Keywords: Hyperspectral cathodoluminescence maps, ALD Atomic-Layer Deposition, ARC Anti-reflection coating, BSF Back surface field EBL Electron beam lithography, EDS Energy dispersive X-ray spectroscopy PCE Power conversion efficiency, PECVD Plasma-enhanced chemical vapor deposition QFLS Quasi-fermi level splitting SEM Scanning-electron microscope, SMU Source monitoring unit, STEM Scanning transmission electron microscope, TCO Transparent conductive oxide

Investigation of polycrystalline CdSeTe layers by cathodoluminescence

4.1

Titre : Avancées sur les cellules solaires ultrafines en GaAs et en CdTe Mots clés : Photovoltaïque, CdTe, GaAs, piégeage optique, cathodoluminescence Résumé : Les cellules solaires en couche mince, constituées de semiconducteurs à bande interdite directe d'une épaisseur de 2 à 4 µm, constituent une alternative attractive aux cellules conventionnelles formées de 150 µm de silicium. Elles permettent une faible consommation de matériaux et une empreinte carbone réduite. Avec une réduction supplémentaire ultrafines offrent de nouvelles possibilités pour économiser des matériaux rares, réduire le temps et le coût de fabrication, faciliter la collecte des porteurs de charges dans les absorbeurs, et étendre le champ des applications du photovoltaïque à des dispositifs flexibles ou tolérants aux rayonnements. Cependant, des stratégies de piégeage optique sont nécessaires pour compenser la faible absorption des films ultrafins, et elles doivent respecter des contraintes de conception et de fabrication exigeantes.

Cette thèse se concentre sur le développement de cellules solaires ultrafines en GaAs, un semiconducteur monocristallin de haute qualité qui a permis le record d'efficacité des cellules solaires à simple jonction, et en CdTe, un matériau polycristallin peu coûteux et adapté à la fabrication industrielle à grande échelle.

Dans une première partie, les pertes des cellules solaires précédentes en GaAs de 200 nm d'épaisseur sont analysées, et un nouveau design de piégeage nanostructurée à fort indice, qui est encapsulée dans un matériau mésoporeux à faible indice, avec un miroir plan en argent en face arrière.

Cette structure permet d'obtenir un courant de court-circuit théorique de 28,2 mA/cm². Une preuve de concept de cette cellule solaire ultrafine est démontrée, avec une réponse optique multirésonante en accord avec les simulations numériques. Afin d'améliorer encore l'efficacité globale des cellules solaires ultrafines, un nouvel empilement de couches basé sur une hétérostructure n-GaAs/p-AlGaAs de 250 nm d'épaisseur est introduit, le courant de court-circuit théorique atteint 28,9 mA/cm².

Une seconde partie est consacrée aux différents facteurs entravant le développement de cellules solaires CdSeTe/CdTe polycristallines ultrafines. La passivation des défauts par l'incorporation de sélénium est étudiée par cartographie de cathodoluminescence à haute résolution. La diffusion des atomes de sélénium et leur impact sur l'efficacité radiative des alliages CdSeTe sont analysés quantitativement. Ensuite, plusieurs solutions pour un contact arrière ohmique, passivant et réfléchissant sur le CdTe sont explorées. Enfin, la stratégie de piégeage optique multi-résonante développée pour le GaAs ultrafin est adaptée aux contraintes technologiques des dispositifs en CdTe. Par simulation optique, une efficacité de conversion allant jusqu'à 19,5% est prédite dans une cellule solaire en CdSeTe de 200 nm d'épaisseur avec un contact arrière nanostructuré.

Title : Advances for ultrathin GaAs and CdTe solar cells Keywords : Photovoltaics, CdTe, GaAs, light trapping, cathodoluminescence Abstract : Besides conventional 150 µm-thick silicon photovoltaics, thin-film solar cells made of 2-4 µmthick direct-bandgap semiconductors represent an attractive alternative with a low material usage and carbon footprint. With an additional 10-times thickness reduction, ultrathin solar cells offer opportunities to save scarce materials, fabrication time and cost, to ease the carrier collection in defective absorbers, and to extend the scope of PV applications with flexible or radiation-tolerant devices. However, light-trapping strategies are required to compensate for the low absorption of ultrathin films, and they should comply with demanding design and fabrication constraints.

This thesis focuses on the development of ultrathin solar cells made of GaAs, a high-quality monocrystalline semiconductor that holds the record efficiency of single-junction solar cells, and CdTe, a low-cost polycrystalline material suitable for largescale, industrial manufacturing.

First, the losses of previous 200 nm-thick GaAs solar cells are analyzed, and a novel light-trapping design is proposed. It is based on a nanostructured high-index layer embedded in a low-index mesoporous material and a flat silver mirror, resulting in a theoretical short-circuit current of 28.2 mA/cm².

A proof-of-concept device is successfully fabricated. It exhibits a multi-resonant optical response in agreement with numerical simulations. To further improve the overall efficiency of ultrathin solar cells, a new design based on a 250 nm-thick n-GaAs/p-AlGaAs heterostructure is introduced, with a theoretical short-circuit current of up to 28.9 mA/cm².

Secondly, the different factors that hinder the development of ultrathin, polycrystalline CdSeTe/CdTe solar cells are investigated. The passivation of defects by selenium incorporation is studied by high-resolution cathodoluminescence mapping. The diffusion of selenium atoms and their impact on the radiative efficiency of CdSeTe alloys are determined quantitatively. Then, several solutions for ohmic, passivating, and reflective back contacts on CdTe are explored. Finally, the multi-resonant light-trapping strategy developed for ultrathin GaAs is adapted to the technological constraints of CdTe devices. By optical simulation, a conversion efficiency of up to 19.5% is predicted in a 200 nm-thick CdSeTe solar cell with a nanostructured back contact.

A mes adorables neveux et nièce Jacob, Clara et Côme pour vous assurer que l'enjeu écologique est au coeur de nombre d'humains brillants et que je parie sans hésitation sur eux pour vous offrir le meilleur des avenirs.

"The ship is anchor'd safe and sound, its voyage closed and done;

From fearful trip the victor ship comes in with object won" Walt Whitman, 1865

Que la présence systématique des remerciements ne vous trompe, loin des usages de routine et d'automatisme ils sont à mon sens le reflet d'un véritable témoignage de reconnaissance. Pour vos contributions directes ou indirectes je souhaite pleinement vous rendre crédit. L'humble essai de parachever ce manuscrit s'initie donc par ces remerciements, qui ne seront je le crains qu'une maigre esquisse de ce que vous m'avez apporté durant ces trois années et demi de thèse et de la reconnaissance que je vous porte. A l'image de ce manuscrit, commençons par la fin, naviguons dans le début et finissons par l'entre-deux.

Un merci tout particulier aux membres de mon jury, d'avoir accepté d'en faire partie. Merci aux rapporteurs Régis A. et Gwénolé J. d'avoir pris le temps de lire patiemment le contenu de ce manuscrit et de vos retours, merci Jean-Paul K., Inès M. et Negar N. pour vos questions, votre intérêt et pour votre vision enrichissante sur le sujet.

Merci Stéphane et Andrea pour votre complémentarité si parfaite sur ce sujet (mais aussi sur tous les autres !). Merci Andrea pour tes connaissances salle blanche impressionnantes. Merci pour toute ton aide pour les expériences avec une multitude de tests ingénieux dont ce manuscrit peine à rendre compte. Merci pour ta gentillesse à Philadelphie. Merci Stéphane pour ton accompagnement de tout instant, pour ta rigueur scientifique et administrative, et pour l'intérêt non feint et profond que tu portes à nombre de projets. Merci pour l'enthousiasme et les encouragements, clés pour surpasser les moments plus difficiles. Merci à tous les deux pour l'implication et les yeux pétillants quand le sujet vous passionne, pour vos connaissances scientifiques infinies et le temps consacré à répondre aux questions sans jugement. Merci pour les opportunités de présentation de mon sujet et merci pour les challenges. Et aussi, surtout, merci pour vos qualités humaines, votre gentillesse, patience et bienveillance qui vous font passer haut la main dans la catégorie d'encadrant d'exception.

Merci Amaury D. pour ta disponibilité, toujours prêt à accorder quelques minutes (et plus) de ton temps qui résulte en une aide si précieuse et des conseils de grande qualité. Merci pour ta bonne humeur de chaque instant (et puis pour avoir manqué Roland Garros aussi).

Merci à mes tuteurs des enseignements de physique à la faculté de pharmacie de Châtenay-Malabry. Merci Florence d'avoir accepté ma participation et merci Nicolas d'avoir eu la patience de me montrer une infime partie de ton savoir sur la physique des fluides. Ta motivation contagieuse m'aura permis d'apprécier cette discipline toute nouvelle. Merci Sandrine et François-Xavier, Ghozlene et Bertrand pour les discussions et votre accueil. Merci à cette équipe pédagogique incroyable dont l'amour de leur travail et la bienveillance envers leurs élèves ne peut être autre chose qu'un modèle. Merci de m'avoir montré ces qualités, trop méconnues lorsqu'on est encore étudiant.

Et puis il y a l'univers de la salle blanche, complètement nouveau pour moi et je l'admets un peu (très ?) stressant et je voudrais en tout premier remercier deux personnes qui ont rendu ce monde beaucoup plus accueillant. Merci Abdou. Merci pour tous les conseils litho, pour ta gentillesse et ta disponibilité, pour les discussions en salle qui rendent les journées meilleures encore plus quand les manips ne fonctionnent pas. La dette de kebab et de pizza n'a fait que s'allonger au fil des années ! (très hâte d'assister à ta thèse). Et merci Christophe. Merci pour l'accompagnement des premiers pas en salle blanche. Merci de m'avoir appris toutes les bases nécessaires à l'apprentie que j'étais et que je reste, et de m'avoir transmis une partie de ton savoir intarissable sur la techno III-V et la litho. Merci pour la multitude d'anecdotes, le naturel de tes échanges et ton expertise MEB, merci d'avoir tellement contribué à ma prise de confiance dans cet univers. Il est temps d'accepter que je ne pourrais surement jamais vous rendre la pareille.

Merci Nathalie B. (pour les conseils litho et la douceur des échanges), Laurent, Laetitia (pour les formations, les rescousses en dépôt métallique et votre expertise), Alan (pour les multiples dépôts ALD), David (pour les découpes à la scie si précise et ta bonne humeur), Xavier (pour les recuits RTA), Nathalie I. (pour le micro-bonding), Jean-Luc (pour la gravure sèche RIE) d'avoir été des interlocuteurs de choix. Merci Téo de m'avoir formé patiemment sur la 550S (à qui j'ai fini par faire des infidélités) et merci pour les discussions techniques ou non qui rendent le temps de pompage toujours très productif et qui j'espère continueront. Merci Alain pour ta rapidité quel que soit le problème informatique. Merci Ludovic pour les images STEM d'exception et la facilité des échanges si agréable. Merci Ali pour les brèves discussions dans le sas de la salle blanche, rayon de soleil dans les journées. Merci Julie et Régis (IPVF/EDF) pour votre savoir-faire qui m'a dépanné plus d'une fois au C2N.

Quand on commence une thèse on se dit qu'on aura des collègues mais j'ai surtout (uniquement ?) trouvé des amis. Merci à tous d'avoir rendu ces années plus belles par votre présence, vos sourires et votre soutien. Merci Capucine, merci Salim (aka la portée 2019 de Stéphane). Capucine c'est à toi principalement que ces remerciements ne rendent pas honneur merci d'avoir été là, co thésarde d'exception, duo indissociable de Tic et Tac. Merci à vous deux, merci d'avoir fait partie de ce trio qui a commencé presque en même temps, d'avoir supporté mes râlages approchant le niveau olympique, merci pour cette semaine de rédaction à la montagne qui n'aura jamais été si productive. Les crédits de cette thèse vous reviennent en grande partie, merci d'avoir été ma belle rencontre, merci pour tous les bons souvenirs et les suivants.

Merci Linh pour ta présence discrète mais si nécessaire, merci pour le soutien constant dans les moments difficiles en salle blanche et les paroles réconfortantes.

Merci aux autres photovoltaïcien•nes rencontrés au fil des conférences et école d'hiver, Guillaume C, Paul O, Jérémy. Guillaume V, Marie merci d'avoir été les compagnons de fortune des Houches, du weekend suivant et des verres d'un soir qui restent de très beaux souvenirs.

Une dédicace spéciale aux plus anciens (ou presque). Merci Romaric, Olivier, Louis pour votre accueil chaleureux, mention toute particulière à Anatole dont les longues discussions me manquent. Merci Thomas d'avoir été le mentor de la CL et bien plus encore, merci pour toutes les bases si précieuses. Merci Tiphaine d'avoir brisé la glace entre sunlit et odin, que Berlin ne soit qu'une étape dans la suite de notre parcours ensemble, merci pour ton amitié inestimable (et pour le chapitre 2).

Merci à la SUNLIT team: Jeronimo, Sylvain, Maxime L., Oleksandr, Alexandre, Raphaël, Chiara, Emilie, Agathe et Alexis pour les rencontres du lundi matin et les conseils avisés. Merci Joya pour ta gentillesse sans égale et pour le guidage dans les méandres de reticolo. Merci Maxime G. pour le temps passé sur OptiPV à m'expliquer les bases de l'électromagnétisme avec sympathie. Merci Daniel de m'avoir rejoint sur les cellules ultrafines et ses péripéties ainsi que ta relecture de chapitre et merci d'être venu à la soutenance. Carlos tu t'approches du but, courage pour la rédaction mais je suis sûre que tu t'en sortiras très bien comme pour toute ta thèse. Naomie ta présence lumineuse est un plus pour cette équipe je te souhaite le meilleur pour la suite. Oleh bienvenue ! vi Merci aux ODINers et coéquipiers du C2N dont mon plus grand regret est de ne pas avoir sympathisé plus tôt avec vous. Merci Marina, Eduardo, Romain, Cléophanie, Nghia, Amir pour les trajets RER B allégés à tes côtés. Merci Amaury C. pour la culture chorale, Jules pour être un peu trop sur la même longueur d'onde au travail (et pour Zelda aussi). Bon courage à vous deux pour la dernière ligne droite. Merci Paul G. pour tes traits d'esprit subtils. Merci Stefano pour toute la suite sur la CL, le support technique de qualité, pour ton ingéniosité sans égale et pour la chicorée surtout.

Merci Mathieu d'être le co-bureau dont tout le monde rêve.

Je mesure pleinement la chance d'avoir pu travailler à vos côtés ces trois dernières années (bon ok un peu plus), merci pour vos sourires.

Les événements auront fait qu'il n'a pas toujours été aisé de vous voir ces dernières années mais qui n'enlèvent en rien à votre présence délicate et votre amitié qui m'est si précieuse: merci à mes amis d'enfance, merci à mes amis de Strasbourg et à mes comparses de photonique (merci pour votre présence à distance sur zoom, j'étais très touchée). Merci Arnaud pour les bouffées d'oxygène alcoolisé si peu régulières. Merci Adrien d'avoir gardé depuis le Canada un regard bienveillant sur nos travaux respectifs. Merci Darius, co-vacataire, tes mots d'encouragements ont résonné jusqu'à la toute fin et la compréhension mutuelle de nos vies de thésards en a soulagé la charge.

Et puis merci les copains d'avoir été là aux différents stades de la thèse. Solenne, Charlotte et Valentin, merci d'être les loco irremplaçables qui illuminent ma vie, qui la rendent plus drôle, plus belle et plus facile. Merci pour les week-ends de pause, les jeux et les vacances, merci pour les fous rires, les galères, les découvertes et les tristesses. Merci d'aborder la vie et sa réalité à mes côtés.

Merci à ma famille. Merci à mes grands-parents, Marie-Françoise qui par l'injonction de devoir copier mille lignes m'a aidé à continuer la rédaction et dont les discussions littéraires remplies de poésie continuent d'être une bulle d'oxygène et d'hésitation et Jean-Marie dont l'intelligence et l'intuition des choses m'encouragent dans la quête du savoir. Merci à mes parents, prendre tous vos défauts n'aura finalement pas si mal tourné (lol comme disent les jeunes). Merci pour votre amour des sciences et de votre travail qui a déteint (un peu) sur moi. Merci maman pour le soutien sans failles et la relecture de tous les chapitres, merci papa pour la touche discrète qui sublime le résultat, merci pour la canalisation des mauvaises ondes, merci pour votre accompagnement qui a fait que je ne me suis heurtée à aucun mur autre que ceux que je dresse moi-même. Je vous suis si reconnaissante.

Merci à mes deux soeurs adorées Eléonore et Mélisande dont les parcours exceptionnels et les philosophies de vie restent une source quotidienne d'inspiration. Votre présence est indissociable de mon existence, merci pour votre amour.

Merci Benjamin d'être l'élément stable de ma vie, merci d'être si différent et de me rappeler chaque jour la belle diversité du monde qui nous entoure. Merci pour la tolérance, la bienveillance, la tendresse, la confiance et l'amour. Merci d'avoir grandi avec moi. Merci de vivre dans l'oeil du cyclone.

vii viii

General introduction

In its sixth assessment report released in 2022, the Intergovernmental Panel on Climate Change (IPCC) confirmed the urgency of drastically reducing our CO 2 emissions to fight against global warming. To respect the Paris agreement and limit the warming to 1.5°C at the end of the century, carbon neutrality should be reached by 2050. Net zero CO 2 emission scenarios developed by the International Energy Agency (IEA) and IPCC remind us the importance of transforming our energy system (energy related emissions account for three-quarters of global CO 2 emissions) and the necessity to replace fossil fuels by low-to-no-carbon energies. In 2020, renewable energies accounted for 12.6% of the total final energy consumption (see Figure 1). Despite the progress of renewable in the power sector, the share of fossil fuels barely changed, in particular to meet the surge in the final energy demand that never ceases to grow. The global electricity demand increased by 4% in 2021, and is expected to reach 60,000 TWh by 2050. To achieve the reduction of CO 2 emission, both sobriety and the increase of the renewable energy are necessary. In this context solar energy, one of the world's most abundant resources, is arising as an essential alternative.

Over the last decade, the share of solar energy has become the lowest-cost option for new electricity generation in most of the world. Silicon technology accounts for 95% of the PV market for the electricity production with its low-cost commercial photovoltaic modules with efficiencies above 20%. However, due to its intrinsic properties, silicon power conversion efficiency is limited to less than 30% and still requires the use of a large amount of material with solar cell thicknesses of 150-200 µm and a considerable thermal budget for the fabrication of high-purity single-crystal silicon wafers. The PV industry is looking for alternatives to further increase the efficiency and reduce the production cost of solar cells. This leaves room for the thin-film technology based on direct bandgap materials: CdTe, CIGS, III-V (materials composed of one or more elements of column III and V from the periodic table) with high absorption. With absorbers thickness of 3-4 µm the material consumption is drastically reduced, and thin-films are very attractive for both light-weight/flexible applications and for combination with other materials to form multi-junction solar cells.

GaAs is at the forefront of high efficiency PV, holding the efficiency record of single-junction solar cell (29.1%). However, the cost of III-V raw materials still significantly limits the III-V technology industrialization, and restricts its field of application to space-PV. One way of improvements is to reduce the absorber thickness. Today, even state-of-the-art solar cells operate close to the single-pass absorption and efficient light trapping could reduce the solar cells thicknesses of at least one order of magnitude while maintaining equivalent efficiencies, with great advantages in terms of material consumption and industrial fabrication throughput. Novel light-trapping structures are required to compensate for the low single-pass absorption and to maintain a high photocurrent generation. Recent developments in nanophotonics provide new tools to go in this direction. In 2019, a record 19.9%-efficient GaAs solar cell with an absorber thickness of only 200 nm has been demonstrated by our team using multi-resonant absorption [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]. Based on an efficient nanostructured back mirror it was fabricated using low-cost and large-area nanofabrication techniques. Short-term perspectives predict the ability to achieve efficiencies greater than 25%.

It would be of particular interest to transfer this proof of concept to the commercial CdTe or the CIGS technologies, as they both represent most of the remaining 5% of the photovoltaic module production capacity in 2021 (about 4% and 1% for CdTe and CIGS, respectively). The energy-payback time (EPBT) of a PV module is the amount of time required to generate the energy necessary to recover the energy it took to produce the module. For silicon-based modules fabricated and operated in Europe, the EPBT is about 1 year [Fraunhofer 2023] while for thin-film technologies the EPBT is reduced to about 6 months [Attendees of the IW-CIGS Tech9 conference 2019]. The industrial competitiveness of these polycrystalline materials would be stimulated by reducing the thickness of the absorber while maintaining a similar efficiency. This would lower the consumption of scarce materials such as tellurium or indium in adequacy with a sustainable exploitation of the resources, shorten the deposition time, and facilitate the carriers collections with less bulk recombination. State-of-the-art performances with 15%efficient 500 nm-thick CIGS solar cells was achieved in the team [START_REF] Gouillart | Interface engineering of ultrathin Cu(In,Ga)Se2 solar cells on reflective back contacts[END_REF]]. In the case of the CdTe semiconductor, this important thickness reduction has been much less studied in the PV community.

However, CdTe-based solar cells are still far below their theoretical limit (32%) with record devices of 22.1% efficiency. Several challenges arise with a power conversion efficiency severely hampered by both low open-circuit voltage and external radiative efficiency. This is partly justified by their complex polycrystalline nature and the challenge of achieving a proper back contact. In this thesis, the strategy for ultrathin GaAs solar cells is applied as a model to CdTe, but a more general approach is performed in order to investigate the different factors that would prevent high-efficiency ultrathin CdTe devices.

In this thesis, I report different strategies to further optimize, both optically and electrically, the design and the fabrication of ultrathin solar cells made of GaAs and CdTe. The manuscript is structured in six chapters.

Chapter 1 introduces the basic principles of photovoltaic energy conversion and the main characterization techniques used to evaluate the performance of the cells presented in this work. A brief overview of the existing PV technologies is proposed. A greater emphasize is given on the GaAs and the CdTe materials to highlight their potential and the motivation behind the commitment to further reduce the thickness of the absorber.

In Chapter 2, I present the theoretical investigations conducted on ultrathin GaAs solar cells to further optimize their performance through optical simulations. First, the light-matter interaction is described, then an overview of the strategies proposed in the literature to trap light within ultrathin solar cells is given, with an emphasis on the GaAs technology. The development and the resulting state of the art performances obtained with the original ultrathin GaAs record cell developed in our team are reported. Based on an optical loss analysis, I propose a different light trapping strategy with multi-resonant absorption using an improved back side architecture which minimizes the parasitic absorption. Finally, I also propose to modify the stack structure with a 250 nm-thick heterojunction and present the results with an optimized light trapping scheme.

Chapter 3 presents the first results and the fabrication of the second generation of ultrathin GaAs solar cells developed in our group. It is based on the novel back architecture composed of a-Si:H/SiO 2 grating and followed by a flat silver mirror, proposed in Chapter 2. With a simple analytical model of resistive losses, I identify the optimal front grid design for an ultrathin solar cell and experimentally assess the validity of these calculations.

The second goal of this thesis work was to apply similar light-trapping strategies to CdTebased devices through theoretical investigation and experimental work. This aims at demonstrating the universality of the optical light-trapping method based on nanostructured back mirrors regardless of the materials for the next generation of high-efficiency and low-cost photovoltaics. Due to the intrinsic properties of the CdTe that hinder the realization of ultrathin devices, a part of this PhD work was to identify, characterize and attempt to improve the key factors necessary to overcome this efficiency limitation for ultrathin solar cells.

Chapter 4 is dedicated to the characterization of CdSeTe/CdTe thin-films by high-resolution cathodoluminescence (CL). First, the CL setup and its main features are described. Then a systematic study on five different CdSeTe/CdTe samples, with various selenium concentration, is performed to assess quantitatively the positive impact and the passivation effect of the selenium on the samples.

Chapter 5 reports the experimental work conducted on the back contacts of standard and thin CdSeTe/CdTe samples. The voltage deficit issue of CdTe-based solar cells is highlighted, and solutions proposed by the scientific community are discussed. Experimental results using novel back contacts compatible with light trapping strategies are presented. These back contact attempts are transferred to thinner CdSeTe/CdTe samples to evaluate their response. Based on experimental measurements, an optical model of the stack is proposed. Chapter 6 presents a numerical study of light trapping of ultrathin CdTe solar cells. I first report the still marginal investigation of the ultrahin CdTe devices in the scientific community. Then I propose several nanostructured back side architectures to increase the photogenerated current, taking into account the high requirements properties of CdTe back contacts. Finally, the experimental perspectives are discussed.

