advanced X-ray scattering on GaP layers epitaxially grown on silicon for integrated photonic applications[END_REF]. M.Alqahtani, C. Levallois, A. Létoublon, J. Le Pouliquen, R. Piron, S. Boyer-Richard, L. Pedesseau, T. Rohel, R. Bernard, Y. Léger, G. Loget, J. Wu, I. P. Parkin, B. Fabre, N. Bertru and C. Cornet. GaPSb/Si tandem material with APBs for efficient overall water splitting, Talk at journée nano-matériaux de Rennes, January 14, 2020, Rennes, France. -L. Chen, O. Skibitzki, Y. Léger, R. Piron, J. Stervinou, A. Letoublon, R. Tremblay, T. Rohel, R. Bernard, C. Levallois, L. Pedesseau, T. Schroeder, A. Moréac and C. Cornet. Strong electronphonon coupling in 2D-vertical In-rich APB structures, Poser at journée nano-matériaux de Rennes, January 14, 2020, Rennes, France. -L. Chen, R. Piron, O. Skibitzki, Y. Léger, C. Levallois, T. Rohel, L. Pédésseau, R. Bernard, J. Stervinou, A. Letoublon, T. Schroeder and C. Cornet. Electron-phonon interactions around antiphase boundaries in InGaP/SiGe/Si: structural and optical characterizations, Poster at 20th Anniversary of LPQM-Institut d'Alembert, ENS -Paris/Saclay, Cachan 2019, April 12, 2019, Paris, France. R. Piron, C. Levallois, Y. Léger, C. Cornet, Oral presentation for a video reunion about PHD works in front of members of three FOTON labs, June 16, 2019, Rennes. -L. Chen, R. Piron, C. Levallois, Y. Léger, C. Cornet, Oral presentation for a reunion about PL experiment results of III-V/Si materials, May 18, 2018, Rennes. -L. Chen, R. Piron, C. Levallois, C. Cornet. Oral presentation about PHD thesis in front of members of FOTON-OHM, February 22, 2018, Rennes.

e)-Communications with national audience -L. Chen, -L. Chen, M. Alqahtani, C. Levallois, A. Létoublon, J. Stervinou, R. Piron, S. Boyer-Richard, J.-M. Jancu, T. Rohel, R. Bernard, Y. Léger, L. Pédésseau, N. Bertru, J. Wu, I. P. Parkin and C. Cornet. GaPSb/Si electrodes for Solar Fuel Production, Poster at European COST action on Solar Energy, April 10-12, 2019, Bulgaria. -L. Chen, O. Skibitzki, Y. Léger, R. Piron, C. Levallois, A. Letoublon, J. Stervinou, R. Tremblay, T. Rohel, R. Bernard, T. Schroeder, G. Capellini and C. Cornet. study on structural and optical properties of InGap/SiGe/Si, Poster at 3M doctor day, October 29, 2018, Rennes, France. f)-Selected seminars -L. Chen, Y. Léger, L. Pedesseau, J. Even, C. Cornet, Oral presentation for a reunion about calculations about topological properties, September 18, 2020, Rennes. -L. Chen,

Acknowledgements

Firstly, I would like to deeply acknowledge my supervisor Charles Cornet. Thank you for giving me the chances to study these interesting projects and for your continuous support and patience during my PhD studies. I learn a lot from you, not only on the research works but also on daily life. I also thank Christophe Levallois and Rozenn Piron sincerely for your instructions and helps.

Besides, there are many other FOTON members give me great helps, instructions and suggestions. I express my sincere gratitude to all of you, especially Yoan Léger, Jacky Even, Laurent Pedesseau, Antoine Létoublon, Julie Le Pouliquen, Tony Rohel, Nicolas Bertru, and Soline Boyer. I thank all our collaborators, especially Gabriel Loget (ISCR), Rozenn Gautheron-Bernard (IPR), Oliver Skibitzki (IHP). I also thank the committee members Magali Benoit, Xavier Marie, Antonio Tejeda, and Vincent Artero for your good comments, suggestions and questions for further improving the thesis.