Contributions

This PhD thesis is the result of a team work with many collaborators and colleagues from the C2N, and here, I would like to detail their contributions, as well as mine, to the results presented in this manuscript.

I have performed the optical simulations of Chapter 2 and Chapter 6 with the Reticolo code supplied by Philippe Lalanne from Laboratoire Photonique, Numérique et Nanosciences and Christophe Sauvan from Institut d'Optique Graduate School. Daniel Micha and Maxime Giteau have implemented over the years an efficient and user-friendly code to facilitate these optical simulations.

Regarding the fabrication of ultrathin solar GaAs solar cells presented in Chapter 3, the epitaxy of the GaAs layers has been done at the Fraunhofer Institute for Solar Energy Systems. I performed every step of the solar cells transferred on a planar mirror. For the nanostructured cell, the amorphous silicon layer was deposited by Pere Roca i Cabarrocas from the Laboratoire de Physique des Interfaces et Couches Minces. Andrea Cattoni from C2N, performed the nanopatterning using e-beam lithography, and the necessary process optimization prior to the introduction on the ultrathin GaAs sample. The lithography step to open the localized back contacts was performed with the help of Christophe Dupuis.

For the fabrication of the CdSeTe/CdTe solar cells reported in Chapter 5, the growth of the layer has been conducted at the National Renewable Energy Laboratory (NREL). I conducted nearly every step of the solar cells process. However, the alumina deposited by ALD was conducted by Alan Durnez from C2N. The Cu 2 O layer was deposited by Abderrahime Sekkat from Laboratoire des Matériaux et du Génie Physique under the supervision of David Munoz-Rojas. The ITO was deposited by Andrea Cattoni.

I performed the current-voltage and EQE characterization of each cell presented in this work, except for the EQE measurement in Chapter 3 of the nanostructured GaAs solar cell that was performed by Jeronimo Buencuerpo from Institut Photovoltaique d'Ile de France (IPVF). The photoluminescence measurements was conducted with the support of Amaury Delamarre.

I performed all the CL measurements on CdSeTe/CdTe grown by the NREL presented in Chapter 4, with the help of Thomas Bidaud, a former PhD student in the laboratory, who trained me on the set-up and showed me how to analyze the resulting data. Part of the analysis presented in this chapter was conducted by both of us. The last section of this chapter shows CL measurements conducted with Stefano Pirotta from C2N, and he performed the data treatment. Finally the SEM-EDS measurement was conducted by Alexandre Blaizot from IPVF and the STEM-EDS measurements was performed by Ludovic Largeau from C2N. Solar energy has the potential to become one of the major source of electricity in the near future, necessary for the climate transition process. This Chapter introduces the basics of solar cells and reviews the characterization techniques used in this work to evaluate the properties of solar cells. An overview of the existing PV technologies is given with emphasis on the GaAs and the CdTe materials. The motivation of the thesis is presented.

Solar energy as a source of electric power

With the ability to convert the sun into electricity, and an unlimited supply source the photovoltaic technology is a highly attractive source of energy. The solar resource is such that it has been projected that six giant solar plants (with a conversion efficiency of only 8%) covering a total area of 0.1% of the Earth's surface would capture enough sunlight to reach 18 terawatts, which is the order of magnitude of the current total energy need of the world [Matthias 2006]. Photovoltaic is the third most important source of renewable electricity after hydroelectricity and wind power. With a steady growth of about 25% per year since 2019, the cumulative global solar PV capacity has reached 942 GW in 2021 (see Figure 1.1(a)). With scenarios predicting that the share of PV in the energy mix should be more than 30% by 2050, there is an opportunity for the PV technology to grow quickly. The photovoltaic industry is evolving rapidly, and the decrease of the module price is related to the increase in energy production. In fact, as shown in Figure 1. 1(b), every time the cumulative PV module production has doubled in the past 41 years, the price has decreased by about 25%. In addition, the levelized cost of electricity (LCOE), defined as the average cost per kWh produced over the lifetime of a PV system, is now below the cost of conventional fossil fuels.

Key facts on solar energy

-The carbon footprint of the PV technology depends only on the energy required for its fabrication and installation, it has a short payback time of only 1.2 years in Europe (much shorter than the 25 years of guaranteed lifetime) [START_REF] Fthenakis | Updated sustainability status of crystalline siliconbased photovoltaic systems: Life-cycle energy and environmental impact reduction trends[END_REF].

-France has the highest economic potential for rooftop PV production in Europe. To produce 125 TWh/year of PV energy, an area of 625 km 2 is required (with an average irradiation of 1 MWh/m 2 /year and an efficiency of 20%), which is about half of the estimated 1346 km

Basics of solar cells 1.2.1 Operating principle

The fundamental elements needed to fabricate a solar cell are introduced in this section. Based on the photovoltaic effect, the basic principle of solar cells is to convert light into electricity. This effect can be separated in three main steps: the light emitted by the sun is absorbed, it creates carriers which are spatially separated and then selectively collected by an external electric circuit (see Figure 1.2).

The photogeneration of carriers

To achieve this photovoltaic energy conversion, most of the solar cells are composed of semiconductor materials. Due to their electronic band structure, the electron energies are constraint within bands. The lower energy level is the valence band (E v ), where electrons are bounded, and the conduction band (E c ) is the energy level where electrons are considered as free charge carriers. The valence band and the conduction band are separated by the bandgap E g .

Impinging photons at the surface of the solar cell can be reflected at the surface depending on the refractive index of the semiconductor, transmitted through the stack or absorbed. Due to the two-bands system, if the photon energy (hν) is lower than the bandgap E g the photon passes through the material. On the other hand, if the incoming photon energy is equal or higher (i.e. shorter wavelength) than E g , it can be absorbed. Thus, the photogeneration will depend on the incident photon flux but also on the semiconductor absorption. The absorption coefficient (α) which is an intrinsic property of the semiconductor determines the depth at which the light of a particular wavelength can penetrate before it is absorbed. This assumption plays a major role in the material thickness of the solar cell, and will be further detailed in Chapter 2.

Once the photon is absorbed, its energy is transferred to the material and an electron from the valence band is excited to the conduction band resulting in an electron/hole pair (step (a) in Figure 1.2). This electron/hole pair loses all excess energy and thermalizes towards the band edges (step (b) in Figure 1.2).

Charges separation

Once generated, the carriers need to be separated. For a semiconductor, the p-n junction is generally used to spatially separate the electrons and the holes. A n-type (p-type) semiconductor is obtained by introducing donor atoms (acceptor atoms), which increase the electrons (holes) density. The diffusion and drift processes result in a built-in electric field which generates a depletion region with no carriers in it. The photogenerated carriers are separated by the internal electric field existing at the p-n junction, the electrons are swept to the n-type side, while the holes to the p-type side.

Carriers collection

The carriers collection is ensured by the charge transport towards the electrical contacts. The selective contacts are at opposite sides of the device (step (c) in Figure 1.2). The efficiency of the collection process will depend on the diffusion length that should be long enough to reach the depletion region and contribute to the current. Besides, recombination losses (detailed in Section 1.2.2) will limit the collection probability.

Characterization of solar cells

The structural, optical and electrical complexity of photovoltaic devices can be characterized by multiple measurements with the goal of understanding and further improving the efficiency. This section highlights some techniques implemented to evaluate the performance of solar cells.

Electrical characterization

Current density vs Voltage

This is one of the most fundamental PV measurements, referred to as current-voltage acquisition (IV curves). Figure 1.3(a) shows the superposition of the JV curve (the current density J is shown instead of I) in the dark and under illumination (the calibrated AM1.5G is referred to as onesun illumination). Under ideal conditions and when the cell is illuminated, the current-voltage characteristic can be modeled by a diode-like behavior and is shifted by a value corresponding to the light-generated current density (J ph ):

J(V ) = J 0 exp( qV kT
) -1 -J ph (1.1) where q is the elementary charge, k the Boltzmann constant, T the temperature. J 0 is the dark saturation current density, which represents the recombination in a device. By placing a photovoltaic device in an external circuit and measuring the current as a function of the applied voltage in dark and illumination conditions, the four main figures of merit for solar cells are available:

The short-circuit current density, J sc : under illumination, electron/hole pairs are formed in the absorber and collected by the external circuit. J sc is impacted by incomplete absorption, the generation and collection of photogenerated carriers, harmful recombinations and series resistance.

The open-circuit voltage, V oc : the balance between photogeneration and recombination when the cell is not connected to an external circuit.

The fill factor, FF: it determines the maximum power of a solar cell, and is defined as the ratio:

F F = J mp V mp J sc V oc (1.2)
where J mp and V mp are the current density and the voltage at the maximum power point P mp , respectively. FF is highly dependent on series and shunt resistances, as well as collection losses.

For a device with reasonable efficiency, it has a value in the range of 0.75 to 0.85. The power conversion efficiency, η: the portion of energy in the form of sunlight that is converted into electrical power. It is defined as the ratio of the maximum power density of the solar cell to the incident light power density (P inc ).

η = J sc V oc F F P inc (1.3)
Electrical measurements are performed under standard conditions, which consist of an incident power density P inc corresponding to the AM1.5G solar spectrum, with an integrated power of 1000 W/m 2 power. AM1.5G irradiation includes both direct and diffuse components of sunlight. Every solar cell technology has an intrinsic limit to its efficiency. The maximum theoretical efficiency of single-junction solar cells was modeled by Shockley and Queisser in the early 60s, considering the radiative recombination as the only loss mechanism [START_REF] Shockley | Detailed Balance Limit of Efficiency of p-n Junction Solar Cells[END_REF]; [START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF]].

-0.5 0.0 0.5 -40

-20 0 20 Dark Illuminated P mp J mp V mp V oc J sc FF= J 01 J 02 J SC R s V R shunt (a) (b) Voltage (V) Current density (mA/cm²) Figure 1.3: (a)
Example of a JV characteristic measured under dark and illuminated conditions. The V oc , J sc and P mp are indicated and the two regions used to calculate the FF are shown. (b) Equivalent circuit for a solar cell based on a 2-diode model. J sc is the photogenerated current, J 01 and J 02 the saturation currents, associated to diodes with ideality factor of n 1 and n 2 , respectively. R s and R shunt are the series and parallel resistances, respectively. Figure inspired from [START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF].

In real solar cells, the collection efficiency can be limited by resistive, shunt and recombinations losses. To characterize the different losses an electrically equivalent model is suggested in Figure 1. 3(b), and the non ideal current-voltage behavior of a solar device can be written as:

J = J 01 exp q(V -R s J) n 1 kT -1 + J 02 exp q(V -R s J) n 2 kT -1 + V -R s J R shunt -J ph (1.4)
J 01 and J 02 are the saturation current densities, R s and R shunt are the series and parallel resistances, respectively, and J ph is the photocurrent, which is equal to 0 for dark JV. n 1 and n 2 are ideality factors, which are set to 1 and 2, respectively, in the following to compare the saturation current density between all the characterized devices. Note that ideality factor values are generally comprised between 1 and 2, indicating which recombination type is dominant.

Depending on the device under investigation, the photocurrent can be voltage-dependent. This effect is visible in Figure 1.3(a), where the slope of the JV curve under illumination is steeper than in the dark over the voltage range from -0.5 V to 0.5 V. This emphasizes the usefulness of the dark JV measurements free of this phenomenon. There are several reasons for this voltage-dependent collection, such as interface recombination at the back interface or short diffusion length. The shifting approximation in equation 1.1 is no longer valid, unless the light generated current density is corrected by a voltage-dependent collection function of charge carriers. Moreover, in real solar cells, possible resistive, shunt and recombination losses must be considered and can be characterized (Figure 1.3(b)).

The current-voltage characteristics of the devices were measured at the C2N using a Oriel LCS-100 solar simulator and a Keithley 2400 as a source monitoring unit (SMU) to supply the voltage and measure the generated current connected in a two-probe configuration.

External Quantum Efficiency

The external quantum efficiency (EQE) is a spectrally resolved method for determining the J sc of a solar cell. It measures the probability of collecting a photogenerated electron per incident photon at a given wavelength. The J sc is obtained by integrating over all the photon energies the product of the EQE and the incident photon flux. Assuming that the cell response is independent of the incident power, the J sc is expressed by:

J sc = q λ 2 λ 1 EQE(λ)ϕ sun (λ)dλ (1.5)
where q is the elementary charge, and ϕ sun is the spectral photon flux of the AM1.5G solar spectrum at 1000 W/m 2 . For a GaAs cell, the wavelength range would be in between 300 nm and 1000 nm. [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF] The analysis of the EQE spectra provides insights into the origin of the collected current and the current losses. At short wavelengths, incident photons can be absorbed in the front layer of the solar cell stack and be visible in the EQE spectrum, indicating its partial contribution to the collected current. For the photon energies below the absorber bandgap, there is a sharp drop in the EQE, sometimes with small sub-bandgap absorption. Figure 1.4 shows the EQE measurement of a thin CdTe-based solar cell (1.2 µm) with corresponding current losses determined by optical simulation. Incomplete absorption loss is visible in the long wavelength part, due to an absorber thickness thinner than the standard device (4 µm).

At the C2N, the EQE is measured with a home-built setup using a Xenon light source coupled to a Horiba Triax monochromator. Incident light is focused on the sample at normal incidence using a microscope objective, and the solar cell is electrically contacted using twopoint probes connected to a SMU. The spectral response of the solar cell is derived from two successive measurements: first the current delivered by a calibrated Si diode of known spectral sensitivity is measured over the spectral range of interest, providing the photon flux and the number of monochromatic incident photons. The same scan is performed for the solar cell in the focal plane of the objective, in place of the diode. While the measurement is typically conducted at 0 V, it is possible to apply a voltage bias, to investigate its effect on the charge collection efficiency.

Luminescence properties

In a semiconductor, when an electron is excited into the conduction band, it can recombine through different recombination mechanisms:

-Defect assisted recombination or Shockley-Read-Hall (SRH) recombination: the electron relaxes in a trap state in the forbidden region in the bandgap due to defects in the material and recombines with a hole. The energy is dissipated as heat.

-Auger recombination: the energy released during the hole/electron recombination is transferred to a third charge carrier, for instance an electron in the conduction band. This electron thermalizes back down to the conduction band edge.

-Radiative recombination: an electron recombines radiatively with a hole, emitting a photon with an energy close to E g .

Luminescence resulting from radiative recombination depends on the band structure of the material and provides information about its electronic properties. The electrons in the semiconductor can be excited by different energy sources such as light (photoluminescence), the injection of excited charges (electroluminescence), or high-energy electrons (cathodoluminescence) [START_REF] Yacobi | Cathodoluminescence Microscopy of Inorganic Solids[END_REF]. The cathodoluminescence setup used during my thesis will be described in Chapter 4.

Overview of PV technologies

A wide variety of photovoltaic technologies are being developed, and regardless of their maturity, they are subject to continuous progress. Their laboratory efficiencies continue to increase over the years, as shown in the National Renewable Energy Laboratory's efficiency chart * or in the review by Green et al. [Green et al. 2022a]. Figure 1.5 compares the power conversion efficiencies for cells and modules of the major PV technologies.

With a 95% share of the PV market, the crystalline silicon is the leader of mainstream PV (referred to as 1 st generation of solar cells) [Fraunhofer 2023]. Its abundance on Earth, its nearly optimal bandgap energy for solar cells, and its low cost justify its significant expansion and the lack of room for competition from other semiconductors. Commercial panels available on the market have an efficiency between 19% and 22%. Despite its indirect bandgap, the silicon module has one of the highest efficiency (see Figure 1.5).

Thin-films semiconductors (CdTe, GaAs, CIGS, referred to as 2 nd generation of solar cells) are an attractive alternative to silicon cells. They have a direct bandgap and a higher absorption coefficient than Si, resulting in low material consumption as they require only a few µm of semiconductor as compared to Si which requires hundreds of µm. Thin films are well suited for building-integrated photovoltaics, and are ideal candidates for lightweight and flexible applications [START_REF] Ramanujam | Flexible CIGS, CdTe and a-Si:H based thin film solar cells: A review[END_REF]. New materials such as perovskite material show fast rising efficiencies since 2010. Yet stability issues and light degradation remain a very limiting drawback to their large-scale industrialisation. With substantially less material than Si, GaAs technology offers the highest-efficiency of single-junction solar cells (29.1%). However, III-V solar cells are expensive, and their commercial use is limited to spatial applications in the form of multijunction solar cells. Polycrystalline, inorganic thin-film solar cells represent 5% of the total PV market and are dominated by the CdTe technology (4%), with Cu(In,Ga)(Se,S) 2 -CIGS-materials for the remaining 1%. In the following, we will present both CdTe ad GaAs technologies in detail to emphasize their potential and their significant position in the development of PV.

Conclusion of the section

We have presented the very basics of the physics of solar cells and the main characterization techniques that will be used in the rest of the manuscript. A brief overview of the different PV technologies is given, emphasizing the predominance of silicon in the world market. Nevertheless, two other materials stand out: GaAs, with its record-breaking cell, and CdTe, with its high-performance and low-cost modules, which occupy a prominent place in the thin-film sector.

Overview of GaAs and CdTe solar cells 1.3.1 GaAs and III-V solar cells

General properties

GaAs is one of the most prominent materials for optoelectronic applications, especially photovoltaics. With its direct bandgap of 1.42 eV, this semiconductor could reach an efficiency of 33.2%, very close to the maximum Shockley-Queisser limit for single junction cells (about 34%). Its efficient optical absorption allows to design efficient solar cells with only 2-3 µm of active layer. The GaAs semiconductor has several intrinsic properties that put it at the forefront of PV technology. It is lattice-matched to a wide variety of binary and ternary compounds, enabling the growth of heterostructures with nearly perfect interfaces [START_REF] Bett | III-V compounds for solar cell applications[END_REF]. Its high-quality material offers the possibility to implement light trapping techniques, which enhance the photon recycling effect. It has a higher temperature stability and a higher electron mobility than silicon, turning it into the adequate candidate for space applications and concentrated PV. In the following, we will present the common structure of a single-juntion GaAs solar cell, with the different fabrication techniques proposed over the years.

Common architecture

The structure of a typical GaAs solar cell is shown in Figure 1.6(a). The usual structure of single junction GaAs consists of a few micrometer thick p-n homojunction made of a thin ndoped emitter and a thicker p-doped base. The Fermi level of the n-doped (p-doped) layer is closer to the conduction (valence) band than the Fermi level of the intrinsic semiconductor. At equilibrium, the Fermi level must be the same within the junction, resulting in a bending of the energy bands at the junction and leading to a built-in potential V bi . The band diagram of an ultrathin GaAs solar cell at short circuit, under 1 sun illumination, is given in Figure 1.6(b). The quasi-Fermi levels E Fn and E Fp characterize the electron and hole energy distribution when the semiconductor is illuminated. The chemical potential, ∆µ is defined as ∆µ = E F n -E F v = qV [Würfel et al. 2016]. The separation and the transport of the charge carriers is driven by the gradients of quasi-Fermi energies within the cells. This p-n junction is cladded between two higher bandgap III-V materials (typically AlInP, AlGaAs, or GaInP). The very thin layer on top of the active layer is called the window layer, while the layer below the absorber is called the back surface field (BSF). They act as passivation layers and carrier barriers, that repel unwanted charges and prevent detrimental recombinations. Ohmic contacts are achieved on highly doped GaAs (to avoid resistance) added on top of the window layer/substrate. They are typically made of annealed metal alloys such as AuGe/Ni. While many convincing ohmic contacts have been developed for GaAs based cells, it remains a hot topic in some respects: low cost contacts or formation at lower temperature are some examples [Wang et al. 1996b]; [START_REF] Huo | Comparison of Ti/Pd/Ag, Pd/Ti/Pd/Ag and Pd/Ge/Ti/Pd/Ag contacts to n-type GaAs for electronic devices handling high current densities[END_REF]]. Finally, anti-reflection coatings (ARC) are added at the front, necessary to produce high-efficiency solar cells. To reduce the reflectivity over a wide spectral range a double layer ARC (DLARC) is required [START_REF] Richards | Single-material TiO2 double-layer antireflection coatings[END_REF]]; [START_REF] Jung | Design and fabrication of multi-layer antireflection coating for III-V solar cell[END_REF].

GaAs growth and transfer

To achieve high efficiencies, III-V cells require high quality monocrystalline material with a complex layer stack and rely on expensive fabrication techniques. Semiconductor layers are grown sequentially in the same epitaxial chamber, by molecular beam epitaxy (MBE) or by metal organic chemical vapor deposition (MOCVD). Both techniques give rise to very good cells [START_REF] Pelzel | A comparison of MOVPE and MBE growth technologies for III-V epitaxial structures[END_REF]], but they require thick GaAs substrates, which represents a significant additional cost. This prevents GaAs from entering the market for terrestrial PV applications. The cost of III-V modules is estimated to be between 100-200 $/W compared to 0.2-0.3 $/W for silicon [START_REF] Horowitz | A Techno-Economic Analysis and Cost Reduction Roadmap for III-V Solar Cells[END_REF]; [START_REF] Smith | Photovoltaic ( PV ) Module Technologies : 2020 Benchmark Costs and Technology Evolution Framework Results[END_REF]]. We shed light on alternative technologies developed by the III-V community to effectively overcome the cost barrier through new growth techniques and substrate reuse.

Hydride vapor phase epitaxy (HVPE) is a fast and potentially lower cost semiconductor growth technique. Expensive metalorganic precursors are replaced by less expensive elemental sources [START_REF] Simon | III-V-based optoelectronics with low-cost dynamic hydride vapor phase epitaxy[END_REF]. The development of dynamic-HVPE has allowed the deposition of highly uniform films, the increase of the deposition rate of GaAs layers (300 µm/h compared to 100 µm/h for MOCVD) [START_REF] Metaferia | Gallium arsenide solar cells grown at rates exceeding 300 µm h-1 by hydride vapor phase epitaxy[END_REF]] and the growth of low defect dual junctions. The highest efficiency obtained so far for a GaAs single junction reaches 25.5%.

Epitaxial growth of III-V materials on low-cost substrates is also currently developed. For instance, proof-of-concept devices with an efficiency of 7.6% have been demonstrated based on single-crystal-like GaAs thin films (directly grown on metallic tapes), but they suffer from a very low V oc [START_REF] Pouladi | High-efficiency flexible III-V photovoltaic solar cells based on singlecrystal-like thin films directly grown on metallic tapes[END_REF].

The other alternatives are thin-film transfer and substrate reuse. They are enabled by several strategies such as epitaxial lift-off (ELO), spalling, cleavage of lateral epitaxial films for transfer [START_REF] Bozler | Ultrathin, high-efficiency solar cells made from GaAs films prepared by the CLEFT Process[END_REF] or remote epitaxy [START_REF] Kim | Remote epitaxy through graphene enables two-dimensional materialbased layer transfer[END_REF]. The ELO method is based on the high etching selectivity of AlAs (or AlInP) over GaAs [START_REF] Yablonovitch | Extreme selectivity in the lift-off of epitaxial GaAs films[END_REF]. The active layer is released from the substrate by chemical etching, and the almost unaffected substrate can then be reused. The resulting thin-films are material of choice for flexible and ultralightweight devices [START_REF] Geelen | Epitaxial lift-off GaAs solar cell from a reusable GaAs substrate[END_REF]]; [START_REF] Cheng | Epitaxial lift-off process for gallium arsenide substrate reuse and flexible electronics[END_REF]. Spalling, on the other hand, consists of exfoliating the epitaxial III-V layers from the parent substrate [START_REF] Bedell | Kerf-less removal of Si, Ge, and III-V layers by controlled spalling to enable low-cost PV technologies[END_REF]; [START_REF] Shahrjerdi | High-efficiency thin-film InGaP/InGaAs/Ge tandem solar cells enabled by controlled spalling technology[END_REF]]. The latter can be reused after a polishing step, that will limit its viability. To date, ELO has been identified as the most promising technique in terms of cost reduction for high volume manufacturing [START_REF] Ward | Optical analysis of the photon recycling effect in InGaAs/GaAsP multiple quantum well solar cell with light trapping structure[END_REF], and for highly flexible cells [START_REF] Cho | Ultrathin GaAs Photovoltaic Arrays Integrated on a 1.4 µm Polymer Substrate for High Flexibility, a Lightweight Design, and High Specific Power[END_REF].

Photon recycling

Besides the economic perspective, the transfer technique enables to process the back side of the devices, which turns out to be a remarkable asset for the III-V technology. Thus, in 2006, Bauhuis et al. reported a 26.1% efficient cell using ELO with a full-plate gold mirror at the back side. High care was taken with the metal contacts, and Pd/Ge/Au front contact annealed at low temperature was preferred [START_REF] Bauhuis | 26.1% thin-film GaAs solar cell using epitaxial lift-off[END_REF]]. In 2011, using a similar ELO technique, Alta Device Inc. reported a major breakthrough for GaAs solar cells with a 29.1% efficient device (J sc = 29.78 mA/cm 2 , V oc = 1.127 V, FF= 0.867) [Green et al. 2022b].

In Ref [START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF]] which reports the previous record GaAs cell (27.6%), the authors justify their high V oc = 1.107 V and FF = 0.841 with low dark saturation currents extracted using the two-diodes model. The resulting J 01 = 6×10 -18 mA/cm 2 and J 02 = 1×10 -9 mA/cm 2 , also suggest low shunt and series resistance. More importantly, it demonstrates the impact of the photon recycling on the V oc .

Under illumination, the absorption of photons will produce electron-hole pairs in semiconductors that can recombine radiatively. While most of the re-emitted photons escape out from the device, some may be re-absorbed in the main absorber. This re-emission, re-absorption process is called the photon recycling. Efficient recycling allows a higher carrier density to be built up in the device, because the re-emitted luminescent photons have a second opportunity to be absorbed and generate electron-hole pairs. Consequently, the quasi-Fermi level splitting increases, leading to higher V oc .

Miller et al. demonstrated the importance of having both a high internal luminescence efficiency and a very high reflectivity of a planar or textured mirror on the back side to achieve a significant V oc gain [START_REF] Miller | Strong internal and external luminescence as solar cells approach the Shockley-Queisser limit[END_REF]]. Thus, non-radiative recombination losses in the material can be dramatically limiting. Numerical models have evaluated the impact of photon recycling, and quantitatively estimated a 2% relative increase over conventional cells [START_REF] Walker | Impact of Photon Recycling on GaAs Solar Cell Designs[END_REF].

From homojunction stack to multijunction

Deep junction devices, are an interesting alternative to the more common shallow-junction devices (homojunction located near the back or the front side, respectively) [START_REF] Bauhuis | Deep junction III-V solar cells with enhanced performance[END_REF]; [START_REF] Geisz | Enhanced external radiative efficiency for 20.8% efficient single-junction GaInP solar cells[END_REF]]. This concept is also encountered in heterostructure devices. Hwang et al. achieved an efficiency of 28.7% (J sc = 30.0 mA/cm 2 , V oc = 1.108 V, FF= 0.865) with a n-GaAs/logarithmic-shape bandgap graded layer of Al x Ga 1-x As/p-AlGaAs stack an efficiency of 28.7% [START_REF] Hwang | Bandgap grading and Al0.3Ga0.7As heterojunction emitter for highly efficient GaAs-based solar cells[END_REF]]. The obtained V oc was higher compared to homojunction cells (∆V = 20 mV), and owes it to an improved built-in potential. While III-V compounds are largely represented by GaAs for single-junction solar cells, binary and ternary III-V semiconductors are widely used in multi-junction devices. To overcome the Shockley-Queisser limit and exploit a broader range of the solar spectrum with low thermalization losses, the idea is to combine several semiconductors stacked on top of each other. Figure 1.7 shows the solar energy that can be theoretically used by single and triple junction cells. For the multi-junction example, the top cell, with a higher bandgap, will absorb high energy photons and transmit the rest of the spectrum, lower energy photons will be absorbed deeper and so on. If the cells are connected in series (two terminals), the total J sc of the structure is limited by the lowest current supplied by any of the cells, while the total V oc is the sum of the individual V oc . Because of their high cost, III-V multi-junction cells are mostly applied for space application or under concentrated sunlight. Notable results have shown a triple-junction efficiency of 39.5% with GaInP/GaAs/GaInAs with multiple quantum wells [France et al. 2002]. Moreover, a 4-junction (GaInP/GaInAs/GaInAsP/GaInAs) developed at Fraunhofer ISE has demonstrated an efficiency of 47.6% under concentrated sunlight of 665 times the solar spectrum AM1.5D.

So far, GaAs (and III-V materials in general) cells have been developed mainly for spatial application, taking the lead in this field. Even with a high cost, GaAs remains highly relevant for ultrathin/ flexible devices. GaAs solar cells can also serve as a reference to develop new concepts that can be applied to other thin-film technologies.

CdTe thin films solar cells

CdTe is a II-VI semiconductor with a nearly ideal direct bandgap for PV application (1.5 eV), an absorption coefficient over 10 4 cm -1 in the visible region (λ=600 nm), and with a maximum PV conversion efficiency that could reach up to 32% [START_REF] Shockley | Detailed Balance Limit of Efficiency of p-n Junction Solar Cells[END_REF]. Figure 1.8(a) displays a schematic of a typical CdTe device structure. CdTe-based solar cells are fabricated in a superstrate configuration where the sun comes from the supporting glass (as opposed to the substrate configuration). A transparent conductive oxide (TCO) is deposited on a low cost glass substrate, and used as the front contact. This is followed by a p-n heterojunction, with an n-type material as the window layer and the CdTe as the p-type absorber layer, and finally a back contact is added. Over the last twenty years, the efficiency of CdTe-based solar cells has undergone a series of improvement steps (see Figure 1.8(b)), giving way to new structures and fabrication processes that differ significantly from the first generations.

Growth technique

The first generation is based on a CdS window layer with a relatively high bandgap of about 2.4 eV to avoid too much parasitic absorption in this layer, and a 4 µm-thick CdTe absorber [START_REF] Britt | Thin-film CdS/CdTe solar cell with 15.8% efficiency[END_REF]]. Due to the stability and robustness of CdTe, successful deposition can be achieved by a variety of fabrication techniques. CdTe can be deposited by radio frequency sputtering [START_REF] Gupta | All-sputtered 14% CdS/CdTe thin-film solar cell with ZnO: Al transparent conducting oxide[END_REF], chemical bath deposition [START_REF] Britt | Thin-film CdS/CdTe solar cell with 15.8% efficiency[END_REF], or MOCVD [START_REF] Zoppi | Grain and crystal texture properties of absorber layers in MOCVD-grown CdTe/CdS solar cells[END_REF]], however the most successful method is by close-space sublimation (CSS). Using CSS, the substrate (at a temperature between 450°C and 600°C) and the source are kept very close to each other in a high pressure chamber, allowing the material transport (bottom-up evaporation) [START_REF] Swanson | Single vacuum chamber with multiple close space sublimation sources to fabricate CdTe solar cells[END_REF]. Note that vapor transport deposition (VTD) is the industrialized CSS method, with a material transport by gases. Unlike other thin-film technologies as the CIGS, that require a high degree of control over elemental composition, CdTe can be deposited very rapidly at rates of micrometers per minute over large substrates. High temperature substrate deposition (>450°C) yields the best CdTe-based efficiencies [START_REF] Romeo | Low substrate temperature CdTe solar cells: A review[END_REF].

Post-deposition treatment

Despite the variety of possible deposition methods, the CdTe solar cells started with relatively low efficiencies. Critical parameters such as very low doping, back contacts, or the polycrystalline nature of CdTe with grain boundaries (GB) have limited the performance. GBs are the result of the misorientation between two grains, and are highly defective areas. They result in dangling bounds and traps that act as non-radiative recombination centers of carriers, which will significantly reduce the power conversion efficiency of the solar cells [START_REF] Moseley | Recombination by grainboundary type in CdTe[END_REF]. It has been demonstrated that the grain size is a critical parameter in terms of its effect on device performance, and in particular, on the V oc [START_REF] Jensen | Long carrier lifetimes in large-grain polycrystalline CdTe without CdCl2[END_REF]. Note that monocrystalline CdTe solar cells have achieved V oc higher than 1 V [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1V barrier[END_REF]], but such technique is not compatible with large-scale manufacturing.

To address the defective structure of the fast-deposited CdTe, a CdCl 2 post-deposition treatment has been introduced, and is partly responsible for the first wave of improvements highlighted in Figure 1.8. This can be done by physical vapour deposition or by immersion in a CdCl 2 -methanol solution. Although the positive effect of such treatment is undeniable, its exact impact is still under investigation [START_REF] Barkhouse | Wet chemical etching of cadmium telluride photovoltaics for enhanced open-circuit voltage, fill factor, and power conversion efficiency[END_REF]; [START_REF] Johnston | Photoluminescence imaging of large-grain CdTe for grain boundary characterization[END_REF]; [START_REF] Bidaud | Imaging CdCl2 defect passivation and formation in polycrystalline CdTe films by cathodoluminescence[END_REF]. CdCl 2 has been shown to increase the grain size by recrystallisation, and passivate GBs (Cl ions segregate at GB and passivate deep defect levels) [START_REF] Amarasinghe | Obtaining Large Columnar CdTe Grains and Long Lifetime on Nanocrystalline CdSe, MgZnO, or CdS Layers[END_REF]; [START_REF] Moseley | Recombination by grainboundary type in CdTe[END_REF]; [START_REF] Major | Control of grain size in sublimation-grown CdTe, and the improvement in performance of devices with systematically increased grain size[END_REF]; [Poplawsky et al. 2016b]. Improvements in carrier lifetime have also been demonstrated [START_REF] Metzger | CdCl 2 treatment, S diffusion, and recombination in polycrystalline CdTe[END_REF]]; [Shah et al. 2021b].

Progress has also been made in doping: CdTe grown at high substrate temperature usually tends to have a p-type conductivity due to Cd vacancies (V cd ), but this intrinsic doping is very low, about 10 13 cm -3 . For p-type configuration CdTe can be doped with group I or V elements, by post-deposition treatment (CuCl or more recently CuCl 2 ). This V Cd configuration is very suitable for Cu doping and the most efficient devices have been obtained by doping CdTe with copper, which has been confirmed by many literature studies [Bastola et al. 2019b]; [Munshi et al. 2018c].

Industrialisation and Te scarcity

For about a decade, the efficiency of CdTe-based devices was constant (see the second area in Figure 1.8(b)). However, fully automated in-line machines have been developed that can produce large CdTe modules from glass to panel in 3.5 hours with 18% efficiency † . CdTe technology has a reduced fabrication time and cost compared to silicon and has the best energy payback time (see Figure 1.9).

Although CdTe module production has an unmatched low price within the thin-film domain, the favourable forecast for CdTe technology can be challenged due to the limited supply of tellurium (Te), which is classified as a scarce element. In fact, Te is isolated as a by-product of mineral refining of copper (Cu), a more common element. CdTe PV is estimated to be materially limited to about 10 GW of annual production [START_REF] Woodhouse | Perspectives on the pathways for cadmium telluride photovoltaic module manufacturers to address expected increases in the price for tellurium[END_REF].

The CdTe PV industry has developed to over 6 GW in 2020, and the PV sector currently accounts for about 30-40% of global Te consumption [START_REF] Nassar | Global tellurium supply potential from electrolytic copper refining[END_REF]. As the efficiency of CdTe modules is lower than that of c-Si, its competitiveness relies heavily on its low cost production [START_REF] Smith | Photovoltaic ( PV ) Module Technologies : 2020 Benchmark Costs and Technology Evolution Framework Results[END_REF]]. The large-scale development could be threatened by the increasing price of scarce materials, but most of the published economic models assume that its effect is unpredictable (note that the Te cost/kg has increased by one order of magnitude since 2000).

A positive scenario expects an annual production of CdTe PV in the order of 50 GW/yr by 2030 at reasonable cost [START_REF] Fthenakis | Sustainability evaluation of CdTe PV: An update[END_REF]. However, most studies infer that by 2030, CdTe module should have an efficiency of at least 18% with 1 µm-thick absorber layer and the recovery of by-products should be optimized [START_REF] Mcnulty | Byproduct critical metal supply and demand and implications for the energy transition: A case study of tellurium supply and CdTe PV demand[END_REF]. [START_REF] Smith | Life cycle assessment of functional materials and devices: Opportunities, challenges, and current and future trends[END_REF]

Structural changes

The last wave of improvements (third zone in Figure 1.8(b)) involve major changes in the structure of CdTe-based solar cells. The CdS window layer has been replaced by the highly transparent MgZnO material (bandgap of 3.3 eV with Mg 0.15 Zn 0.85 O) [START_REF] Ablekim | Interface Engineering for 25% CdTe Solar Cells[END_REF]; [START_REF] Bittau | Analysis and optimisation of the glass / TCO / MZO stack for thin film CdTe solar cells[END_REF]]. The J sc is increased with MZO [Munshi et al. 2018c] because photogenerated carriers generated in CdS were not collected. In addition, the MZO bandgap can be tuned to match the band alignment with CdTe, allowing better V oc and FF [START_REF] Wang | Exploring window buffer layer technology to enhance CdTe solar cell performance[END_REF]]. Although great progress has been made to address the low doping issue with copper, its instability leads to a long-term degradation. Metzger et al. obtained highly doped CdTe layers (1 × 10 17 cm -3 ) with group V dopants using VTD. This results in a 20.8% efficient device with carrier lifetimes similar to Cu-doped CdTe devices (tens of ns).

Numerous investigations on the back contact have been performed to improve the V oc and will be developed in Chapter 5. Finally, the addition of selenium in the absorber layer has led to exceptional results, with a record cell efficiency of 22.1% (J sc = 31.69 mA/cm 2 , V oc = 0.887 V, FF= 0.785). The positive impact of Se will be discussed and investigated in Chapter 4.

Despite the multiple improvements in efficiency, cost-effectiveness production, and continuous characterization efforts, the performances of CdTe solar cells are still below expectations. The choice of the back contacts, the reduction of the defects in the bulk, the GB improvements and the formation of high-quality interfaces in the stack remain challenging and especially affect the V oc . Therefore, in this work, we will keep in mind the complexity of the device.

Aim of the thesis: towards ultrathin solar cells

Thin-films represent a promising option towards the fabrication of increasingly efficient solar cells but also the production of low-cost modules with significantly less material than the silicon technology. The substantial reduction of the thickness of the absorber would improve the thin-films technology, leading to additional benefits such as a higher industrial throughput (fabrication time saving), and a lower use of toxic elements. Subsequently, we are developing cells with reduced thickness, with the objective of achieving ultrathin cells with a target thickness about 10 times less than conventional cells.

The GaAs semiconductor offers to date ultimate possibilities of high-efficiency with reduced absorber thickness, and in 2019, a 19.9% efficiency record device with a 205 nm-thick GaAs absorber was demonstrated in our team [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]. The ever-increasing progress in ultrathin GaAs solar cells has driven us to develop new architectures to achieve ever higher efficiencies. In the first part of this thesis work, the main limitations of the first generation of ultrathin GaAs solar cells are theoretically investigated. To overcome the 20% efficiency limit we propose to modify the architecture of the back side to enhance light absorption within the absorber with less parasitic losses. An alternative strategy relying on a heterostructure is examined (Chapter 2). A proof-of-concept device with a new back side design is presented and analyzed in Chapter 3.

A second goal of this thesis is to apply these new light-trapping concepts to the low-cost CdTe technology. However, CdTe-based solar cells suffer from the low V oc which limits significantly their efficiency. Bulk defects, passivation interface and back contact issues are barriers to the achievement of efficient ultrathin CdTe solar cells. In Chapter 4, high-resolution cathodoluminescence mapping is used for the study of high-quality CdSeTe material, highlighting the essential role of selenium for the passivation of defects. In Chapter 5 several back contacts are explored for samples with different absorber thicknesses. Ohmic, passivating and optically reflective back contacts are experimentally studied to overcome the low V oc issue while being optically interesting for light-trapping implementation. In Chapter 6, the light-trapping strategy successfully implemented for GaAs is transferred to ultrathin CdTe solar cells. As introduced in Chapter 1, light absorption in semiconductors plays a key role in solar cells. We have already shown that there are many benefits to be gained from thinning the absorber thickness, regardless of the material. However, reducing the absorber thickness requires a light-trapping strategy to compensate for the low single-pass absorption in very thin films. A wide range of light trapping techniques have been developed in the community, to ensure high photocurrent and decent efficiency.

In this chapter we study ultrathin GaAs cells and offer some solutions to further increase the record efficiencies achieved so far. In section 2.1, we present an overview of light trapping strategies developed for photovoltaic solar cells with a state-of-the-art review of notable results. In particular we highlight the ultrathin GaAs cells developed by our team, for which a record efficiency for a 205 nm thick GaAs solar cell has been reported. Broadband light trapping was achieved using a nanostructured mirror on the backside, and resulted in a certified efficiency of 19.9% [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]. The simulation tools used for the corresponding optical calculations will be presented shortly. In section 2.2, we present a detailed optical loss analysis, and we propose an improved optical design with an increased theoretical short-circuit current mainly due to a lower parasitic absorption in the back mirror. In the final section 2.3, based on the detailed optical loss analysis and the improved optical design presented in section 2.2, we present a new solar cell architecture (III-V solar cell stack and nanostructured mirror) that integrates all the know-how the group has matured during the last years.

Light absorption in semiconductors

Light-trapping in semiconductors

As introduced in Chapter 1, light-matter interaction plays a fundamental role in solar cells: when a photon hits the surface of the semiconductor, it can be absorbed or not depending on its energy. However, the absorption in the active layer is also highly dependent on the optical properties of the material and its thickness. The absorption coefficient, α, is directly related to the imaginary part of the refractive index, κ, and is wavelength dependent: α = 4πκ λ . Figure 2.1(a) illustrates the behavior of the absorption coefficient for a variety of semiconductor materials of interest for PV. Compared to indirect bandgap silicon, direct bandgap semiconductors (GaAs, CIGS and CdTe) have a relatively higher absorption coefficient with a sharp edge. The simplest absorption model, valid for a beam of light impinging with an intensity I 0 in a semi-infinite medium, dictates that the light intensity I decreases exponentially with the distance (d) from the air/semiconductor surface. This is known as Beer-Lambert's law, stating:

I(d, λ) = I 0 e -α(λ)d
Nonetheless, for more complex structures, different intensity profiles can be expected due to the presence of light rays coming from reflection and transmission in other interfaces and layers. In the following, we describe different architectures for light-trapping and how the absorptivity can be expressed. It is calculated from the complementary part of the absorption, assuming zero front reflection (perfect ARC), i.e., A = 1-I/I 0 . A practical definition of the absorption is given by:

A(λ) = 1 -e -F α(λ)d (2.1)
where F defines the optical path enhancement factor, used as a figure of merit to assess the light-trapping efficiency.

Single-pass absorption, F=1

Assuming a semiconductor slab of thickness d, with perfect ARCs (perfect front/back transmission) and no backside reflection, the single-pass absorption at any wavelength is expressed by: For single-pass absorption, F=1. The penetration depth (δ) of a photon into the semiconductor before its absorption is determined by the relation δ = 1/α. High energy photons (blue) are absorbed at a shallow distance compared to lower energy photons (red) due to the higher absorption coefficient at shorter wavelengths (see Figure 2.2). This effect is directly visible in Figure 2.1(b) which illustrates the absorptivity of light in a single-pass configuration (solid lines) by GaAs slabs of different thickness. The long wavelength part of the spectrum becomes increasingly less absorbed with decreasing thickness. For a 200 nm-thick slab, about 36% of the spectrum is lost compared to a semi-infinite GaAs slab, while about 97% of the light is absorbed in a 2 µm-thick GaAs slab (the typical thickness of a GaAs solar cell).

A(λ) = 1 -e -α(λ)d (2.

Double-pass absorption, F=2

To prevent the transmission of red photons at the rear side, adding a perfect back mirror results in a double-pass absorption (photons entering the slab are reflected by the mirror) which is expressed by: Here F= 2. In Figure 2.1(b), the single-pass (solid lines) and the double-pass (dashed lines) absorption are plotted, calculated using equations 2.2 and 2.3 for d = 200 nm, 1 µm and 2 µm.

A(λ) = 1 -e -2α(λ)d (2.
For the 200 nm-thick slab, even with a planar perfect back mirror, the margin for improvement is still significant in particular in the long wavelength spectral range.

Lambertian scattering

Substantial light absorption is key to solar technology, and light-trapping strategies are being explored to compensate for poor material absorptivity or insufficient layer thickness. The idea is to increase the optical path length, or in other words, the distance that photons can travel within the absorber slab before escaping out of the device. In Lambertian light-trapping, one of the surface of the absorber is randomly textured so that light scattering follows the Lambert's cosine law and thus exhibits Lambertian reflectance: the brightness is independent from the observer view angle, the perfect "matte surface". In combination with a perfect ARC and a perfect back mirror (Figure 2.4) this design traps light in the absorber by total internal reflection, and gives rise to an intensity enhancement F= 4n 2 , where n is the refractive index of the absorber. This value was calculated by Yablonovitch, assuming weakly absorbing indirectbandgap semiconductors (such as silicon) in the ray optics regime. In other words, for a solar cell operating in a perfect Lambertian light-trapping scheme, a light ray is expected to undergo 4n 2 ≃50 passes inside the silicon absorber before escaping. The absorption under the Lambertian scattering model can be written as [START_REF] Green | Lambertian light trapping in textured solar cells and light-emitting diodes: Analytical solutions[END_REF]]; [START_REF] Yablonovitch | Statistical Ray Optics[END_REF]:

A(λ) = α(λ)d α(λ)d -1/F (2.

Multi-resonant absorption

In addition to incoherent scattering [START_REF] Campbell | Comparison of cathodoluminescence and photoluminescence of CdSeTe films grown on Si by molecular beam epitaxy[END_REF]; [START_REF] Depauw | Sunlight-thin nanophotonic monocrystalline silicon solar cells[END_REF]]; [START_REF] Eyderman | Light-trapping and recycling for extraordinary power conversion in ultra-thin gallium-arsenide solar cells[END_REF], advanced light trapping strategies such as plasmonic effect have been developed [START_REF] Collin | Nanostructure arrays in free-space: Optical properties and applications[END_REF]]; [START_REF] Atwater | Plasmonics for improved photovoltaic devices[END_REF]; [START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]. A subwavelength metal nanogrid in a metal /semiconductor/metal configuration resulted in high optical absorption in a 25 nm-thick GaAs absorber [START_REF] Massiot | Metal Nanogrid for Broadband Multiresonant Light-Harvesting in Ultrathin GaAs Layers[END_REF]]. However, plasmonic effect implies great losses in the metal.

Light-trapping strategies relying on coherent scattering (optical mode excitation) have also been investigated. The idea of nanophotonic light-trapping is to design in a controlled way the light-trapping scheme of the solar cell in order to couple the incident light into a discrete number of resonant or guided modes in the absorber. Depending on the arrangement and on the size of the nanostructures, several resonant modes can arise increasing the absorption in active layers. A comprehensive report of optical resonant modes is given in Ref. [START_REF] Brongersma | Light management for photovoltaics using high-index nanostructures[END_REF]. Materials with a high refractive index (e.g. semiconductors) can induce strong reflections, and when shaped into periodic nanostructures, provide ways to manipulate light and thus the optical path length. For example, engineered periodic patterns supported with highly reflective material can lead to Fabry-Perot resonances and guided modes in the slab of interest. The combination of several types of mode gives rise to multi-resonant absorption, and enhanced optical absorption in the long-wavelength range. All those resonances can be distinguished by their electric field response to the angle of incidence. Therefore, to fully understand this lighttrapping configuration and the mechanisms of optical resonances, the wave optics perspective is also required. We briefly present the main two types of resonant modes that occur in solar cells when implementing the light-trapping strategy developed in our team.