29166.

b)-Patents -L. Chen, S. Li, B. Fan et al. A new real-time controllable device and method for separation of micro-droplets based on lithium niobate wafer sandwich structure (CN105413767A). -L. Chen, S. Li, B. Fan et al. A new method for transporting micro-droplets based on Y-cut lithium niobate wafer (CN105510615A). c)-Invited talks -C. Cornet, L. Chen, A. Ruiz, I. Lucci, A. Beck, L. Pedesseau and Y. Léger. "Monolithic integration of III-V semiconductors on silicon for photonics and solar hydrogen production", Invited talk at SPb Photonic, Optoelectronic, & Electronic Materials (SPb-POEM), April 24-27, 2019, Saint-Petersburg, Russia. -C. Cornet, L. Chen, I. Lucci, A. Ruiz, A. Beck, L. Pedesseau, and Y. Léger. "GaP-based materials for photonics and energy", Invited talk at SNAIA 2018 -Smart Nanomaterials Conference, December 10-13, 2018, Ecole normale supérieure de Chimie-Paris. d)-Communications with international audience -L. Chen, O. Skibitzki, Y. Léger, C. Levallois, R. Piron, T. Rohel, J. Stervinou, A. Letoublon, T. Schroeder and C. Cornet. Photoluminescence of 2D-vertical In-rich APBs embedded in InGaP/SiGe/Si. Poster at 20th International Conference on Molecular-Beam Epitaxy (ICMBE 2018), September 02-07, 2018, Shanghai, China. -L. Chen, O. Skibitzki, Y. Léger, C. Levallois, R., T. Rohel, J. Le Pouliquen, A. Létoublon, T. Schröder and C. Cornet. Excitons bounded around In-rich antiphase boundaries. Poster at 34 th International Conference on the Physics of Semiconductors (ICPS 2018), July 29 to August 3, 2018, Montpellier, France. -L. Chen, O. Skibitzki, R. Piron, J. Stervinou, A. Létoublon, R. Tremblay, T. Rohel, R. Bernard, C. Levallois, Y. Léger, T. Schroeder, G. Capellini and C. Cornet. Antiphase boundaries in InGaP/SiGe/Si: structural and optical properties. Poster at European Materials Research Society-Spring Meeting 2018 (E-MRS 2018 Spring Meeting), June 18-22, 2018, Strasbourg, France.

79

Bandstructure calculations

Electronic bands were calculated along different k-paths (Z(0,0,0.5))-Г(0,0,0)-U(0.5,0,0.5); R(0.5,0.5,0.5)-Г(0,0,0)-S(0.5,0,0.5)), where the axis x, y, z of the reciprocal space correspond to the [001], [1][2][3][4][5][6][7][8][START_REF] Holt | Extended Defects in Semiconductors: Electronic Properties, Device Effects and Structures[END_REF][START_REF]Structures of Element and Compound Semiconductors[END_REF], [START_REF] Hankin | Flat band potential determination: avoiding the pitfalls[END_REF] crystallographic directions, based on the Heyd-Scuseria-Ernzerhof (HSE) hybrid functional. The band structure plot for the APBstructure is shown in Fig. 3.14, where the states are weighted by their spatial localization at the APB plane. The color map from blue to red underlines the increase of localization effect of the states considered around the APB. method needs very high vacuum environment, which can enable long mean-free path of atoms or molecules. Like this, the evaporated atoms can avoid interacting with each other or vacuum-chamber gases until they reach the substrate. This is why we call it "beam".

In order to monitor the growth in-situ, a Reflection High-Energy Electron Diffraction (RHEED) system consisting of a RHEED Gun and a fluorescent screen are equipped on both the chambers, operating at 30 keV for UHVCVD and 12 keV for MBE. A computer is used to control the temperature of the substrate and evaporation sources, allowing precise control of the thickness and composition of samples. Title : Optoelectronic, vibrational and transport properties of III-V/Si antiphase boundaries for photonics and solar energy harvesting Keywords : III-V/Si semiconductors, antiphase boundaries, 2D materials, vertical homovalent singularities, optoelectronic properties Abstract : This thesis aims to investigate the specific optoelectronic properties of III-V/Si Anti-Phase Boundaries (APBs) and its use for energy harvesting devices theoretically and experimentally. Strong electron-phonon coupling around stoichiometric APBs are first demonstrated due to simultaneous confinement of charge carriers and phonons in the same region, based on structural and optical characterizations and density functional theory calculations. The GaPSb/Si tandem materials association for solar water splitting is then studied. Combining ellipsometry measurements and tight binding calculations, the bandgap of GaPSb alloys in the whole Sb range and band lineups of GaPSb/Si with water redox levels are obtained, which shows the potential of the GaPSb/Si association for the hydrogen evolution and oxygen evolution reactions. Then a GaP0.67Sb0.33/Si photoanode with an almost optimal bandgap combination (1.7eV/1.1eV) was investigated for photoelectrochemical (PEC) water splitting with promising performances related to efficient sunlight spectrum absorption. Finally, the transport and PEC properties of III-V/Si with vertical non-stoichiometric APBs are investigated from experimental characterizations and first-principle calculations. We demonstrate that epitaxial III-V/Si materials with vertical non-stoichiometric APBs are hybrid structures, composed of bulk photo-active semiconductors with 2D topological semi metallic vertical inclusions, enabling simultaneously good photo-activity, efficient charge transport and separation, and interesting ambipolar properties.