-Fabry-Perot resonances Fabry-Perot (FP) resonances result from the interferences caused by multiple round trips of the incoming light in the semiconductor, between the top surface and the metallic back reflector (or reflecting interfaces, the reflection coefficient, R is greater than 0 from the reflection interface contrast) (see Figure 2.5). This phenomenon is very similar to the one occurring in optical cavities of lasers. Its significance is featured, for example, in solar cells with multiple quantum wells [START_REF] Behaghel | Absorption enhancement through Fabry-Pérot resonant modes in a 430 nm thick InGaAs/GaAsP multiple quantum wells solar cell[END_REF].

FP resonances conditions require that the light, after one round trip in the cavity, returns in phase at the same location. Two waves interfere constructively when their phase shift is an integer multiple of 2π. We consider the light impinging normally in an absorber of thickness d and refractive index n, in between two semi-infinite media of refractive indices n 1 on the front side and n 3 on the back side. This leads to the phase shift presented in Equation 2 where k = 2πn λ 0 is the vertical component (projection in z-direction) of the wavevector in the layer. λ 0 is the wavelength, q is an integer defining the FP order. ϕ 21 and ϕ 23 are the phase changes acquired upon reflection at both ends of the layer cavity. For planar layers, they can be calculated from the Fresnel coefficient of reflection depending on the wave polarization. Note that for a multi-layer cavity of thickness d, we consider for the phase shift an averaged index defined as

n average = 1 d i h i n i (2.6)
where i is the number of layers, with their respective thickness h i

-Guided-modes resonances

We can also highlight guided-mode resonances, that result from the interaction between a waveguide and a grating. Note that Bloch modes for photonic crystal configuration or coherent multiple scattering with nanorod gratings are also part of the resonance mechanisms [START_REF] Collin | Nanostructure arrays in free-space: Optical properties and applications[END_REF]. A photon undergoing total internal reflection in a slab is an easy way to imagine a waveguide. It shows how light is coupled in an optical fiber: the refractive indices of the three media are tuned to suit the requirements, and the photons are trapped in the slab. Periodic nanostructures, on the other hand, act as gratings. Electromagnetic fields take on the same periodicity as the media. When hitting a grating, the wavefront is modified and its power is split into multiple discrete directions called diffraction orders (since the field should be continuous). When periodic nanostructures acting as a grating are placed next to a slab waveguide, guidedmode resonances arise. The incoming incident plane wave is diffracted into multiple modes, and at appropriate angles (the diffracted angle matches that of the guided-mode), can be coupled into a guided-mode (note that the wave also couples back out of the guide). In other words, the additional momentum induced by the diffractive nanostructures allows the coupling to electromagnetic guided waves in the absorber. When the resonance conditions are not fulfilled, the plane wave is just passing through (see Figure 2.6(a)). The structure can be considered as a slab characterized by an average refractive index. On the other hand, at resonance, part of the incident wave is coupled into a guided mode in the active layer, which is more intense than the initial wave, and the response is as shown in Figure 2.6(b). For a p-periodic crossed grating (2D-grating), light is diffracted in many directions. Taking into account the phase matching condition at the boundaries, all the possible transverse wave vectors that can exist inside the grating are defined by:

⃗ k x (q 1 ) = ⃗ k inc x + q 1 2π p ⃗ e x ⃗ k y (q 2 ) = ⃗ k inc y + q 2 2π p ⃗ e y ⃗ k t (q 1 , q 2 ) = ⃗ k inc t + q 1 2π p ⃗ e x + q 2 2π p ⃗ e y (2.7)
where ⃗ k inc t is the transverse component in x and y directions of the incident wave. q 1 and q 2 are the orders of the diffracted waves and can only take integer values. ⃗ e x and ⃗ e y are unit vectors in x and y directions. By generating guided-mode resonances in the active layer, the absorption is enhanced compared to a metallic reflector alone. But even more significant absorption can be achieved by partial overlapping of multiple resonances, which occurs when nanostructures are strategically designed. By tuning the nanostructure geometry, resonances are combined, leading to a broadband absorption enhancement from all modes (Figure 2.6(c)). Associated theoretical works suggest that such a light trapping could even go beyond the Lambertian scheme [START_REF] Collin | New limits for light-trapping with multi-resonant absorption[END_REF]; [START_REF] Brongersma | Light management for photovoltaics using high-index nanostructures[END_REF] Absorption in a semiconductor absorber solar cell is first determined by its intrinsic material properties (absorption coefficient). The addition of ARC coatings, back mirror and surface texturation can affect the way light enters, propagates, scatters, is trapped, and ultimately escapes from the semiconductor solar cell. Three models of light-trapping are considered to assess the performance of solar cells. The optical path length in the absorber can be increased of a factor F= 2 using a perfect ARC and back mirror (double-pass absorption), and up to F= 4n 2 using Lambertian scattering coupled with a perfect back mirror. Investigations on multi-resonant absorption have shown that coherent scattering is a highly attractive alternative to promote absorption in the active layer using a discrete number of resonances. Massiot et al. present an extensive review of the optical performance of Si, GaAs and CIGS ultrathin single junction solar cells, and the short circuit-current is used to assess the implemented light-trapping performance [START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]. In this part, we emphasize some key advances for silicon and GaAs solar cells.

State-of-the-art of ultrathin GaAs solar cells

Given the importance of silicon, light-trapping was first explored to increase the photocarrier generation rate in silicon solar cells whose bandgap is indirect [Wang et al. 1996a]; [START_REF] Blakers | 17% Eficient thin-film silicon solar cell by liquid-phase epitaxy[END_REF]. For solar cells thicker than 10 µm, the confinement of light is mostly designed with the combination of an ARC and a front texturation such as pyramids [Wang et al. 1996a]. However, for thinner cells this strategy is more difficult to implement due to a rapid wet etching. Although all the numerical studies exceed the double-pass absorption limit (some of them even exceed the Lambertian limit [START_REF] Bhattacharya | Towards 30% Power Conversion Efficiency in Thin-Silicon Photonic-Crystal Solar Cells[END_REF]; [START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]), only three significant experimental results have overcome this limit with devices thickness of 0.86 µm, 2.5 µm and 10 µm [START_REF] Depauw | Sunlight-thin nanophotonic monocrystalline silicon solar cells[END_REF]; [START_REF] Branham | 15.7% Efficient 10-µm-Thick Crystalline Silicon Solar Cells Using Periodic Nanostructures[END_REF]; [START_REF] Gaucher | Ultrathin Epitaxial Silicon Solar Cells with Inverted Nanopyramid Arrays for Efficient Light Trapping[END_REF]]. The light-trapping scheme is based on a subwavelength periodic patterning (either in the front or back side or both) coupled with a back side reflector. The spectral response in the long-wavelength range has been improved in conjunction with the increase in optical path length. Note that simulations with photonic crystals have demonstrated promising theoretical J sc of 21.8 mA/cm 2 with a ∼200 nm Si device [START_REF] Mallick | Optimal light trapping in ultra-thin photonic crystal crystalline silicon solar cells[END_REF].

Enhanced absorption in ultrathin GaAs solar cells

Strongly reducing the GaAs absorber offers remarkable alternatives for expensive GaAs cells, that could be used more thoroughly, with realistic potential for terrestrial applications [START_REF] Bosi | The potential of III-V semiconductors as terrestrial photovoltaic devices[END_REF]]. However as for thin silicon absorber, the light-trapping requires some investigations. Figure 2.7 takes up the figure of the review from Massiot et al. for GaAs, and summarizes the experimental short-circuit current J sc (filled circles) and the numerical short-circuit current J th (open circles) of standard GaAs to ultrathin GaAs solar cells reported in the literature. J sc values are compared with the three reference models described in section 2.1: single-pass absorption (continuous line), double-pass absorption (dash line) and Lambertian scattering model (dotted line).

First attempts with metallic nanoparticles on the front side, scattering the light in the semiconductor (confine sunlight by diffracting it beyond the total internal reflection angle), showed only a small increase of the J sc with very poor efficiency [START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]; [START_REF] Liu | Surface plasmon enhanced GaAs thin film solar cells[END_REF]]. EQE performance were improved at long wavelengths, but overall the devices suffered from parasitic absorption and the lack of a reflective rear side. Hence, the light-trapping developed using a front subwavelength nanostructuration with a dielectric TiO 2 material but combined with a back surface reflector resulted in a very high J sc = 22 mA/cm 2 for a 200 nmthick solar cell [START_REF] Lee | High Performance Ultrathin GaAs Solar Cells Enabled with Heterogeneously Integrated Dielectric Periodic Nanostructures[END_REF]. The etching of a window layer to fabricate a nanostructured front side was also reported [START_REF] Liang | Highefficiency nanostructured window GaAs solar cells[END_REF].

On the other hand, the back side of the solar cells can also be patterned. Significant improvements have therefore been demonstrated for 300 nm-thick GaAs transferred cells with a metallic rough back mirror [START_REF] Yang | Ultra-thin GaAs single-junction solar cells integrated with a reflective back scattering layer[END_REF]; [Eerden et al. 2020a]; [START_REF] Woude | Ultrathin GaAs solar cells with a high surface roughness GaP layer for light-trapping application[END_REF]. The combination of an AlGaAs contact layer easily chemically etched and a silver mirror deposited on localized contacts that enhances carrier collection [START_REF] Vandamme | Ultrathin GaAs solar cells with a silver back mirror[END_REF]] results in a J sc = 24.8 mA/cm 2 [Eerden et al. 2020a]. Few grating implementations have been reported in the literature for GaAs solar cells compared to silicon, but its potential has been demonstrated [START_REF] Paetzold | Nanoscale observation of waveguide modes enhancing the efficiency of solar cells[END_REF]. Moreover, a recent numerical study on the optimization of grating parameters for solar cells helps to grasp their full impact on the device efficiency [START_REF] Bläsi | Light trapping gratings for solar cells: an analytical period optimization approach[END_REF]. A 200 nm thick GaAs cells with a nanostructured TiO 2 /Ag mirror exciting guided modes in the waveguide-like structure has demonstrated a J sc of 24.64 mA/cm 2 and will be extensively detailed in the case study. Sayre et al. presented a 80 nm-thick cell exceeding the single pass limit with a rear side nanophotonic grating [START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF].

More recent studies have demonstrated the beneficial effect of nanostructuring the semiconductor layer at the back side of the cells associated with a flat mirror [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF]. Using this strategy, Micha et al. reported a theoretical J sc of 27.2 mA/cm 2 for a 100 nm-thick GaAs cell [START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF].

The examples developed above highlight two main methods of light trapping: random nanotexturing, which is cheaper but with non-optimized light scattering, and periodic nanostructures, which provide excellent and controlled scattering, but are more expensive. In their study on periodic vs random texturing Battaglia et al. insist on the importance of the modal structure and the necessary high number of coupled guided modes in the absorber, where periodic geometry should be superior [START_REF] Battaglia | Light trapping in solar cells: Can periodic beat random?[END_REF]]. Nevertheless, it has been suggested that breaking the symmetry would give access to asymmetric modes in the solar cells [START_REF] Yu | Fundamental limit of nanophotonic light trapping in solar cells[END_REF]]. Taking the best of both strategies, high diffractive efficiencies and broadband response, quasirandom (QR) structures have emerged. A new approach that transforms photonic crystals into quasi-random structures by modifying the spatial-Fourier space predicts a J th = 25.3 mA/cm 2 for a 100 nm-thick device. Resonances are enhanced and the result is robust to angle variation [START_REF] Buencuerpo | Engineering the reciprocal space for ultrathin GaAs solar cells[END_REF]. Buencuerpo et al. experimentally applied this QR photonic crystals pattern to a 260 nm-thick GaAs solar cell and achieved a J sc = 26.38 mA/cm 2 [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF]. In 2022, transparent semiconductor/dielectric quasi-periodic nanostructures have been shown to enhance the modal structure with absorption over a wide range [Camarillo Abad et al. 2022]. Despite these state-of-the-art results, it has also been demonstrated that periodic gratings and photonic crystals give excellent results [START_REF] Buencuerpo | Optically-thick 300 nm GaAs solar cells using adjacent photonic crystals[END_REF]] and sometimes outperform the QR strategy [START_REF] Xiao | Paths to light trapping in thin film GaAs solar cells[END_REF]; [START_REF] Camarillo Abad | Light management in ultra-thin solar cells: a guided optimisation approach[END_REF]. Note that the QR structure was metallic in one case and not optimized in the other. Both technologies (periodic vs QR) offer equally high performance, but perhaps QR designs offer more flexibility in terms of the conditions required to achieve optimum performance.

Table 2.1 summarizes the performance of the mentioned ultrathin GaAs cells. Emphasis is placed on the cells stack and on the type of nanostructuration implemented. Figure 2.7 outlines that even with surface texturization, it is still extremely challenging to surpass the numerical prediction for double-pass absorption with experimental solar cells, associated to electronic degradation. The record device J sc is placed as a comparison (noted as k). [START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. Short-circuit current of ultrathin GaAs solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models previously described. The values mentioned in the manuscript are in red and labeled. a. [START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]], b. Liu et al. [Liu et al. 2011], c. [START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF]], d. Camarillo Abad et al. [Camarillo Abad et al. 2020], e. Camarillo Abad et al. [Camarillo Abad et al. 2022], f. [START_REF] Buencuerpo | Engineering the reciprocal space for ultrathin GaAs solar cells[END_REF], g. [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF], h. van Eerden et al. [Eerden et al. 2020a], i. [START_REF] Lee | High Performance Ultrathin GaAs Solar Cells Enabled with Heterogeneously Integrated Dielectric Periodic Nanostructures[END_REF] , j. Chen et al. [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF], k. [START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF], l. [START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF], m. [START_REF] Massiot | Metal Nanogrid for Broadband Multiresonant Light-Harvesting in Ultrathin GaAs Layers[END_REF]. In 2019, our team in collaboration with the Fraunhofer Institute for Solar Energy (ISE) demonstrated a 19.9% efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and using a nanostructured silver mirror on the backside [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]; [START_REF] Chen | Ultrathin and nanowire-based GaAs solar cells[END_REF]]. Sections 2.2 and 2.3 will be dedicated to the improvement of this specific ultrathin cell. Therefore, care is taken to describe the stack, the cell fabrication, and to present the numerical and experimental results.

-Epitaxial stack:

The III-V semiconductor solar cell was epitaxially grown by Metal-Organic Chemical Vapour Deposition (MOCVD) at the Fraunhofer ISE on a GaAs substrate. This technique makes use of ultra pure gases injected near atmospheric pressure into the reactor. As the precursors (trimethyl-gallium Ga(CH 3 ) 3 and arsine (AsH 3 ) for GaAs) approach the surface substrate, they are thermally decomposed. The remaining atoms incorporate the elements into a new epitaxial layer by surface reaction, leading to crystal growth. Complex engineering of the substrate temperature, the pressure and the III/V ratio results in a high-quality control of the material growth.

The precise structure of the grown solar cell (layer material, thickness and doping level) is described in Table 2.2. The single-junction GaAs cell consists of a relatively common structure as described in section 1.3. Grown on a few hundred µm thick n-type GaAs(100) substrate, the stack consists of GaAs buffer, an AlGaAs etch stop to ease the removal of the substrate, n+GaAs/GaInAs contact layers, an n-AlInP window layer, a 205 nm GaAs pin homojunction as a main absorber: the 5 nm intrinsic GaAs is added for growth purposes only, a p-AlGaAs back surface field (BSF) and a p+GaAs contact. Wide-gap AlGaAs and AlInP alloys are used to passivate the GaAs surface and act as minority carrier blocking layers to reduce surface recombination losses. The high Al composition in the AlInP alloy aims to reduce parasitic absorption in the window layer. The solar cell was epitaxially grown "upside-down" (windows/ni-p/BSF) with the idea to fabricate the back-contact/mirror, add a foreign substrate, remove the GaAs substrate and finish the solar cell with a front busbar and ARC, as described in detail in the following section. After receiving the epitaxially grown solar cell, the fabrication starts with the rear side processing. First, localized metallic back contacts (5×5 µm, spaced with a period of 50 µm) are defined onto the p+GaAs back contact (9 × 10 18 cm -3 ) layer by optical lithography and lift-off of Ti/Au. The uncovered area of p+GaAs was etched in a mixture of citric acid and hydrogen peroxide. Etching is naturally stopped at the top of the AlGaAs layer (BSF). After the fabrication of the localized contacts, the nanostructured dielectric mirror was fabricated by directly embossing a TiO 2 sol-gel derived film followed by Ag deposition. To emboss the TiO 2 sol-gel, we used degassing assisted patterning (DAP) [START_REF] Luo | Degassingassisted patterning of cell culture surfaces[END_REF], a modified version of soft nanoimprint lithography that allows rapid embossing of sol-gel derived films [START_REF] Dalstein | Nanoimprinted, Submicrometric, MOF-Based 2D Photonic Structures: Toward Easy Selective Vapors Sensing by a Smartphone Camera[END_REF]. The Ag mirror side of the sample is bonded to a glass host substrate using OrmoStamp (micro resist technology GmbH), a flexible hybrid inorganic/organic polymer, and the GaAs substrate etched in a ammonia, hydrogen peroxid and water solution and the AlGaAs etch stop removed in HF:H 2 O. After this step, the III-V layer stack order is inversed and front metallic contacts (Ni-Au-Ge alloy) are fabricated using steps similar to those used for the localized back contacts. Uncovered areas of n+GaAs front contact (-9 × 10 18 cm -3 ) and n-InGaAs were etched in a mixture of citric acid and hydrogen peroxid. The final solar cells with areas of 1×1, 2×2 and 3×3 mm 2 were delimited with a photolithography mask and wet etching and a MgF 2 /Ta 2 O 5 DLARC was deposited. The final cell results are shown in Figure 2.8. More detailed information on cell fabrication are provided in the next chapter.

Material

-Cell performances:

The numerically optimized TiO 2 /Ag grating on the back of the solar cell enhances broadband absorption through multiple resonance overlaps (each of which has been identified). This results in a cell with a certified efficiency of 19.9% under AM1.5G illumination, where J sc = 24.64 mA/cm2, V oc = 1.022 V and FF = 0.792 (see Figure 2.8). Numerous light-trapping strategies including random texturing, periodic and quasirandom nanostructures have been implemented to enhance the absorption in ultrathin GaAs devices. The best results are obtained when scattering is coupled to a reflective backside. Simulations predict excellent short-circuit current values, however overcoming the double-pass absorption limit remains challenging for experimental devices (see Figure 2.7).

Electromagnetic simulations tools used in the thesis

OptiPV

The numerical simulations are performed using OptiPV, a numerical tool code developed by D. Micha and M. Giteau within our group for modelling and optimizing solar cell structures [START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF]. It accurately simulates the optical response of solar cell stacks with either only planar layers using the transfer matrix method (TMM) or nanostructured layers using the rigorous couple wave analysis (RCWA). It is mainly employed to improve the optoelectronic properties of devices and maximize the absorption in the active layer. In this code, several programs are linked together to optimize the computation time and the relevance of the simulated data.

OptiPV thus integrates Reticolo for RCWA, developed and supplied by P. Lalanne and J-P. Hugonin, and Solar Cell Capacitance Simulator (SCAPS) for the drift-diffusion model, written by M. Burgelman [START_REF] Burgelman | Advanced electrical simulation of thin film solar cells[END_REF]]; [START_REF] Burgelman | Modelling polycrystalline semiconductor solar cells[END_REF].

RCWA for optical absorption

Numerical methods that discretized Maxwell's equations have the advantage of being versatile and can be used for a wide variety of geometries. To perform absorption calculations with complex nanostructures in one or more layers, such as periodic diffraction gratings, we used the RCWA. It is based on an expansion of the electromagnetic fields based on Fourier harmonics and is fully adapted to the scattering problem with periodic structures. Reticolo * , a program written in Matlab and based on RCWA, is used to accurately simulate the diffraction grating problem [START_REF] Silberstein | Use of grating theories in integrated optics[END_REF]. Typical structure's schemes simulated with Reticolo are described in Figure 2.9. The cell is represented by a multilayer system (zdirection) with one layer or more, containing periodic sub-wavelength patterns. The structured layer(s) consists of two different materials: one is a nanopillar with refractive index n i,j embedded in a lattice of another material whose optical index is n i . For the whole study, we focus mainly on a two-dimensional grating with square blocks in the grid, using the same period p x = p y and side width d x = d y in the y-and x-directions. With such a design, the solar cell is independent of the polarization of the incident light at normal incidence. In this case, it means that waves polarized in TE mode (electric field is perpendicular to the plane of incidence) and in TM mode (magnetic field is perpendicular to the incidence plane) have a similar response.

RCWA calculates the electromagnetic fields by solving the Maxwell's equations, assuming an infinite xy horizontal plane. Each cell layer is discretized in the z-direction; in each of these equally thick slabs, the permittivity is constant with z. In the grating layer, the relative dielectric permittivity function ε(x,y) (transverse direction) is expanded in a Fourier series of the form:

ε r (x, y) = ∞ n=-∞ ∞ m=-∞ ε(m, n) exp(2iπ( mx p x + ny p y )) (2.8)
where ε(m, n) is the Fourier component of the relative permittivity in the grating, m and n are integers. The electromagnetic fields are computed. They are decomposed by a given number of Fourier harmonics, F x and F y , along the x and y axes respectively. Once the electric and magnetic fields are coupled in a linear differential equation [START_REF] Lalanne | Computation of the near-field pattern with the coupled-wave method for transverse magnetic polarization[END_REF]], it can be solved similarly to the eigenmode problem of a matrix (size of the matrix is 2 × (2F x +1) × (2F y +1)).

In the following, we will use 15 harmonics, unless otherwise stated. Once the eigenmodes in all layers are known, a scattering matrix method (series of propagating and evanescent plane waves) is used to solve the continuity equations at the interfaces of the layer stack. The wave propagation through the whole structure is fully calculated. Finally, the light absorption in the planar layer is derived from the Poynting's vector calculation, while the absorption in the grating can be obtained by integrating the electric field over the whole volume. We have access to the light reflection on the device and the absorption layer by layer as a function of the wavelength. Reticolo allows the numerical simulation of a large panel of periodic nanostructures including 1D architecture, but usually at the expense of computation time. For such models we will prefer the Transfer Matrix Method. It allows the calculation of the reflection and transmission coefficients of a stack of thin layers (thickness d ≈ λ). It relies on matrix products, which makes it relatively fast.

For the last section of this chapter, a large number of parameters have to be optimized for the different simulated structures. As already mentioned, due to time constraints, it is not possible to calculate every scenario. In this respect, the optimizations are carried out using a genetic algorithm (GA) that was implemented by J. Zeitouny in our team. The method is based on mutation and crossing operators which are applied to expand a population of individual solutions. Then, the method performs a selection on individuals based on their performances. The GA repeatedly modifies a population of individual solutions, over successive generations the population reaches an optimal solution.

Conclusion of the section

Typically, absorbers are kept relatively thick to ensure the absorption of most of the incident photons in a single-pass. However, with a suitable light-trapping scheme, thinner solar cells can be made optically thick. The absorber thickness can be significantly reduced by more than an order of magnitude while preserving high current densities. This approach has been extensively developed for c-Si and CIGS cells, but here we focus mainly on GaAs results. To counterbalance low absorption, light scattering has been widely developed. However, strategies involving periodic patterning have shown major enhancement in absorption due to multi-resonances. Periodic nanostructures enable incident light to be coupled into resonant modes in the structure, leading to a partial overlapping of resonances and a broadened absorption spectrum. This technique, developed by our team, has led to a record-breaking ultra-thin solar cell with an efficiency of 19.9%, based on a 205-nm-thick GaAs absorber. Prior to fabrication, the RCWA method is used to calculate the absorption with periodic diffraction gratings in the cell stack.

In the following, the original 205 nm-thick GaAs solar cells developed in the team will evolve with respect to many physical characteristics. We sketch the different axis of work in Figure 2.10

Optical losses and advanced light-trapping in 200 nm-thick GaAs solar cell

A detailed loss analysis on the ultrathin record device has been conducted and provides pathways to higher efficiencies [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]. In the following, the main factors limiting the efficiency will be outlined and we will address these identified challenges. The 205 nm nanostructured cells developed by Chen et al. will be referred as the initial/original structure in the following [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]].

2.2.1

Parasitic absorption losses in state-of-the-art 200 nm-thick GaAs solar cell Figure 2.11(a) shows the calculated absorption in the GaAs absorber and in the other layers of the GaAs solar cell structure presented in Figure 2.8. The identified losses limiting J sc are summarized in Figure 2.11(b).

The simulation of the absorption in the ultrathin GaAs cell, is based on the RCWA method (described in Section 2.1). The x-z plane is considered as the plane of incidence, and impinging waves are linearly polarized. We consider the TE mode i.e. with the electric field transverse to the plane of incidence. The absorption is calculated in each layer over a specified wavelength range with Fourier orders F x = F y = 15. The theoretical short-circuit current density J th is defined as:

J th = q λ 2 λ 1 A(λ)Φ sun (λ)dλ (2.9)
where q is the electron charge, A(λ) is the calculated absorption in the GaAs layer and Φ sun (λ) represents the standard AM1.5G solar spectrum (number of photons per unit time, per unit area and per unit wavelength). λ 1 and λ 2 represent the interval on which J th is calculated. A perfect collection of photogenerated charges is assumed: the absorption of one photon generates an electron-hole pair and any minority carrier generated in the absorber is collected. Therefore, J th is an upper bound in the simulation to characterize the performance of the light trapping architecture.

Jsc (mA/cm 2 )

31.9

25.1 In Figure 2.11(b), the top value corresponds to the calculated radiative limit of 31.9 mA/cm 2 , for a 205 nm-thick GaAs using the Lambertian scattering model [START_REF] Yablonovitch | Intensity Enhancement in Textured Optical Sheets for Solar Cells[END_REF]; [START_REF] Green | Lambertian light trapping in textured solar cells and light-emitting diodes: Analytical solutions[END_REF]]. The experimental J sc (with and without the front contacts shadowing) are highlighted in bold. The J sc = 24.6 mA/cm 2 was measured from the J-V illuminating the entire solar cell, while the best short-circuit current was calculated from the EQE measured in a region without front contact shading (5-6% of the total area) results in J sc = 25.1 mA/cm 2 . The latter is very close to the numerical calculation that predicted 25.6 mA/cm 2 (see Figure 2.11(a)) for the optimized TiO 2 /Ag back mirror. The small discrepancy can be attributed to manufacturing errors during the fabrication that led to geometrical parameters slightly different from the nominal ones used for the electromagnetic simulations. That being said, in the simulation the J th was calculated assuming a perfect collection of photo-generated carriers inside the GaAs absorber while in the experimental results we cannot exclude to collect also some of the carriers photo-generated in the AlInP window layer. Note that, although the GaAs bandgap is around 872 nm, the theoretical short-circuit current (J th ) in the active layer is calculated by integrating over the whole spectral range. The EQE data in Figure 2.8(b) confirms a relatively important sub-bandgap absorption, justifying this methodology. Moreover, it has also been demonstrated that the Franz-Keldysh effect has a significant impact on the EQE of ultrathin GaAs cells. Subbandgap absorption is observed and is attributed to the electric field-induced modification of the near-bandgap absorption coefficient, leading to oscillations in the EQE above the bandgap [Eerden et al. 2020b].

The optical losses include parasitic absorption in the window layer (dJ = 1.3 mA/cm 2 ), in the BSF (dJ = 1.2 mA/cm 2 ), in the Ag mirror (dJ = 2.7 mA/cm 2 ) and due to reflection (dJ = 1.1 mA/cm 2 ). Each contribution to the parasitic absorption is highlighted in Figure 2.11(a) with colored area. As already mentioned, the losses in the window layer may be slightly overestimated by comparing the experimental EQE (red dashed curve) with the simulation, where the EQE at short wavelengths outperforms simulated absorption inside the GaAs (red solid curve). The total current density J tot and the current density loss dJ in the silver back mirror were calculated in the range 300-910 nm. This study highlights that the parasitic absorption losses in the nanostructured metallic reflector is the main contributor to J sc losses. Leaving aside the losses inherent to the structure, J sc would also improve by reducing reflection losses in the longwavelength spectral region, where the ARC is less effective. In the following paragraph, we will address this challenge by modifying the architecture of the back mirror to reduce its parasitic losses while reducing the reflection losses in the red part of the spectrum.

A detailed analysis of the parasitic absorption losses conducted on the record ultrathin GaAs cell shows that reducing the losses in the nanostructured silver mirror and the reflection losses in the long-wavelength spectral region can improve the J sc by 3.8 mA/cm 2 without changing the current architecture of the III-V solar cell.

Improved light-trapping scheme

Periodic gratings allow strong coupling to optical modes and help to achieve optimal solar cell performance. However, it is shown that metallic grating structures, even with high reflectivity, drastically reduce the expected performance, due to strong parasitic absorption. The latter could be avoided by using dielectric nanostructures in direct contact with the absorber, followed by a back flat metallic reflector or directly using all-dielectric reflectors [START_REF] Lare | Dielectric back scattering patterns for light trapping in thin-film Si solar cells[END_REF]; [Fu et al. 2015]. Distributed Bragg reflectors are also a valid alternative as they offer very high reflectivity and are radiation resistant, a key feature for space applications. However, they consist of multiple layers with different refractive indices and are often too expensive to be considered as a viable solution for GaAs single junctions [START_REF] Tobin | Enhanced light absorption in GaAs solar cells with internal Bragg reflectors[END_REF]; [France et al. 2018]. Interestingly, high refractive index contrast between materials is expected to result in an efficient grating for light trapping in solar cells [START_REF] Brongersma | Light management for photovoltaics using high-index nanostructures[END_REF].

Keeping the original GaAs solar cell structure presented in section 2.1.2, we seek to improve the design of the nanostructured back mirror to reduce the parasitic absorption in the silver. We also expect that this will have a positive impact on the amount of light reflected at the front surface. We retain a similar light trapping strategy based on periodic grating made of a high refractive index (RI) material to enhance the absorption by overlapping guided mode resonances, but coupled with a flat silver mirror. .12: Schematic of an ultrathin GaAs cell with a novel nanostructured back contact including a grating and a planar reflector on the rear side. The scheme on the right shows the details of the nanostructured layer with a-Si nanopillars embedded in a dielectric nanogrid. The red parameters: height of the grating h N , period p and width w of the nanopillars, refractive index and thickness under the grating of the dielectric n and h will be optimized in the numerical simulation.

The new design is shown in Figure 2.12. It consists of a high RI grating in proximity of the BSF (and ultrathin GaAs absorber layer), a low RI spacer and a flat silver reflector. A twodimensional grating with the same periodicity in both x-and y-directions is used to simplify the simulations and maintain a polarization-independent response at normal incidence. In order to efficiently scatter light, the 2D grating is composed of a relatively high RI (comparable to the RI of III-V semiconductor) and negligible extinction coefficient (to avoid parasitic absorption). This high RI grating must be sufficiently separated from the flat silver mirror to avoid the excitation of surface plasmon polariton inside the metal. The use of a low RI material as spacer in between the grating and the silver mirror should result in a high RI contrast (to increase the scattering efficiency of the grating) and reduce the losses in the flat silver mirror [START_REF] Ma | Phase shift and penetration depth of metal mirrors in a microcavity structure[END_REF]]. Set aside III-V semiconductors such as AlInP, AlGaAs and GaP, transparent and high RI materials, in the visible spectral range of interest here, are rare. Titanium dioxide (TiO 2 ) and Lead titanate (PbTiO 3 ) have relatively high RI (respectively n = 2.52 and n = 2.61 at λ=800 nm) * , but a grating made of these materials won't ensure a sufficiently high scattering efficiency. For this reason, we have explored the use of a-Si:H which has a much higher RI (n= 3.87 at λ = 800 nm) and a relatively high bandgap of about 1.77 eV (700 nm). Nonetheless, it should be noted that the a-Si:H bandgap is not high enough to avoid parasitic absorption. Most of the photons in the 350-700 nm spectral range will be absorbed by the GaAs absorber in a single-pass, so that we can expect to have a relatively modest parasitic absorption within the a-Si:H grating. The a-Si:H grating can be embedded (planarized) using a low RI dielectric such as mesoporous SiO 2 synthetized by sol-gel method, which can attain a RI about n=1.2 providing a porosity approaching 50%.

In the following, the figure of merit J th is used to optimize the back side design by numerical calculations. The goal is to maximize J th in the absorber layer.

The simulated stack is shown in Figure 2.12, with the different layers, their respective thick-ness and the top view schematic of 4 unit cells of the nanostructured layer. We have considered a grating composed by an array of a-Si:H squared nanopillars in order to limit the parasitic absorption inside the a-Si:H grating: similar scattering properties can be optimized using the complementary pattern design (a-Si:H square nanogrid), but it would result in greater a-Si:H parasitic absorption due to the greater amount of a-Si:H used in this geometry. Compared to the initial structure, a new ARC stack was optimized with thicknesses of 90/60 nm for MgF 2 /Ta 2 O 5 .

In order to optimize the geometry of the ultrathin cells under the aforementioned conditions, five main parameters of the rear side design require optimization and are highlighted in red in Figure 2.12. The grating thickness, i.e. the nanostructures height (h N ), the pattern period (p), the width of the a-Si:H nanopillars (w), the refractive index (n) of the dielectric embedding the a-Si:H grating and distance of the a-Si:H grating from the mirror (h), composed of the same dielectric and hereafter referred to as the "spacer". The refractive indices used in the simulation are plotted in Figure 2.13. For the GaAs, the data are extended in the sub-bandgap region with an exponential Urbach tail of energy width 20 meV and 10 meV for p-GaAs and n-GaAs respectively. The dielectric material is simulated using a constant refractive index and non-absorption (extinction coefficient equal to 0) over the entire spectral range of interest (a fair approximation for mesoporous SiO 2 ). [START_REF] Gao | Exploitation of multiple incidences spectrometric measurements for thin film reverse engineering[END_REF]], Al 0.51 In 0.49 P [START_REF] Schubert | Open-air printing of Cu2O thin films with high hole mobility for semitransparent solar harvesters[END_REF]], Al 0.4 Ga 0.6 As and MgF 2 [START_REF] Palik | Handbook of Optical Constants of Solids Cited on page 39[END_REF]], GaAs [START_REF] Sturge | Band Alignment for Rectification and Tunneling Effects in Al2O3 Atomic-Layer-Deposited on Back Contact for CdTe Solar Cell[END_REF]]; [START_REF] Palik | Handbook of Optical Constants of Solids Cited on page 39[END_REF], Ag [START_REF] Jiang | Realistic Silver Optical Constants for Plasmonics[END_REF]],a-Si:H.

Once the structure is optimized (J th maximized), the new design of ultrathin GaAs cells with aSi/dielectric grating and planar Ag mirror can be described as follows:

-The grating height, h N = 100 nm -The period, p = 750 nm -The a-Si:H nanopillars width, w = 340 nm (representing 45% of the period) -The dielectric nanogrid refractive index, n =1.2 -The dielectric thickness below the grating, h = 180 nm Note that the lower value of the refractive index was limited to 1.2 to keep the experimentally achievable value realistic. Sol-gel layers of porous SiO 2 can meet such optical index and fit the requirements necessary for nanopattering using, among others, the nanoimprint lithography (NIL) technique [START_REF] Dalstein | Nanoimprinted, Submicrometric, MOF-Based 2D Photonic Structures: Toward Easy Selective Vapors Sensing by a Smartphone Camera[END_REF]; [START_REF] Bottein | Environment-controlled sol-gel soft-NIL processing for optimized titania, alumina, silica and yttria-zirconia imprinting at sub-micron dimensions[END_REF]]. The parameters are varied in practical steps, taking into account the experimental resolution limit. Figure 2.14(a) shows the absorption spectrum in the GaAs absorber of the optimized a-Si:H/SiO 2 /flat Ag mirror design (solid yellow line) and the absorption spectrum in the GaAs absorber of the original structure using a TiO 2 /Ag back mirror (dashed yellow line). The calculated theoretical short-circuit current is 28.23 mA/cm 2 , giving a J th gain of 2.63 mA/cm 2 over the original architecture. In the same figure are plotted the absorption spectrum of the a-Si:H array of nanopillars (purple line) and the parasitic losses in the silver flat mirror (gray line). The new design provides further improvement in light absorption with additional peaks from λ=700 nm to 900 nm and a drastic decrease of the parasitic absorption in the mirror. In fact, the absorption in the GaAs is almost equal to the total absorption in this region. The calculated short circuit current loss in the silver layer drops by more than 2 mA/cm 2 as compared to the previous structure (see Figure 2.11). In addition, the J sc lost by reflection decreases dramatically, accounting for only 0.4 mA/cm 2 of the total loss in the new structure. The expected losses in the a-Si:H are relatively small, accounting for 0.61 mA/cm 2 and will be further described in the following. The updated estimated loss analysis for J sc is shown in Figure 2.14(b). Significant progress was made in improving J sc and reducing parasitic absorption. To decrease the parasitic losses in the silver mirror, the TiO 2 /Ag nanostructured mirror is replaced with an a-Si:H/dielectric grating embedded in a low RI material and coupled with a planar Ag mirror at the back side. The new design of the ultrathin GaAs cells leads to a theoretical short-circuit current density in the active layer of 28.23 mA/cm 2 . Despite some parasitic absorption in the a-Si:H, the net J sc increase corresponds to 2.63 mA/cm 2 , which would correspond to an increase in the efficiency of about 2.9%.

Influence of the dielectric refractive index and the period

To glean insights into the effect of the variation of the dielectric index and of the period on our improved device, the specular reflectance 1-R of the cell with a-Si:H/SiO 2 / flat Ag mirror was calculated. In fact, the absorption in each layer is not calculated due to time constraints and the total absorption of the whole stack is derived from 1-R. The ultrathin and optimized GaAs cell is unchanged while the dielectric index is varied from 1.1 to 2 in steps of 0.01. Figure 2.15(a) shows the total absorption as a function of the wavelength and the dielectric real index. The corresponding current density, J tot , derived from the total absorption is plotted as a function of the index. The maximum J tot value is found for n = 1.19. It is very close to J tot at 1.2, confirming our previous calculation. J tot continues to decrease for higher indices, followed by a slight surge to about 1.9. The absorption map suggests fewer resonances for higher dielectric indices in the [600-880] nm range.

The effect of the grating period is also analyzed in detail. 1-R is plotted as a function of the wavelength and grating period from 400 to 900 nm in Figure 2.15(b). In this particular case, the fill factor (f f = pxpy dxdy ) is kept constant at 45%, i.e. the width of the nanopillars is updated for each simulated period value. Again, J tot is plotted against the period on the right of the figure. It emphasizes that the grating period does not play an important role below a wavelength of 600 nm: the vertical resonances are clearly independent of its value. On the other hand, beyond 600 nm, the variation of the grating period strongly favors the formation of different resonances that overlap and generate broadband absorption.

Comparing this dispersion diagram with the one calculated in Figure 3e by Chen et al. [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF] for the original ultrathin GaAs structure, it is straightforward that more resonances are confined in the spectral range of interest (the vertical dashed line indicates the GaAs bandgap) for the new rear design. The evolution of J tot with respect to the grating period variation shows that the maximum value is reached at p = 748 nm, very close to our simulation results at p = 750 nm. Very small variations of J tot are visible in the interval 700-750 nm and it demonstrates that a variation of p =750±10 nm generates less than 2% loss for J tot . While higher absorption occurs as the period increases, there is a steep decrease above p = 750 nm. The diffraction line for p = λ is indicated by a dashed line. The absorption loss induced by the a-Si:H layer is visible between 600 and 650 nm regardless of the period.

Parasitic losses in the a-Si:H grating Figure 2.15(b) highlights that there is still a non negligible loss in the a-Si:H (dJ=0.61 mA/cm 2 ) due to the semiconductor absorption for wavelengths smaller than 700 nm. Although amorphous silicon has a very high refractive index, it is not the optimal material for the grating due to its absorption. It could be avoided by using alternatives such as very high bandgap semiconductors such as silicon or gallium nitride (SiN or GaN), or high index dielectric layers such as silicon carbide deposited by PECVD or titanium oxide. However, compared to a-Si:H whose refractive index n(λ=800 nm) = 3.87, these materials display much lower values in the range 2 to 2.6 at λ = 800 nm [START_REF] Luke | Broadband mid-infrared frequency comb generation in a Si3N4 microresonator[END_REF]; [START_REF] Barker | Infrared lattice vibrations and free-electron dispersion in GaN[END_REF]; [START_REF] Singh | Nonlinear optical properties of hexagonal silicon carbide[END_REF]; [START_REF] Devore | A Fast and Easy-to-use Modeling Tool for Scanning Electron Microscopy and Microanalysis Users[END_REF]. A RI of 3.25 obtained at λ = 800 nm for a silicon-rich nitride film is worth mentioning [START_REF] Ng | Exploring High Refractive Index Silicon-Rich Nitride Films by Low-Temperature Inductively Coupled Plasma Chemical Vapor Deposition and Applications for Integrated Waveguides[END_REF], with the extinction coefficient dropping quickly to zero at λ = 550 nm.

Preliminary calculations have been performed, where the amorphous silicon nanopillars are replaced by a non-absorbing material with a refractive index of 2.6. The best optimization so far gives a J th = 27.65 mA/cm 2 , below the J th obtained with the a-Si:H/SiO 2 /flat Ag mirror back structure.

The absorption loss in a-Si:H challenge will be addressed in the next section 2.3 where an alternative ultrathin GaAs solar cell is proposed and the a-Si:H grating replaced with epitaxially grown wide-bandgap III-V semiconductors. 2.16(a) presents the total absorption as a function of the incidence angle for both TE and TM polarizations. At short wavelengths, little variation with the incidence angle is observed. It becomes more important above 600 nm, especially for the TM polarization. As for the initial structure, we see in Figure 2.16(b) that with our improved back grating, the short-circuit current J th is above 27 mA/cm 2 up to an angle of 60°.

Angular dependence

Conclusion of the section

Based on numerical simulations, a detailed absorption loss analysis was performed on the 205 nm-thick GaAs cell. It was shown that, apart from small J sc losses due to contacts shading and fabrication, the losses can be divided into two groups. First, the parasitic absorption resulting from the choice of the cell layers, i.e., the window layer and the back surface field. Second, the absorption lost in the nanostructured mirror and reflection, which is directly related to the back-mirror architecture.

In this section we have modeled a modified back-mirror design to further optimize the devices while keeping a light-trapping scheme based on a 2D square periodic patterning.

The nanostructured mirror architecture is now composed by an a-Si:H grating and a flat silver mirror opportunely spaced from the a-Si:H grating using a low RI dielectric material. The best structure shows a predicted short-circuit current J th of 28.2 mA/cm 2 , corresponding to a 2.63 mA/cm 2 improvement as compared to the original design. This improved short-circuit current derives from reduced parasitic losses inside the silver mirror and reduced reflection losses. We have shown the effect of the grating period on the resonances and a relatively large leeway on the choice of the refractive index of the low RI dielectric material. Moreover, the calculated short-circuit current is essentially insensitive to the variation of the angle of incidence, with a small decrease above 60°. The a-Si:H parasitic loss accounts for 0.61 mA/cm 2 and are relatively small, the use of a III-V semiconductor with a higher bandgap (AlGaAs with Al>50%) could reduce even further these losses.

New optoelectronic design for ultrathin GaAs solar cell

In the previous section, the new back-mirror design has proven its worth, with a J th gain of 2.63 mA/cm 2 , as compared to the original TiO 2 /Ag back mirror. Thus, a safe and effective strategy is to keep the grating/flat mirror configuration, but the parasitic absorption inside the a-Si:H is still relatively important. It should be noted that parasitic absorption occurs equally in the AlInP window layer (dJ= 1.3 mA/cm 2 ) and in the AlGaAs BSF (dJ= 1.2 mA/cm 2 ). However, these two layers are necessary for passivation and selective collection of carriers. The AlInP window layer composition is as transparent as possible, and its thickness is similar to the ones used for conventional thicker GaAs solar cell. On the other hand, the AlGaAs BSF can be more easily optimized to decrease the parasitic absorption. Moreover, the novel back structure only improved the J th , with no impact on the V oc . Keeping in mind these considerations, an improved stack for the ultrathin GaAs solar cell is proposed. Based on our previous experience and recent breakthroughs in the community, we propose to replace the current design, comprising a n-AlInP window layer / 205 nm-thick GaAs n-i-p homojunction / p-AlGaAs BSF, with a modified stack comprising a n-AlInP window layer / 200 nm-thick n-GaAs emitter / p-AlGaAs base / p-AlGaAs BSF. The light-trapping scheme is now composed by a wide-bandgap AlGaAs grating, a low RI spacer and a flat silver mirror. The p-i-n homojunction is updated with a p-n heterojunction for a higher V oc , while the rear side design maintains a flat silver mirror to avoid parasitic losses.

Heterojunction solar cell for enhanced V oc

Although it is expected that ultrathin devices will have a higher V oc than standard devices, there is still a 100 mV difference compared with the record standard cell (V oc = 1.127 V) [Green et al. 2022a]. A possible approach to increase the V oc is to modify the absorber from homojunction to heterojunction, adjusting its thickness to achieve a deep junction (close to the back mirror). Moreover, outstanding outcome obtained for V oc using a heterostructure in [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF] tends to confirm the viability of such option.

Investigations by Steiner et al. have shown that the simulated V oc increases when the back contact layer is thinned down (reduced parasitic absorption) and when the absorber is reduced [START_REF] Steiner | Optical enhancement of the open-circuit voltage in high quality GaAs solar cells[END_REF]]. On the one hand, this highlights the advantage of bringing the back mirror closer to the absorber. On the other hand it highlights the interest in reducing the thickness of the device, although unfortunately no clear trends could be demonstrated experimentally. Besides the impact of the backside texturing on photon recycling, it has been demonstrated that photon recycling is also affected by the front contact grid [START_REF] Eerden | Comprehensive analysis of photon dynamics in thin-film GaAs solar cells with planar and textured rear mirrors[END_REF]; [START_REF] Gruginskie | Limiting mechanisms for photon recycling in thin-film GaAs solar cells[END_REF]. A textured back mirror minimizes the photon recycling (many scattering events lead to a higher escape probability compared to a planar mirror) and the reflectivity criteria of the mirror are relaxed. Note that the study did not present any results on periodic textured rear mirrors. Nevertheless, it does not undermine this method since it leads to a better power conversion efficiency [START_REF] Eerden | Comprehensive analysis of photon dynamics in thin-film GaAs solar cells with planar and textured rear mirrors[END_REF]; [Watanabe et al. 2017]. Chen et al. evaluated the V oc loss at 40 meV for a Lambertian light trapping compared to a perfect flat mirror [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF].

Our novel structure is based on n-GaAs/p-Al x Ga 1-x As heterojunction and the GaAs absorber's thickness is 200 nm. The same 25 nm-thick AlInP window layer is maintained, with a DLARC of MgF 2 /Ta 2 O 5 (78/48 nm). The choice of the AlGaAs base thickness is discussed below, in view to increase the V oc while reducing parasitic losses in this layer.

Enhanced optical absorption

As previously discussed, a high refractive index contrast in the grating is essential to boost solar absorption through efficient diffraction. The easiest way to achieve this, is to use a wide-bandgap semiconductor for the grating and low RI dielectric materials or the spacer [START_REF] Pearce | Designing transparent nanophotonic gratings for ultra-thin solar cells[END_REF]]; [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF]]. Still, this configuration requires a trade-off between the parasitic absorption inside the grating and the enhanced absorption via the resonances activation in the active layer.

Wide-bandgap III-V semiconductor materials lattice-matched on GaAs are the natural choice for the grating (provided they can be patterned), with no deposition steps required other than the initial epitaxy [Eerden et al. 2020a]; [START_REF] Yang | Ultra-thin GaAs single-junction solar cells integrated with a reflective back scattering layer[END_REF]]. Since Al x Ga 1-x As is a part of the heterojunction, we conveniently keep the same material for the grating part, but with an optimized stoichiometry. In fact, when the aluminum content is varied, it will change the lattice constant negligibly, while strongly affecting the absorption response. Figure 2.17(a) shows a schematic of the solar cell design using an AlGaAs base and a nanostructured AlGaAs grating (with Al>40%). For the rest of the study, the rear side grating is optimized using the RCWA. The same parameters as before are varied: grating height, period, nanopillars width, refractive index of the dielectric and its thickness in between the grating and the mirror (red parameters in Figure 2.17(a)). These multiple possibilities were handled using a the previously mentioned genetic algorithm keeping the objective to maximize J th in the absorber. Although it does not prevent the problem of local maximum convergence, the risk is reduced thanks to the experience of expected results (gain from multiple runs) and the help of the initial population implementation. Note that J th represents the collected current from the GaAs absorber only, this assumption will be discussed later. For each numerical simulation presented in this section, the red parameters (back side parameters) are optimized according to the stack configuration unless stated otherwise. This means that each calculated J th is expected to be the best result.

The objective is to further reduce the parasitic absorptions (grating, mirror, window layer, BSF) as well as to increase the V oc . The new solar cell design consists of a DLARC, a n-AlInP window layer, a 200 nm-thick n-GaAs emitter, a p-Al x Ga 1-x As base, a Al x Ga 1-x As (x>0.4) nanostructured grating embedded in a low RI index dielectric and a flat silver mirror. The thickness and stoichiometry of the Al x Ga 1-x As base and back architecture (Al x Ga 1-x As grating and BSF) require a fine optimization to maximize the photogenerated current.

Aluminum content variation in Al x Ga 1-x As grating

As with the a-Si:H grating, there is parasitic absorption in the Al x Ga 1-x As nanopillars. But in this case, the aluminum composition in the nanostructures can be varied to reduce it. The base consists of a 100 nm-thick Al 0.4 Ga 0.6 As layer. Figure 2.17(b) shows the absorption spectra in the active layer and in the grating as a function of the aluminum content present in the nanostructures. Table 2.3 reports the theoretical collected current in the active layer, and the parasitic absorption in the 100 nm p-Al 0.4 Ga 0.6 As layer and in the grating.

As expected, the higher the aluminum content in the grating, the lower the parasitic absorption, resulting in a higher J th . For the structure with the highest Al concentration, J th reaches 27.95 mA/cm 2 . The grating with 40% Al has the highest parasitic absorption compared to the others. Surprisingly, the grating with 50% Al displays a slightly lower J th : the bandgap shift to higher energy is clearly visible, but this translates in a modest reduction of the parasitic losses in the Al x Ga 1-x As grating and a degraded light-trapping inside the GaAs absorber in the 650-900 nm spectral range. Overall, there is undoubtedly a decreased parasitic absorption inside the Al x Ga 1-x As grating when the aluminum content is increased from 40% to 70%. However, the net J th gain is limited to ∼0.12 mA/cm 2 due to concomitant effect on the light-trapping efficiency in the red part of the spectrum. Interestingly, the grating parameters do not vary much regardless of the Al content, which can be explained by the relatively similar shape of the Al x Ga 1-x As refractive index over the visible spectral range. Furthermore, we find that the parasitic absorption in the mirror remains small as expected (maximum value is dJ = 0.44 mA/cm 2 ).

These simulations highlight the key role of the p-Al x Ga 1-x As base in terms of parasitic absorption. To further increase the theoretical short-circuit current, two adjustments may be considered or combined: increasing the Al content in the base and/or reducing its thickness.

Alteration of the base's properties

Base thickness

Two constraints limit a major reduction of the base thickness. The first one is that the base with a limited doping level should be thick enough to facilitate lateral conduction [START_REF] Sotoodeh | Empirical low-field mobility model for III-V compounds applicable in device simulation codes[END_REF]]. Since we use localized contacts in our structure, this criterion is of utmost importance. The second is related to fabrication concerns. The nanostructuration of the grating will require an etching step (possibly dry etching for the best control) that can potentially damage the base, so that we may want to keep a relatively thick base to have some margin in case a "dead-zone" will be produced in the base upon dry etching of the grating.

Figure 2.18 shows the theoretical absorption spectra inside the GaAs, the p-Al 0.4 Ga 0.6 As base and the Al 0.6 Ga 0.4 As BSF for two different configurations: one with a 80 nm-thick base and a 20 nm-thick flat Al 0.6 Ga 0.4 As and one with a 50 nm-thick base and a 50 nm-thick flat Al 0.6 Ga 0.4 As. The base+BSF stack thickness is kept constant at 100 nm. This structure provides benefits on all fronts:

-It emphasizes the selective role of the BSF, which acts as a barrier to prevent minority carriers recombination, while facilitating the lateral conduction of the carriers -It provides a safeguard in the fabrication process -Higher Al content in the BSF minimizes the parasitic absorption, and therefore increases J th 300 400 500 600 700 800 900 1000 0.0 In the next simulations, the AlGaAs nanopillars are kept at 60% Al, a good compromise between the parasitic absorption reduction in the grating and its ease of oxidation. Besides, when Al content is increased, the series resistance in the layer increases as well, which is detrimental for the lateral conduction. To compensate for this effect the layer must be thicker. The J th exceeds 28 mA/cm 2 for both structures. Although no difference in J th is observed between the two structures, Configuration 1 (200/50/50 nm) offers a more suitable response to the absorption reduction in the base.

Base content

In this paragraph, we investigate how the Al content in the base layer affect its absorption properties and ultimately the absorption inside the GaAs absorber. Figure 2.19(a) shows the absorption in the GaAs as a function of the Al content from 40% to 70% in the base. For these simulations, we keep a 100 nm-thick Al x Ga 1-x As base (where the Al content is varied) and the grating's parameters are optimized for Al 0.6 Ga 0.4 As nanopillars.

Base material

We also considered the use of a GaInP base instead of Al x Ga 1-x As. With a GaInP base, a 80% Al content in the BSF is required to have a J th exceeding 28 mA/cm

Final considerations on the optoelectronic design

The whole stack harmony is based on a justified trade-off. We have considered a solar cell structure with a heterojunction to improve the V oc , and higher Al content in the base/BSF/ nanostructured grating to reduce their parasitic absorption while considering that too high Al contents can be detrimental to lateral conduction with lower carrier mobility [START_REF] Sotoodeh | Empirical low-field mobility model for III-V compounds applicable in device simulation codes[END_REF]] and the base thickness cannot be reduced too much due to high resistances. The new structure consists of an n-GaAs/p-Al 0.6 Ga 0.4 As heterojunction of 200/50 nm. This is followed by a 50 nm-thick Al 0.6 Ga 0.4 As BSF, and a 100 nm-thick periodic grating formed by Al 0.6 Ga 0.4 As nanopillars embedded in a SiO 2 layer (n=1.4, [START_REF] Bottein | Environment-controlled sol-gel soft-NIL processing for optimized titania, alumina, silica and yttria-zirconia imprinting at sub-micron dimensions[END_REF]; [START_REF] Dalstein | Nanoimprinted, Submicrometric, MOF-Based 2D Photonic Structures: Toward Easy Selective Vapors Sensing by a Smartphone Camera[END_REF]). Finally, we have added a transparent dielectric spacer (SiO 2 ) between the grating and the silver planar mirror to avoid parasitic absorption inside the mirror. Table 2.4 summarizes and compares the best simulations obtained so far, showing the corresponding theoretical current in the main absorber and rear layers. From an optical viewpoint, we can assess that for the BSF layer and for the grating, the higher the Al content, the lower the parasitic absorption. However, the difference is not large and lowering the absorption in these two layers lead to an increase of absorption in the base in any case. In addition, a (relatively) moderate Al concentration reduces stress and prevents oxidation, which justifies keeping the Al content at 60%. Keeping the nanostructured Al x Ga 1-x As layer at 100 nm provides efficient and robust light scattering. Since the role of the base is primarily electrical, care is taken to ensure that the three back layers discussed are sufficiently doped for majority carriers collection through localized contacts. Nevertheless, a significant part of the absorption occurs in the base and we have addressed this particular point in the discussion.

Base BSF Nanostruct. J th [mA/cm 2 ] dJ base [mA/cm 2 ] dJ BSF+ nanostruct. [mA/cm 2 ] J tot [mA/cm
Finally, in calculating the theoretical sort-circuit current, we have assumed that photons absorbed in the non-photoactive AlInP window layer are lost, but this assumption can be tempered based on the certified EQE of the initial ultrathin structure (see Figure 2.8(b)). Absorbed high-energy photons are partially collected, suggesting that the window layer also participates to the photogenerated current. Given this result, it is also reasonable to assume that some of the photons absorbed in the base of the heterostructure are also collected and contribute to the short-circuit current. Although a base with 60% Al increases J th by 0.4 mA/cm 2 in the main absorber (n-GaAs), a lower bandgap would be preferable. This would prevent the minority carriers generated in the base from going to the potentially damaged BSF/grating interface, as we aim to collect them. If we assume to collect the photogenerated carriers in the emitter and in the base, a J th ≈ 28.9 mA/cm 2 is expected for configurations 1 and 2. If we assume that it is possible to collect most of the minority carriers from the base, both heterostructures GaAs/Al 0.4 GaAs and GaAs/GaInP with Al 0.6 Ga 0.4 As nanostructures, have a remarkably high J th = 28.9 mA/cm 2 . The total thickness of the active layer would be 250 nm.

On the other hand, if no efficient minority carriers collection is expected, the heterostructure from configuration 2 is preferred due to its higher bandgap. No absorption in the base would results in suppressing or at least drastically reducing J 02 , leading to higher FF and V oc . The final trade-off of such structure would remain between the possible efficiency gain due to the increase of FF and V oc on the one hand and the increase of J sc on the other. Free-carrier absorption effect has been shown to strongly influence J sc and V oc for ultrathin GaAs solar cells using highly doped layers for light trapping structure. However, it starts to have a significant impact when the layers exceed 800 nm with doping in the order of 1 × 10 19 cm -3 [D' Rozario et al. 2022].

As an aside, and as observed with the a-Si:H/dielectric grating, the fraction of loss due to reflection is significantly reduced. The final stack is grown at the Fraunhofer ISE. The doping in all the layers is not discussed here and is left to the discretion of the Fraunhofer ISE.

Required accuracy in the experimental planarization of the nanograting

The analysis of the spacer between the grating and the planar mirror has been left aside, but it plays a large and influential role in the experimental steps. This spacer mitigates the parasitic absorption in the mirror and planarizes the periodic patterns on the rear side, resulting in a low roughness mirror. For most of the structures tested, the ideal spacer thickness is around h = 90 nm, slightly below the h NN = 100 nm. It may be necessary to slightly increase the spacer thickness, to finely planarize the nanostructure before the metallic mirror. To gain insight into its impact, J th in the main absorber layer for configuration 1 (see Fig. 2.18(b)) is plotted in Figure 2.21 as a function of the spacer thickness (in our case, a dielectric material of n=1.4). Each parameter is kept constant while only the spacer thickness is varied. Although there is no drastic change in the J th value for a 10 nm thicker spacer compared to the optimized value, this observation is no longer true for larger variations. The J th plummets with increasing spacer thickness. The small J th recovery at h = 340 nm is not sufficient to reach the optimized value. This parameter sensitivity is highlighted, and care should be taken during this experimental step. We go one step further and assume that the deposited mirror is not perfectly flat due to a non-ideal planarization. The effect of a rough mirror on the absorption in the GaAs layer (Figure 2.22) is simulated in a very coarse way by adding a higher dielectric thickness at the level of the grating nanopillars in AlGaAs. This thickness varies from 0 nm to 40 nm. The calculations are performed on n-GaAs/p-Al 0.4 Ga 0.6 As (200 nm-100 nm) with Al 0.6 Ga 0.4 As nanostructures. The effect on the resonances is minor, but they become increasingly pitted leading to an overall decrease in absorption. Although a small distortion does not lead to a consistent loss in the J th , a larger distortion results in an overall loss ∆J th =0.7 mA/cm 2 . However, such a large distortion is very unlikely using the planarization technique based on sol-gel derived SiO 2 that we have developed.

Conclusion of the section

In order to increase the V oc , the structure of the ultrathin GaAs cell has been modified into a heterostructure. Not only the parameters of the Al x Ga 1-x As/dielectric nanostructured back side have been optimized, but also the base properties were investigated. Reducing the Al x Ga 1-x As layer or increasing the Al content leads to less parasitic absorption. Several structures, balancing optical and electrical constraints, have been proposed. Assuming a carrier collection only in the 200 nm-thick GaAs, the best device (Configuration 2) has a J th equal to 28.5 mA/cm 2 . Considering the collection of photo-generated carriers in the base (250 nm-thick device) all the four sustained structures presented can reach a J th of about 28.9 mA/cm 2 .

Summary

We have conducted an extensive analysis of the absorption loss in the ultrathin record cell. Once the origin of the optical losses identified, two possible pathways toward more efficient devices were explored.

We have modelled ultrathin GaAs solar cells with a periodic 2-dimensional grating consisting of a-Si:H semiconductor nanopillars embedded in a low RI dielectric transparent material. A flat mirror is deposited afterwards which greatly reduces the losses in the silver compared to the nanostructured one of the original device. RCWA method coupled with a genetic algorithm is used to calculate the optical absorption and to optimize the geometry of the nanostructured back side. We have observed that Fabry-Perot and guided-mode resonances result in multiple overlapping absorption peaks, which broadened the absorption spectrum. The best result leads to a J th up to 28.2 mA/cm 2 , a gain of 2.6 mA/cm 2 compared to the original structure.

With the aim to further increase light absorption, but also to improve the V oc , we have designed a new heterostructure solar cell. Based on the same light-trapping strategy, but replacing the a-Si:H grating with a more transparent Al x Ga 1-x As material, we calculated a J th equal to 28.5 mA/cm 2 for a 200 nm-thick device. Taking into account the collection of photocarriers inside the base, several ultrathin configurations can reach a short-circuit current of J th of 28.9 mA/cm 2 (which corresponds to a 250 nm-thick absorber layer)

Both architectures have been designed keeping in mind the experimental constraints. The improved J th are plotted in the [START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. Short-circuit current of ultrathin GaAs solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models previously described. The values mentioned in the manuscript are in red and labeled. a. [START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]], b. [START_REF] Liu | Surface plasmon enhanced GaAs thin film solar cells[END_REF], c. [START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF]], d. Camarillo Abad et al. [Camarillo Abad et al. 2020], e. Camarillo Abad et al. [Camarillo Abad et al. 2022], f. [START_REF] Buencuerpo | Engineering the reciprocal space for ultrathin GaAs solar cells[END_REF], g. [START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF]], h. van Eerden et al. [Eerden et al. 2020a], i. [START_REF] Lee | High Performance Ultrathin GaAs Solar Cells Enabled with Heterogeneously Integrated Dielectric Periodic Nanostructures[END_REF] , j. Chen et al. [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF], k. [START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF], l. [START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF], m. [START_REF] Massiot | Metal Nanogrid for Broadband Multiresonant Light-Harvesting in Ultrathin GaAs Layers[END_REF]. The blue star is obtained by modifying the back architecture of the record cells, and the purple star by transforming the homojunction to a heterojunction. Chapter 2 paved the way for the realization of more efficient ultrathin GaAs cells. Using the same 205 nm layer stack as the original record-breaking GaAs solar cells [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]], a J sc of 28.2 mA/cm 2 was calculated with a new nanostructured back side based on amorphous silicon semiconductor. In this chapter, we present the fabrication and the optimization of these specific cells.

In addition to the J sc enhancement, we propose a model to improve the V oc and the FF by modifying the front contact grid design and increasing the size of the solar cells (Section 3.1). Then, in Section 3.2, we detail the experimental steps to fabricate ultrathin GaAs cells with a nanostructured a-Si:H/SiO 2 grating. First results are reported in Section 3.3, and their performances are compared to experimental devices with different light-trapping structures. Finally, we discuss the J sc losses using simulations and evaluate experimentally the effect of increasing the size of the solar cells.

Optoelectronic improvement

In this section we model the resistive losses due to the front contact grid and propose a new design compared to the ultrathin record cell to improve the performance.

Electronic characteristics and edges recombination

Recombination current at the perimeters degrades the performance of GaAs solar cells [START_REF] Demoulin | Influence of Perimeter Recombination on High-Efficiency GaAs p/n Heteroface Solar Cells[END_REF], and efficiency loss in small cells has been linked to the unpassivated edges [START_REF] Ochoa | 3-D modeling of perimeter recombination in GaAs diodes and its influence on concentrator solar cells[END_REF]]; [START_REF] Sheldon | GaAs Passivation with Trioctylphosphine Sulfide for Enhanced Solar Cell Efficiency and Durability[END_REF]]. [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF] demonstrated that the V oc decreases with the cell size and is correlated with an increased dark current density J 02 . For a cell of 1 cm 2 surface area, the J 02 is reduced to about 1 × 10 -9 mA/cm 2 , and when the edge recombinations are suppressed, the V oc and the FF are expected to increase up to 1.045 V and 82.6%, respectively [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]].

Metallic contacts losses analysis

Front contact resistive losses

The front contact grid also contributes significantly to losses in ultrathin GaAs devices. Typical front contact designs for GaAs cells are sketched in Figure 3.1(a,c) [START_REF] Flat | Triple-junction solar cells with 39.5% terrestrial and 34.2% space efficiency enabled by thick quantum well superlattices[END_REF]]; [START_REF] Cariou | III-V-on-silicon solar cells reaching 33% photoconversion efficiency in two-terminal configuration[END_REF]]. Lateral resistive losses can occur according to the geometry of the front grid, which requires optimization. To calculate the resistive losses, we developed a theoretical model based on simulations reported in the literature [START_REF] Honsberg | Photovoltaics Education Website Cited on[END_REF]] * , [START_REF] Liu | Optimization of grid design for solar cells[END_REF]]; [START_REF] Mäckel | Analytical models for the series resistance of selective emitters in silicon solar cells including the effect of busbars[END_REF]]. We consider a simple structure of four layers extracted from the ultrathin GaAs cell, whose properties are summarized in Table 3.1. The limitation due to the top contacts can be categorized into four main resistive losses:

-The grid shadowing -The lateral conduction in the emitter -The transport from the finger to the busbar -The contact resistance

The complexity of the optimization depends on the required trade-offs between all the parameters involved. The equations on which the theoretical model is based are developed in the Appendix A. [START_REF] Sotoodeh | Empirical low-field mobility model for III-V compounds applicable in device simulation codes[END_REF]].

The power losses calculated with equations A.1, A.2, A.4 and A.5 are plotted as a function of the finger spacing S in Figure 3. 1(b,d). It shows the behavior of each loss for two different front grid designs. The black line corresponds to the total loss: the sum of each loss contribution. Figure 3.1(b) shows the losses corresponding to the front grid that was used for the record ultrathin GaAs cell shown in (a). It shows that the shadowing (yellow line) is the main source of losses, which is proportional to the ratio of covered and uncovered area.

The minimum of the total loss is about 6% and almost constant in the range S=[600 µm-900 µm]. This result has been observed experimentally with similar J sc regardless of the finger spacing. The relative losses associated with the lateral conduction in the emitter (blue line) and the transport in the finger (green line) remain small compared to the shading. However, beyond S = 1000 nm, they influence the total loss, which starts to rise. The optimized grid design consists of a surface area of 1 cm 2 with modified parameters to reduce the shading (see Figure 3.1(c)). The widths of the fingers and the two busbars are set to 5 µm and 30 µm respectively, and the pads cover about 0.3% of the surface. The finger width (w f ) is kept at 5 µm, because a smaller value may be challenging from the process perspective. Figure 3.1(d) shows that the total loss is reduced to a little more than 3% for a finger spacing of 400 µm. The shadowing loss has been greatly reduced and is similar to the resistive losses in the fingers. The cell size has a strong influence on the resistive losses in the finger and to compensate for it, the contact thickness is increased to 2 µm. Under the same conditions, for a 220 nm-thick contact, the minimum total loss is equal to 7.7% for S = 140 µm. Note that for both configurations, the effect of the contact losses (red line) is small.

Conclusion of the section

It has been shown that increasing the cell size and optimizing the front grid design could increase the FF and V oc of ultrathin GaAs solar cells. The front grid design of GaAs ultrathin solar cells has been improved using a simple model based on four main losses. Taking into account the edge recombination losses, a new front grid architecture is presented for a 1 cm 2 cell. The front total losses of the original design are reduced by half.

Fabrication process

In this section, we describe the detailed fabrication process of transferred ultrathin GaAs solar cells. The as-grown stack is described in Table 2.2. We introduce the fabrication steps of a nanostructured a-Si:H/SiO 2 back side followed by a planarization (spacing layer) and an Ag planar mirror. Figure 3.2 shows a schematic of the solar cell fabrication sequence.

Localized back contacts

A direct correlation between the back reflectance and cell performance has been demonstrated [START_REF] Gruginskie | Limiting mechanisms for photon recycling in thin-film GaAs solar cells[END_REF]. Therefore, to overcome the trade-off between high reflectivity and good ohmic contacts we use localized ohmic contacts that cover about 1% of the back surface, while the other 99% is covered by the silver mirror [START_REF] Vandamme | Ultrathin GaAs solar cells with a silver back mirror[END_REF]]; [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]; [START_REF] Hwang | Bandgap grading and Al0.3Ga0.7As heterojunction emitter for highly efficient GaAs-based solar cells[END_REF]]. The localized contacts are defined by photolithography. The mask consists of 5 x 5 µm 2 squares regularly spaced with a period of 50 µm in both the in-plane x-and ydirection. Ti/Au (20/200 nm) contacts are deposited by electron-beam assisted evaporation after a deoxidation of the p-GaAs contact layer in dilute HCl (37%) solution (HCl:H 2 O 1:3). The contact layer is kept under the metal contacts and removed elsewhere by selective chemical etching. The solution consists of a mixture of citric acid at 1 g/L and hydrogen peroxide (30%) at a volume ratio of 5:1, and the etching naturally stops at the AlGaAs back surface field layer. Figure 3.3(a) shows a microscope image of the metallized back contact on the p+GaAs layer.

Nanostructured a-Si:H layer

Immediately after the contact layer removal and deoxidation step in dilute HCl, 100 nm of hydrogenated amorphous silicon (a-Si:H) is deposited on the entire back surface by Plasma-Enhanced Chemical Vapor Deposition (PECVD) at the LPICM [START_REF] Desthieux | Development and characterization of Fired Passivating Contacts for p-type silicon solar cells fabrication[END_REF]]; [START_REF] Yu | Understanding light harvesting in radial junction amorphous silicon thin film solar cells[END_REF]]. The deposition rate was evaluated at 0.097 Å/s, resulting in a deposition time of about 16 min. The refractive index of this layer is determined by ellipsometry and compared with the one used for the simulations in the previous chapter. We find a very close agreement between the two data sets, confirming both our hypothesis and the required quality of the a-Si layer.

To improve the adhesion of an electron-sensitive resist on the a-Si:H, a 5 nm SiO 2 layer is sputtered on the a-Si:H surface. The sample is then coated with a negative resist: ma-N 2401, and patterned in an electron-beam lithography (EBL) machine operating at 100 kV. The patterning of a 1.2x1.2 cm 2 surface, exposed at 20 nA, with squares spaced by a period of 750 nm takes no more than 1.5 hours. The size of the squares depends on the injected electron dose, while the distance between them is accurately controlled by the e-beam. After EBL, the pattern is developed in a ma-D 525 solution for 30 s, and transferred to a-Si:H by reactive ion etching. A mixture of SF 6 and CHF 3 gases (flow rate ratio of 4: 24) is used with a bias voltage of 240 V. The etching time is about 9 min to completely remove the a-Si:H layer down to the AlGaAs layer. The resist is then removed with acetone, and to avoid any residual resist particles, 2 min of plasma air is carried out. Figure 3.3(b) shows the SEM image of the sample after the nanostructuration of the a-Si:H layer. The patterning is not perfectly homogeneous over the surface, but there is enough nanopillars close to each other to ensure that the grating diffraction will take place.

Planarization by SiO 2 and Ag mirror

A sol-gel of silicon oxide is prepared from a solution EtoH/HCl (0.1M)/MTEOS/TEOS at a ratio of 3.5/0.5/0.5/0.5 g. A CTAB surfactant was added in order to improve the wettability of the substrate (0.245 g). The prepared solution is filtered through a 0.45 µm PTFE membrane to remove the aggregates. The sol-gel solution is spin-coated onto the nanostructures at 1500 rpm for 10 s. To stabilize the sol-gel film, a soft annealing at about 130°C is performed to remove the volatile solvents. Complete condensation of the inorganic network at 350°C for 30 s causes the film to shrink [START_REF] Bottein | Environment-controlled sol-gel soft-NIL processing for optimized titania, alumina, silica and yttria-zirconia imprinting at sub-micron dimensions[END_REF]. This results in a ∼280 nm-thick porous SiO 2 layer, or a layer of 180 nm above the nanostructures, with a refractive index comprised between 1.23 and 1.21 over the spectral range(see Figure 3.4(a,b)). Subsequently, the planarized surface was protected with a photoresist mask of interest, opening only the area of the localized ohmic contacts. To avoid any detrimental contact between the metallic mirror and the AlGaAs layer, the opening is inscribed in the localized contact (see Figure 3.4(c)). The nanostructured a-Si:H and the SiO 2 layer above the contacts were etched by RIE (SF 6 :CHF 3 , flow rate 4:16) for 15 min 30 s. The photoresist was removed in acetone. The following metal stack Ti/Ag/Ti/Au/Ti (1/200/10/100/20 nm) was deposited by electron-beam assisted evaporation using a rotating stage. The thin Ti layer, while detrimental to the high reflectivity of Ag, is necessary to promote adhesion between SiO 2 and the metallic mirror. Thus, a nanostructured a-Si:H/SiO 2 is obtained, followed by a planar Ag mirror that also acts as a back electrode.

Bonding and substrate removal

The Ag mirror-side of the sample was bonded to a glass host substrate using OrmoStamp®, a UV-reticulated flexible hybrid inorganic/organic polymer. A glass substrate was chosen because it is transparent to UV light. Both the glass and the sample go through a plasma air and then the glass is deposited on a hot plate at 120°C. OrmoStamp® is spin-coated on the Ag-side of the sample at 3000 rpm for 30 s and then applied to the glass substrate. The sample is bonded to the glass after the reticulation of OrmoStamp® under UV light for 20 min.

The GaAs substrate removal is carried out by chemical etching. Most of the etching solutions are a mixture of an oxidizer agent and a chemical etchant (acids or bases) diluted in deionized water. First, the material is being oxidized (the most commonly used oxidizer is the hydrogen peroxide (H 2 O 2 )). Then the oxidized compounds are dissolved in the solution. On the one hand, the etching rates will depend on the oxidation speed or the dissolution reaction, on the other hand it also depends on the crystallographic orientation of the semiconductor. The wet chemical etching of III-V compounds has been widely reviewed in the literature [START_REF] Clawson | III-V semiconductor chemical etching[END_REF]], and Table 3.2 offers a (non-exhaustive) summary of wet chemical etching behavior for four III-V compounds and alloys. We note that As-and P-based III-V semiconductors act differently. The GaAs substrate was etched in two different solutions. The first solution NH 4 OH (30%):H 2 O 2 (30%) in proportion 1:4 achieves a fast chemical etching (8 to 15 µm per minute depending on the solution temperature). It is selective over AlGaAs, but the selectivity is not high enough, so that 300 nm of Al 0.85 Ga 0.15 will be etched within a few minutes. An image of the etched GaAs substrate is shown in Figure 3.4(d). The central area with purple and green colors corresponds to the AlGaAs etch stop layer. The solution is changed to a high-selectivity (citric acid):H 2 O 2 solution (5:1). The monohydrous citric acid powder is dissolved in water at a weight/volume ratio of 1 g per 1 ml. The sample can stay in the solution for over 3 hours while it homogenizes the sample surface by removing the remaining GaAs. While the first solution requires constant stirring to maintain a constant pH at the sample surface, the citric acid based etch is not affected by the stirring at all. Finally, the etch-stop layer AlGaAs is removed in pure HCl(37%). After this step, we have a transferred cell, since the stacking order is reversed. The substrate removal takes about 4 hours in total and would benefit from optimization. The sensitive phase where the etch stop layer appears could be avoided with a different strategy, and the reaction rate of the second solution could be accelerated by heating up.

Front contacts and mesa etching

As previously discussed, the thickness of the fingers and busbar plays a fundamental role in power losses. The target thickness is more than 1 µm thick in order to reduce the resistance and improve charge collection. We have developed two different lithography processes for thick front contacts, and the detailed steps are developed in Appendix A.1.2. The front contacts are spaced by 600 or 800 µm, and multilayers of Ni/Au/Ge/Au/Ni/Au (4/10/60/110/10/1000 nm) are deposited on n+GaAs. There is no post thermal annealing, as the metal diffusion would degrade the ultrathin GaAs absorber.

The final solar cells of 1x1, 2x2 and 3x3 mm 2 are delimited with photolithography mask and mesa etching. The AZ10xt positive photoresist is chosen because of its nominal thickness of 7 µm. It allows a good coverage of the entire cell with thick contacts. The sample is then subjected to a series of wet chemical etchings down to the a-Si:H/SiO 2 surface. The n+GaAs contact layer and the InGaAs capping layer are removed in a H 3 PO 4 :H 2 O 2 :H 2 O, with a volume ratio of 3:1:40. It stops naturally at the AlInP layer, so the solution is changed to pure HCl, and the window layer is dissolved within a few seconds. We remove the active GaAs layer and the BSF AlGaAs layer using the same phosphoric acid based solution as before for several minutes. After the resist removal, the n-GaAs contact layer and the InGaAs capping layer between the fingers of the front metal electrode are selectively removed with citric acid and hydrogen peroxide (5:1). This solution is preferred because phosphoric acid could create pinholes in the layer and damage the cell. Figure 3.5(a) shows a picture of the sample with a nanostructured back side (C2853b-12), and Figure 3.5(b) shows the labeling of the solar cells.

After measuring the solar cell performance, a MgF 2 /TiO 2 (98/47 nm) double-layer antireflection coating can be deposited.

Conclusion of the section

We have described the main steps to fabricate ultrathin GaAs solar cells. At the back side we added a nanostructured grating made of a high refractive index semiconductor a-Si:H with a SiO 2 sol-gel of refractive index near 1.2. Then a planar Ag mirror was introduced due to its high optical reflectivity. We thus combined a-Si:H/SiO 2 grating and Ag mirror with localized ohmic contact to efficiently collect photo-generated charge carriers. A sample with several solar cells and low defects was successfully fabricated (see Figure 3.5(a)). Note that, due to time constraints, the new front grid design for large cells was not implemented and the nanopatterning was performed by EBL instead of nanoimprint lithography

Characterization of ultrathin GaAs solar cell

Preliminary results

In this section, we present the first results obtained for ultrathin solar cells with a nanostructured a-Si:H/SiO 2 back side, followed by a flat Ag mirror. Note that no ARC was deposited, yet. Figure 3.6(a,b) shows the JV curves (under one sun illumination and in the dark) and the EQE of an ultrathin GaAs solar cell. From JV measurements we find for this cell a J sc =14.91 mA/cm 2 , a V oc =1.017 V and a FF= 0.786. However, from the EQE measurement and without shading, the extracted J sc is equal to 17.39 mA/cm 2 , resulting in a solar cell with an efficiency of 13.9%. This difference in J sc of almost 2.5 mA/cm 2 can be partially justified by the front grid shadowing, but not only. An underestimation of the J sc , intrinsic to the solar simulator used to characterize the cells, may also be a reason. The multiple resonances and the large absorption peak in the range of 800-900 nm confirm the light trapping effect and the positive impact of the nanostructured back side on the absorption of the ultrathin device. A summary of the photovoltaic parameters (J sc , V oc , FF, and 1 sun conversion efficiency) of all the cells is shown in Figure 3.7(a-d) as a function of the cells size. For simplicity, the cell size is denoted A, B, and C for 3x3mm 2 , 2x2mm 2 , and 1x1mm 2 , respectively.

Note that the J sc presented here is measured under a solar simulator and normalized to the total area of the solar cells, and may be underestimated. The best measured current density reaches 16.4 mA/cm 2 for the largest cells. Overall, the smallest cells have the lowest J sc values, but this disparity is strongly related to the front grid shadowing which is about 8% compared to 6% for the A and B groups. Moreover, this observation is corroborated with cells of size A and B because the cells with the lowest J sc for these groups are the ones with the highest shading (see Figure 3.5(b): the cells have different finger spacing). Very good V oc values are obtained, and as expected, edge recombinations play an important role in the performance: the V oc decreases with the cell size and is correlated with an increase of the saturation current density J 02 [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]. The extraction of the parameters J 02 and R s is described in Appendix A.1.3. On the other hand, the FF is slightly higher for C cells. x -7

x 3 mm² 2 x 2 mm² 1 x 1 mm² The very first results obtained for ultrathin GaAs solar cells with a nanostructured back side, a-Si:H/SiO 2 grating and a flat Ag, are presented. We have demonstrated our proofof-concept, by fabricating working devices with promising performances, and efficiencies between 11% and 13%. Despite encouraging FF and V oc , the J sc values remain slightly lower than predicted by the simulations. Without a double layer anti-reflection coating, we expected a J sc close to 20.6 mA/cm 2 for a perfectly fabricated device. This observation is discussed in the next section.

3 x 3 mm² 2 x 2 mm² 1 x 1 mm² (a) (c) (b) (d) (e) (f) (A) (B) (C) (A) (B) (C) (A) (B) (C) (A) (B) (C) (A) (B) (C) (A) (B) (C)

Performance assessment

In this section, we evaluate the performance of our new back-side architecture of ultrathin GaAs cells, focusing on the J sc parameter. We compare these cells to other experimental devices with the same 205 nm-thick GaAs layer but different light trapping strategies. We then analyze their J sc value and we present their expected performance after the deposition of a double layer anti-reflection coating (DLARC).

Spectral response comparison

To highlight the efficient absorption of the fabricated nanostructured cells, and to illustrate the absorption models (single-pass, double-pass and light-trapping) developed in the previous chapter, we compare the best EQE of four different ultrathin GaAs cell structures. show the results for a solar cell on a GaAs substrate, and for a cell transferred on a flat Ag mirror respectively. Figure 3.8(c) compares the EQE measurements for two cells with a nanostructured rear side. On the one hand, the device with a nanostructured TiO 2 /Ag mirror (blue line, Structure T), corresponds to the latest record cell for 205 nm-thick GaAs devices to date (without the DLARC). The data are taken from the PhD thesis of H.L. Chen [START_REF] Chen | Ultrathin and nanowire-based GaAs solar cells[END_REF]]. On the other hand, we show the device with a nanostructured a-Si:H/SiO 2 grating and a flat Ag mirror (red line, Structure S).

The increased absorption is evidenced with each new structure. The addition of a planar mirror at the back side almost doubles the J sc value: sharp Fabry-Perot resonances lead to higher absorption in the GaAs active layer (Figure 3.8(b)). The addition of a nanostructured back side leads to a strong absorption improvement in the low absorption region (long wavelengths), and overall the multi-resonant response enhances the photocurrent. While the structure T slightly outperforms structure S in terms of J sc , the spectral response of both structures is very close. The EQE of the structure S shows two small drops compared to the structure T. The first one, between 520 nm and 600 nm is related to the parasitic absorption in the amorphous silicon, and the second one is between 700 nm and 750 nm. The last major difference is in the long wavelength part. Structure S has a higher but sharper EQE peak than structure T, as predicted by the simulation (see Figure 2.12). This results in a reduced sub-bandgap absorption. Figure 3.8(c) confirms the correct introduction of light trapping via the a-Si:H/SiO 2 grating and the Ag flat mirror.

J sc investigation

Based on the simulations, there is still a J sc loss of about 2 mA/cm 2 (J th predicted at 20.65 mA/cm 2 , and best J sc determined by EQE is 18.47 mA/cm 2 ). However, these predictions do not take into account the real shape of the a-Si:H/SiO 2 grating, and the fabrication constraints. We introduce now several modifications to our original simulation model. They are listed in Table 3 Each parameter is modified based on control measurements during the fabrication, except for the SiO 2 spacer thickness (the layer in between the grating and the flat mirror). This layer could not be measured and it was varied in the simulation to better match the experimental EQE. Initially optimized at 180 nm, its thickness was varied between 80 nm and 400 nm, and although we fixed it at 280 nm for the final calculation, there remains a large uncertainty in its thickness. Note that an error of 100 nm for this spin coated sol-gel is a very likely, because this step is not yet entirely optimized. 3.3). (a) with no DLARC, (b) with a DLARC. Equivalent currents at 1 sun (J sc for EQE, J th for absorption in the GaAs absorber) are given.

In Figure 3.9(a), we compare the EQE with the calculated absorption in GaAs using the modified parameters to better describe the experimental conditions. Overall, all the resonance features are in agreement with the EQE measurement, although there is a clear shift between the two optical responses (∼30 nm). The two sharp resonances in the red region are not resolved in the experiment (λ= 760 nm). This might be related to the diffraction efficiency, which depends strongly on the geometry of the grating with a sharp square shape assumed in the simulation. We attribute the EQE drop between 700 nm and 750 nm to the thickness of the SiO 2 spacer, since we do not see it when the structure is calculated with a 180 nm-thick spacer. The calculated absorption results in a slightly lower J th of 18.25 mA/cm 2 than the best experimental J sc deduced from EQE (18.47 mA/cm 2 ). Figure 3.9(b), shows the theoretical EQE after a DLARC deposition (98 nm/47 nm MgF 2 / TiO 2 ). We observe that the maximum absorption at the resonant wavelength is reduced, but instead a more efficient broadband multi-resonant absorption is obtained. The current density is estimated to be 24 mA/cm 2 , which would result in devices with efficiencies higher than 19%.

With the optical simulations we managed to identify the parameter responsible for most of the loss. Taking into account the experimental parameters values (updated structure in Table 3.3) but keeping the spacer thickness (h spacer ) at 180 nm, the absorption losses in the GaAs layer results in a J sc loss of 0.6 mA/cm 2 . However, simulating again the updated structure but with a 280 nm-thick spacer it results in a J sc loss of 1.8 mA/cm 2 . By simply adjusting the spacer thickness, the J sc could easily reach 27 mA/cm 2 with a DLARC. Note that the loss contribution of the 1 nm titanium layer is 0.45 mA/cm 2 , (when the DLARC is deposited), and this should also be addressed in the future.

We have evidenced the multi-resonant light-trapping strategy of ultrathin GaAs solar cells with a nanostructured back side consisting of an a-Si:H/SiO 2 grating and a flat mirror. We demonstrated its superior absorption compared to ultrathin cells deposited on a flat mirror alone, with an enhanced absorption at long wavelengths. To explain the lower than expected value of J sc , we have simulated again the absorption in the GaAs layer, with a structure closer to reality. With a DLARC deposition a J sc of 24 mA/cm 2 is expected, resulting in an efficiency higher than 19% for the best cells. From optical simulations, we have identified the parameter responsible for most of the loss: the thickness of the SiO 2 spacer. By simply adjusting the the SiO 2 spacer thickness, the J sc could be increased by almost 3 mA/cm 2 , leading to a new record for a 205 nm-thick GaAs solar cell.

Impact of the cell size and its front grid design

In the previous subsection, we used an early design and defined cells of maximum size 3x3 mm 2 . However, we have demonstrated in section 3.1 that increasing the cell size and modifying the front grid design can reduce relative losses and have a positive impact on the V oc and the FF. The experimental results of this study are presented here. This optimization has been carried out on ultrathin GaAs cells transferred on a flat Ag mirror and the different parameters of each group, denoted as A, B, C and D are presented in Table 3 The cells performances are summarized in Figure 3.11. A, B and C cells are plotted as references, to evaluate the positive effect of the grid modification and the size increase. The J sc presented is measured under a solar simulator and normalized to the total area of the solar cells. It might be underestimated, but does not prevent a quantitative comparison between the cells. D1 and D2 cells offer a slightly better J sc , which can be directly attributed to a lower percentage of shading. The V oc of D1 cells outperforms that of A, B and C cells, reaching values as high as 1.032 V, confirming the detrimental impact of the edge recombinations for small cells. Rs (ohm.cm²) However, this trend is not followed by D2 and D3 cells. On the one hand, for D2 cells the degradation during the fabrication justify these results and the heterogeneity between the cells.

A B C D1 D2 D3 A B C D1 D2 D3 A B C D1 D2 D3 A B C D1 D2 D3 A B D1 D2 C D3 A B D1 D2 (a) (c) (b) (d) (e)
On the other hand, we attribute the low performance of D3 cells to the front contacts that are too thin (200 nm) and generate high ohmic losses. This emphasizes the importance of thick contacts as cell size increases. Unfortunately, larger cells have very low FF regardless of the design. High serie resistance up to two times higher than for small cells may be partly responsible for this response. Greater deoxidation prior to metal deposition may help to minimize the resistance between the metal and the GaAs layer.

We have successfully increased the size of ultrathin GaAs solar cells: efficiencies similar to smaller cells are presented. However, the degradation of the FF is limiting, in view of the J sc and V oc we can expect better performance.

We have experimentally demonstrated the benefit of increasing the size of solar cells with higher J sc and V oc . The increase in V oc proves that we can overcome the losses due to edge recombinations with larger cells. The importance of the front contact thickness has also been emphasized.

Conclusion of the section

Ultrathin GaAs solar cells with a nanostructured a-Si:H/SiO 2 grating and a flat mirror have been successfully fabricated. The light trapping strategy has been highlighted with EQE measurements showing a strong absorption enhancement in the high part of the spectrum. After a DLARC deposition, efficiencies up to 19% are expected. Moreover, an optimized SiO 2 spacer thickness would result in a J sc of 27 mA/cm 2 (with DLARC). Increasing the cell size has experimentally confirmed a way to further increase the V oc . 68 3.4. Summary

Summary

Based on the model developed in Chapter 2, we have successfully fabricated ultrathin (205 nm) GaAs solar cells with a periodic 2-dimensional nanostructured grating made of a-Si:H/SiO 2 on a planar silver mirror. These working devices have very encouraging V oc and FF, and the multi-resonant light trapping is highlighted. It is responsible for the absorption enhancement in the high wavelength region. By adapting the parameters to simulate the real structure, we expect a J sc close to 24 mA/cm 2 after a DLARC deposition. We evidence that modifying only one parameter during the fabrication (the thickness of the SiO 2 spacer) would result in a J sc of more than 27 mA/cm 2 . Finally, we have increased the size of an ultrathin solar cell on a planar silver mirror and demonstrated an improvement in V oc . Combining the positive effects of the increased cell size and the nanostructured back side would result in devices with efficiencies higher than 22% (considering a J sc = 27 mA/cm 2 , V oc = 1.032 V, and FF= 0.80 (highest V oc and FF obtained so far with our cells)). With the idea of transferring the light trapping concepts developed for ultrathin GaAs solar cells we now turn our attention to CdTe-based solar cells.

In Chapter 1 we described the CdTe semiconductor with its nearly ideal bandgap of 1.5 eV, and sufficient charge carrier lifetime and mobility as a promising material for solar cells. However, this technology still suffer from poor V oc due to the nature of the polycrystalline thin films. Finite grain sizes and grain boundaries due to fast deposition, bulk defects, and interfaces are some of the factors limiting performances. While the role of interfaces will be strongly emphasized in the next chapter, we will focus here on the material characterization. Recently, CdTe-based solar cells have been improved by adding selenium alloys: the J sc is improved and surprisingly no V oc loss is reported despite the bandgap lowering. We describe the positive influence and passivation role of Se in Section 4.1. Yet, a complete and quantitative study of the effect of the selenium distribution within the absorber is still missing. In this chapter, we propose a quantitative characterization of the CdSe x Te 1-x material by cathodoluminescence to address this particular aspect. We conduct a systematic study on a batch of CdSe x Te 1-x samples with several selenium contents at different temperatures, and we quantify the passivation of Se and the diffusion of Se atoms along grain boundaries and in grain interiors.

State-of-the-art of CdSe x Te 1-x material for photovoltaics

Over the last ten years, despite an efficiency hindered by the presence of defects in grain interiors and at grain boundaries (GB), CdTe solar cells improved their efficiency up to 22.1% [Green et al. 2022b]. The latter record was driven by the addition of selenium composition grading in the layer, between the front buffer layer and the CdTe. It has been shown that the Se alloy in CdTe significantly improves the short-circuit current [Munshi et al. 2017]. Due to bowing effects, which will be discussed in more details in Section 4.3, CdSe x Te 1-x alloys exhibit reduced bandgaps (down to 1.4 eV) and this results in an enhanced photocurrent in the red part of the spectrum. The EQE measurements of CdTe-based solar cells with Se compositions ranging from 0% to 40% as shown in Figure 4.1(a), indicate a higher photocurrent with an increased absorption from 800 nm as the Se content increases.

A recent study has demonstrated the key role of CdTe in a polycrystalline CdSeTe/CdTe graded bilayer, the CdTe acts as an electron reflector layer improving the V oc [Shah et al. 2021a]. Then, instead of replacing the CdTe layer with uniform CdSeTe, a gradient of selenium is added. It can be introduced into CdTe-based solar cells in several ways. A CdSe x Te 1-x compound can be deposited by CSS prior to the CdTe growth [START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF], or a thin CdTe layer is annealed in a selenium atmosphere followed by a thicker CdTe layer deposition [START_REF] Artegiani | A new method for CdSexTe1-x band grading for high efficiency thin-absorber CdTe solar cells[END_REF]. So far, the highest J sc (31 mA/cm 2 ) was achieved by the NREL, depositing CdSe and CdTe layers by thermal evaporation [START_REF] Ablekim | Thin-Film Solar Cells with 19% Efficiency by Thermal Evaporation of CdSe and CdTe[END_REF]] (see Chapter 5). The majority of the Se diffusion within the CdTe film occurs during CdCl 2 annealing, resulting in a CdSe x Te 1-x /CdTe structure. Typically the Se content is comprised between 0% and 40%, because higher concentrations may promote a phase change from zinc-blende to photo-inactive wurtzite, in the sense that we have a carrier collection deficit [Poplawsky et al. 2016a]. Although the cell J sc increases due to Se incorporation, this effect alone does not explain such a high performance improvement. In addition to higher photocurrent, the voltage deficit is maintained or even improved despite a lower bandgap [Munshi et al. 2017]; [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]]; [START_REF] Bothwell | Characterization of thin CdTe solar cells with a CdSeTe front layer[END_REF]]. One explanation is that higher level of Se content seems to create a grading in electron affinity, driving electrons to the front and holes to the back [START_REF] Fiducia | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF]. However, numerous studies strongly correlate this effect with superior minority carrier lifetime in CdSeTe thin films [START_REF] Kuciauskas | Radiative Efficiency and Charge-Carrier Lifetimes and Diffusion Length in Polycrystalline CdSeTe Heterostructures[END_REF]] compared to CdTe [START_REF] Metzger | CdCl 2 treatment, S diffusion, and recombination in polycrystalline CdTe[END_REF]]. The enhanced lifetime is evaluated from a few ns [START_REF] Metzger | Time-resolved photoluminescence studies of CdTe solar cells[END_REF]]; [START_REF] Jensen | Long carrier lifetimes in large-grain polycrystalline CdTe without CdCl2[END_REF]] up to hundreds of nanoseconds (exceeding 200 ns) in CdSeTe cells [START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]]; [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF]]; [START_REF] Ablekim | Exceeding 200 ns Lifetimes in Polycrystalline CdTe Solar Cells[END_REF]]; [START_REF] Li | Low-temperature and effective ex situ group V doping for efficient polycrystalline CdSeTe solar cells[END_REF]]. In the following, we report experimental characterization and numerical results performed by the scientific community on CdSe x Te 1-x thin films, which have largely contributed to a better understanding of the effect of the selenium on cell performance.

The excellent optoelectronic properties of polycrystalline CdSeTe have been strongly correlated to the defects passivation in the CdTe bulk by the Se alloy. In 2019, Zheng et al. [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF]] and Fiducia et al. [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]] have performed cathodoluminescence (CL) experiments on beveled CdSeTe/CdTe samples, which allow probing optical properties along the sample thickness and direct comparison between the CdSeTe and CdTe regions. Zheng et al. characterized the chemical changes throughout the devices with Auger-electron spectroscopy and corroborated these compositional changes with CL showing an enhanced CL intensity in the CdSeTe region relative to CdTe [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF]]. Fiducia et al. with a complementary nano secondary ion mass spectrometry (SIMS) map at the same location as the CL map, evidence the correlation between Se content, CL intensity at the nanoscale and longer diffusion lengths. This increase in the radiative efficiency indicates a passivation of the bulk defects by the selenium [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]. They also showed the expected red-shift of the CL maximum intensity peak in the Se rich region, indicating a lower bandgap. They argue that the selenium diffuses into the CdTe mostly during the CdCl 2 annealing step. After evidencing with SIMS that the selenium diffuses into the CdTe GI/GB mainly along CdSeTe GBs [START_REF] Fiducia | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF]]; [Shah et al. 2021b], experimental evidence of Se reducing harmful non-radiative recombination at grain boundaries has been confirmed using a cooled STEM-CL system [START_REF] Fiducia | Selenium passivates grain boundaries in alloyed CdTe solar cells[END_REF]]. Density function theory modeling has been used to investigate the effect of Se on the electronic properties of CdTe GB. They propose that the co-doping of selenium and chlorine at a CdTe grain boundary reduces the density of harmful mid-gap defect states [START_REF] Guo | Effect of selenium and chlorine co-passivation in polycrystalline CdSeTe devices[END_REF]]; [Shah et al. 2021b]; [START_REF] Wang | Enhanced performance of Se-Alloyed CdTe solar cells: The role of Se-segregation on the grain boundaries[END_REF].

Extensive studies have shed light on a large number of questions about the real effect of the selenium alloy on CdSeTe cells and have highlighted its critical passivation effect in the bulk and at grain boundaries. However, no detailed and quantitative study of the dependence of the passivation effect on the Se content has been conducted. Furthermore, the nature and distribution of the passivated defects have not yet been demonstrated.

Conclusion of the section

The introduction of Se in CdTe-based solar cells has led to record efficiency devices. These cells have an increased J sc and maintain similar V oc as CdTe cells despite a lower bandgap, which is attributed to the higher lifetime of minority carriers compared to CdTe devices only. Several studies have contributed to a better understanding of the effect of selenium by demonstrating its passivation role on the bulk and on grain boundaries. However, no quantitative results have been reported on the diffusion of Se atoms at GB and in GI, and on their role in the internal radiative efficiency in CdSeTe layers

Cathodoluminescence characterization of CdSe x Te 1-x layers

In the following, a systematic study of several CdSe x Te 1-x samples with different Se contents is performed to gain insight into the device performance and the effect of selenium. High-resolution cathodoluminescence has been shown to be useful for characterizing polycrystalline thin films by probing the radiative efficiency within the grains, thus providing information on their electronic properties [START_REF] Bidaud | Imaging CdCl2 defect passivation and formation in polycrystalline CdTe films by cathodoluminescence[END_REF]]; [START_REF] Moseley | Cathodoluminescence analysis of grain boundaries and grain interiors in thin-film CdTe[END_REF]]; [START_REF] Moseley | Luminescence methodology to determine grain-boundary, grain-interior, and surface recombination in thinfilm solar cells[END_REF]]; [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]. In this chapter, this technique is used extensively to characterize CdSeTe photovoltaic material. In this section, we will first describe the cathodoluminescence setup and the electronmatter interaction. Then we introduce the fabrication of the five CdSe x Te 1-x samples under investigation, and the experimental conditions. Finally, we present the hyperspectral maps on which the entire analysis of this chapter is based. This tool has the ability to perform time-resolved measurements, but only the continuous * Attolight: https://attolight.com/ operation mode configuration is described here. A detailed schematic of the CL equipment is shown in Figure 4.2.

CL characterization setup

Electron source

The emission system is an electron source, consisting of a Schottky thermal field emission gun with a ZrO coated tungsten tip. Heating of the electron gun provides a continuous emission of electrons which are then gathered and accelerated by electromagnetic lenses. The electron energy or acceleration voltage is adjusted at the top of the column and can be set between 2 kV and 10 kV. The electron beam is controlled and directed by the gun lens, which determines the spread of the electron beam onto an aperture. Diameters of 100 µm, 50 µm, 30 µm, and 25 µm can be selected. Both parameters (gun lens and aperture) are set to change the beam current, resulting in an injection current that can vary over a range of several orders of magnitude, from a few tens of pA to a few tens of nA. The electron beam is then focused onto the sample with an objective lens and corrected for stigmatic aberration. The final spot size of the electron beam is about 10 nm or less. Samples are mounted on a piezoelectric control stage with nanopositioning capability and compatible with a helium cryostat. The temperature is controlled to within 0.1 K over a temperature range of 10 K to 300 K. There is a secondary electron detector in the chamber.

Light collection and detection system

The light emitted by the sample is collected through an achromatic reflective objective with a numerical aperture of 0.72. This objective, with a concept similar to a Cassegrain reflector, provides a constant, high collection efficiency over a large field of more than 150 µm, 40% of the photons emitted by a Lambertian source exit nominally the microscope. The optics collection position is fixed in the column (Figure 4.2), and is optimized only by adjusting the sample in the focal plane (maximum signal). This configuration of the collection system allows a quantitative comparison between samples, since no optical alignment is required except for the sample height (z-direction).

The collected light is focused on the entrance slit of one of the two spectrometers depending on the spectral range of interest: Horiba iHR320 or Andor Kymera 328. The light is dispersed by a reflective diffraction grating. CL spectra are recorded on an Andor Newton silicon CCD camera (1024 x 256 pixels, pixel width 26 µm) for the visible to near-infrared (NIR) range or on an InGaAs array for the NIR-IR spectral range.

The spectral resolution depends on the groove density of the grating. For example, a grating with 150 grooves.mm -1 scatters light at 0.53 nm/pixels. A punctual source at the entrance of the spectrometer can be completely defined by 4 pixels on the CCD, resulting in a spectral resolution of about 2 nm with this grating. In this detection system, the diffraction grating and the CCD camera are wavelength sensitive. A corrected luminescence spectrum is obtained by dividing the recorded spectrum by the response of the system: the external quantum efficiency of both the grating and the CCD camera.

Data acquisition and hyperspectral correction

An illustration of the data acquisition is given in Figure 4.2 [START_REF] Frouin | A path toward 25%-efficient ultrathin GaAs solar cells[END_REF][START_REF] Atwater | Plasmonics for improved photovoltaic devices[END_REF][START_REF] Frouin | CdSe x Te 1-x solar cells: the positive impact of Se analyzed by high-resolution cathodoluminescence mapping[END_REF][START_REF] Frouin | Ultrathin CdSeTe solar cells with optimized rear side contact[END_REF][START_REF] Frouin | Alternative Rear Contacts for Ultrathin CdSe x Te 1-x Solar Cells[END_REF]. The scanned area is a rectangle of size L x and L y . Then, the image of this region is rotated by 90°at the entrance of the spectrometer with size ML x and ML y where M is the magnification of the collection system. After light dispersion, a third image is formed on the CCD sensors and the spectrum is obtained by vertical summation over the 256 pixels of the CCD array. For a light source which is not centered along the x-axis, the recorded spectrum is spectrally shifted (see point B-B 2 in 3 and 4, Figure 4.2). The shift value is evaluated at approximately 0.245 nm per µm. This spectral shift is corrected by subtracting this value along the vertical direction. After this correction, the background is removed from the CL signal. It is then normalized to the exposure time of the CCD camera and spectral corrections are applied. The luminescence signal corresponds to a spectral density of photon flux per unit of wavelength. A CL map or hyperspectral cube is acquired by scanning the surface sample. A luminescence spectrum is recorded at each position of the electron-beam excitation, corresponding to a pixel in the CL map.

Spatial resolution

The spatial resolution depends on the electron-beam spot size which is about 10 nm, nevertheless the interaction volume of the electrons with the material is larger. It depends on the acceleration voltage and the average volumic mass density of the material. With Monte-Carlo simulations (CASINO software [Drouin et al. 2007]), it is possible to simulate the trajectories of the electrons. It assumes a Gaussian-shaped incident electron-beam with a beam radius of 10 nm. The distance between successive elastic collisions depends on the material density, and the atomic fraction. At 6 kV, CASINO simulation shows that 75% of the carriers in CdTe are generated in a pear-shaped excitation volume with a radius of 25 nm and a depth of approximately 75 nm (see Figure 4.3). Since the light is collected from the entire surface, some electro-generated carriers may diffuse and some radiative recombination may occur away from the injection spot. Since the diffusion length of most PV materials is about 1 µm, the CL spatial resolution is generally limited by the diffusion phenomenon, which can be lowered at low temperature. Electron-matter interactions also play a significant role in the carrier injection level and can affect the luminescence spectrum (defect saturation or poor filling). The calculation of the carrier density per second in CdTe material is detailed in the PhD manuscript of T. Bidaud [START_REF] Bidaud | Characterization of nanomaterials by cathodoluminescence for photovoltaic applications[END_REF]] and is estimated to be about 2.1 × 10 25 cm -3 .s -1 ( in standard conditions: e-beam current of 2 nA and diffusion length of 1 µm) for the upper limits . The excess carrier density injected by the e-beam in CL is significantly higher than the density of photognerated carriers under 1-sun illumination (∼10 21 cm -3 .s -1 ). This results in a concentration of injected carriers of ∆n ≃ 10 16 cm -3 with a carrier lifetime of 1 ns, and of ∆n ≃ 10 18 cm -3 with a lifetime of 100 ns, which is higher than the typical value of hole density (p-doped CdTe) of 10 14 -10 15 cm -3 .

The main features of the CL setup were presented. Its specific collection optics allow a high collection efficiency over a large area and a quantitative comparison between samples. It has a high spectral resolution (2 nm with adapted detection system) and a high spatial resolution (down to a few tens of nm), which depends mainly in the diffusion length of the material. The injection levels have been estimated for CdTe semiconductor under standard conditions. Preliminary experiments for an absolute calibration of the photon flux of the CL setup have been initiated [START_REF] Bidaud | Characterization of nanomaterials by cathodoluminescence for photovoltaic applications[END_REF]]. This would allow a direct probing of the emission rate induced by the CL excitation and lead to a quantitative determination of the radiative efficiency.

CdSe x Te 1-x samples fabrication and experimental conditions

Description of the series of CdSe x Te 1-x beveled samples Five samples are fabricated with five different Se contents ranging from 0% to 40%, to evaluate the selenium impact on the optoelectronic properties and structural changes of CdTe-based solar cells. The cells are fabricated in a superstrate configuration at the National Renewable Energy Laboratory (NREL). A soda-lime glass (commercial TEC12D) is coated with fluorine-doped tin oxide (SnO 2 :F) and intrinsic i-SnO 2 . Mg y ZnO 1-y (where y= 0.08) is then deposited by radio frequency sputtering using a mixed-oxide target with a weight ratio of 4wt% and 96wt% for MgO and ZnO, respectively. About 1 µm of CdSeTe is thermally evaporated from a ternary source of CdSe x Te 1-x whose temperature is about 650℃ while the substrate temperature remains around 450℃. Then, without breaking the vacuum, 3-4 µm of CdTe is evaporated. The samples undergo a CdCl 2 annealing treatment at 420℃ for 10 min. Then, the devices are dipped in a 0.1 mmol CuCl 2 solution in water and annealed in a tube furnace at 200℃ for 30 min in air. Finally, a 100 nm Au layer is evaporated to form the back contact. To present an extended cross-section for CL measurements, bevels at 20°were prepared by focused-ion-beam milling in an xT-Nova NanoLab (FEI Company) with a 30 kV ion beam, followed by surface cleaning with a 5 kV beam. The milling allows the exposure of the whole absorber thickness. An additional cleaning is performed by milling at a glancing angle, with Ar+ ions using a JEOL cross-section polisher tool operating at 3 kV (the detailed process is described in [START_REF] Moseley | Cathodoluminescence spectrum imag-ing analysis of CdTe thin-film bevels[END_REF]). Three bevels for each sample have been prepared. A schematic of the CdSeTe/CdTe device structure and 20°bevel, including an outline of the CL measurement area (dotted white line) is shown in Figure 4.4(b). A SEM image of the bevel with the back contact on top is shown next to the scheme. Five CdSe x Te 1-x samples have been fabricated at the NREL (US), with nominal Se content ranging from 0% to 40%. In-depth analysis is possible thanks to the preparation of bevels.

Experimental conditions

High resolution hyperspectral CL maps (256×256 pixels) have been acquired for five different samples. The acceleration voltage of the electron beam was set to 6 kV and the probe current was kept constant to 4 nA within less than 5%. At 6 kV, CASINO simulations have shown that 75% and 95% of carriers are generated in CdTe in a pear-shaped excitation volume within a radius/depth of approximately 25/75 nm and 85/150 nm, respectively. For CdSe 0.4 Te 0.6 , the size of the radius/depth is increased to 40/100 nm and 110/160 nm, respectively. The entire depth of the bevel is scanned and the CL maps size vary from 15.2×15.2 µm 2 to 32×32 µm 2 . The constant excitation and collection efficiencies during the experiments guarantee a quantitative comparison between the different samples. The focus was optimized on the top of the film, near the Au back contact, nevertheless the grains are distinguishable both towards the MZO and near the back contact. The CL intensity reported in the following corresponds to a spectral density of photon flux per unit of energy (counts.s -1 .eV -1 ). The whole acquisition time was about one hour and a half for each hyperspectral map. The same samples have been used for the fabrication of solar cells and have been characterized with a focus on lifetime and V oc measurements. These results have been published in this reference [START_REF] Ablekim | Exceeding 200 ns Lifetimes in Polycrystalline CdTe Solar Cells[END_REF]] to which we are associated, and will be discussed in Section 4.4.2.

Hyperspectral cathodoluminescence maps

The CL panchromatic maps (CL intensity integrated over the full spectrum for each pixel), measured on the five bevels under the same excitation conditions and at room temperature (RT) are plotted in Figure 4.5(a-e). Some CL intensity maps (4.5(a,c,e)) have been realigned because of a drift caused by the charge effect induced by the electron beam scan on insulator layers. The realignment procedure is based on the CL intensity contrast between the CdSeTe/CdTe emitting layer and the back side, each row of pixels is shifted to align along a horizontal axis. The results are checked with the SEM images attesting that the CL map is not distorted.

The panchromatic maps show the evolution of the grain morphology and of the luminescence intensity across the depth of each CdSeTe/CdTe thin film. The grayscale is adjusted for each sample to account for the strong differences in the luminescence intensities. In Figure 4.5(b-e), there is a clear contrast between the top and the bottom of the CL intensity maps, revealing the bilayer structure where the CdTe region can be distinguished from the CdSeTe. The CdSeTe region (bottom of CL maps) exhibits smaller grains as compared to the top CdTe area. Contrary to previous CL studies of pure CdTe thin films [START_REF] Moseley | Luminescence methodology to determine grain-boundary, grain-interior, and surface recombination in thinfilm solar cells[END_REF], we see no obvious correlation between the CL intensity and the grain size. On the other hand, we observe a strong increase of the luminescence intensity with the Se concentration, with a 10-fold enhancement between CdTe (Fig. 4.5(a)) and CdSe 0.4 Te 0.6 (Fig. 4.5(e)). A similar CL intensity contrast between regions with high and low Se content has already been observed by Fiducia et al. in a single CdSeTe/CdTe bilayer and attributed to the passivation effect of selenium [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]]. The main peak energy of luminescence spectra, which corresponds to the maximum peak, is extracted for each pixel of the hyperspectral CL maps, and plotted as peak energy maps in Figures 4.5(f-j) (the method is sketched in Figure 4.6 ). This provides a direct insight into the variation of the selenium concentration. As expected, the peak energy map of the pure CdTe sample (Figure 4.5(f)) is homogeneous over the entire depth. For the other samples, the luminescence redshifts as the selenium concentration increases. In Figure 4.5(g-j), the peak energy maps exhibit two distinct and homogeneous regions corresponding to CdTe (top) with a peak energy around 1.51 eV, and to CdSeTe (bottom) with peak energies from 1.47 eV to 1.41 eV. We note that the CdSe 0.3 Te 0.7 sample differs from the others with bigger grains and a less homogeneous CdSeTe region (Figure 4.5(i)). These panchromatic and peak energy maps acquired by cathodoluminescence provide the basis of the following analysis. From the Figure 4.5, we observe that the high-luminescence region of the panchromatic maps matches well with the CdSeTe region. This will be used in Section 4.3 to study the evolution of the bandgap as a function of both the Se concentration and the temperature. The relation between the luminescence intensity and the selenium concentration will be further investigated in Section 4.4. We also remark a steep transition between the CdTe and the CdSeTe, and small inhomogeneities along the grain structure in peak energy maps. To provide insights into these phenomena, we will analyze in the diffusion of Se atoms in Section 4.5.

Conclusion of the section

In this section we presented the fabrication of the CdSe x Te 1-x sample series under study. The cathodoluminescence setup used to characterize them was introduced, and we described the experimental conditions in which the samples will be analyzed in order to gain insights into the positive impact of Se. We have shown the panchromatic and peak energy maps acquired by cathodoluminescence at room temperature for each sample. The multiple informations contained in these maps will enable the analysis of the role of selenium in CdSeTe photovoltaic devices that will be developed in the following sections.

Selenium and temperature dependence of the bandgap 4.3.1 Bandgap and Urbach tail determination at T=296 K

To glean insights into the bandgap dependence with the Se concentration, CL spectra of the CdSe x Te 1-x regions are averaged over the areas indicated by colored rectangles in the panchromatic maps of . The normalized CL spectra are plotted in Figures 4.7(a) for a sample temperature of 296 K. As expected [START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF]], the main peak exhibits a clear redshift as the concentration in Se increases, from 1.51 eV to 1.41 eV.

Generalized Planck law and absorption model

The following analysis requires a side-note on the luminescence emission of semiconductors. The explanation is based on the modeling developed in references [Chen et al. 2021b]; [Chen et al. 2021a] for carrier density assessment by cathodoluminescence in GaAs. Under excitation electron-hole pairs are generated in the material. The light emitted by the carrier recombination is expressed by the generalized Planck law:

Φ lum = A(ℏω) × (ℏω) 2 4πℏ 3 c 2 × 1 exp( ℏω-(E f c -E f v ) k B T -1) (4.1)
where ℏω is the energy, c the speed of light in vacuum, and E f c and E f v the quasi-Fermi levels of electrons and holes, respectively. A(ℏω) is the absorptivity. Above the bandgap, the absorption coefficient α p is modeled by a parabolic band approximation:

α p (ℏω) = α c (ℏw -E g )/E 0 . (4.2)
For CdSeTe/CdTe we define α c = 30000 cm -1 at an energy E 0 = 200 meV above the bandgap (1.7 eV for CdTe). This value is chosen in accordance with the absorption coefficient values found in the literature [START_REF] Adachi | Optical properties of CdTe: Experiment and modeling[END_REF]]; [START_REF] Treharne | Optical design and fabrication of fully sputtered CdTe/CdS solar cells[END_REF]] and provided by the NREL [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]] (see Figure A.4(a)). In Appendix A.2.1, the absorption coefficients as a function of the selenium content are plotted in Figure A.4(b). The difference between the values is minor, which justifies the α c value regardless of the Se content. α p is convoluted by an exponential decay below the bandgap to account for the Urbach tail (characteristic parameter γ):

α 0 (ℏω) = 1 2γ ∞ Eg α p (E) exp - |ℏω -E| γ dE (4.3)
The absorptivity is modeled by: A = 1 -exp(-αd). The temperature T and the characteristic length d, which corresponds to the emission volume where carriers recombine radiatively, affect mainly the high-energy tail of emission spectra and have little impact on the fitted bandgap and bandtail. They are tuned in the ranges of T= 300-330K and d= 0.1-1 µm, respectively. We assume that the quasi-Fermi levels are deep in the bandgap and no band filling needs to be taken into account. A Gaussian function is added to the generalized Planck law to model the low-energy emission peak attributed to a defect level.

CL spectra fits

CL spectra plotted as solid lines in Figure 4.7(a) are fitted with the generalized Planck law and the absorption model described earlier, to accurately determine the CdSe x Te 1-x main electronic parameters. The resulting fits (dashed lines) are in very good agreement with the experimental data, as shown in Figure 4.7(a). The main fitted parameters are reported in Table 4.1. The bandgap is found 10-15 meV below the peak energy, and the defect level is found 120 meV to 180 meV below the bandgap. The comparison between the absorption coefficient calculated with the equations 4.2 and 4.3 and the one deduced from refractive index for CdTe demonstrates a decent match between them.

Energy (eV)

Absorption coefficient (cm ) x 10 4

Urbach Tail

Putting the x = 0.3 sample apart, the bandtail slightly increases with the Se content, from 10.5 meV for CdTe to 12.2 meV for CdSe 0.4 Te 0.6 (see Table 4.1), whereas the corresponding V oc measured in Ref. [START_REF] Ablekim | Exceeding 200 ns Lifetimes in Polycrystalline CdTe Solar Cells[END_REF]] increases from about 580 mV to ∼830 mV. It shows that the relation between V oc and Urbach bandtail suggested by Wolter et al. is not that straightforward [START_REF] Wolter | How band tail recombination influences the open-circuit voltage of solar cells[END_REF]].

Sample

Room temperature (296K) Exp.

Fit parameters Peak (eV) Bandgap (eV) Band tail γ (meV) Defect (eV) σ (meV)

x se = 0 1.

Determination of the CdSeTe bandgap at various temperatures

We conducted CL hyperspectral acquisitions on each sample at nine different temperatures between 15 K and 296 K. CL maps have been recorded at the same position and with the same injection conditions (6 kV, I= 4 nA) regardless of the temperature. Higher intensities are observed at low temperature (LT), resulting in a lower acquisition time. Similar to Figure 4.7(a), the normalized CL spectra at 15 K are plotted in Figure 4.8(a) as colored solid lines. The sharp excitonic peak redshifts from 1.59 eV to 1.49 eV as a function of the Se content. The same shift is observed for the broad sub-bandgap peaks attributed to A-center defects induced by complexes formed between cadmium vacancies, V Cd , and shallow donors [START_REF] Consonni | Spectroscopic analysis of defects in chlorine doped polycrystalline CdTe[END_REF]]; [START_REF] Giles-Taylor | Light Trapping in Ultrathin CIGS Solar Cells with Nanostructured Back Mirrors[END_REF]; [START_REF] Halliday | A photoluminescence study of polycrystalline thin-film CdTe/CdS solar cells[END_REF]]; [START_REF] Krustok | Recombination velocity less than 100 cm/s at polycrystalline Al2O3/CdSeTe interfaces[END_REF]]; [START_REF] Kraft | Comprehensive photoluminescence study of chlorine activated polycrystalline cadmium telluride layers[END_REF]]; [START_REF] Bidaud | Imaging CdCl2 defect passivation and formation in polycrystalline CdTe films by cathodoluminescence[END_REF] The bandgaps extracted from the averaged CL spectra (in the CdSeTe rich region) at different temperatures are plotted in Figure 4.8(b). At low temperature (T<150 K), they are approximated by the peak energy. No significant difference is found from bandgaps extracted in the CdTe regions regardless of the Se concentration. The proximity of the two sets of data obtained for x Se = 0.2 and x Se = 0.3 suggests that the actual Se content of the CdSe 0.3 Te 0.7 sample might be slightly below the nominal level. The value found for x=0.3 might be also affected by the inhomogeneities (Figure 4.5(d)).

Varshni law E g (T) for different Se alloy

For each Se concentration, these experimental data are fitted with the Varshni empirical law, which links the bandgap energy to the temperature:

E g (T ) = E g (0) - αT 2 T + β , ( 4.4)
where E g (T ) is the bandgap at a given temperature, and α and β are two constants. The three parameters α, β and E g (0) are fitted for each sample, the results are reported in Table 4.3. Interestingly, both α and β parameters are nearly independent from the selenium concentration.

E g (x Se ) and bowing parameters

The bandgap energies at 0 K, deduced from the fit of the Varshni law, are plotted in Figure 4.9 for the different Se concentrations (blue dots). They are compared to room-temperature bandgaps extracted from the fit of CL spectra (red dots), and from the optical data published in [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]] (green dots). CdSe bandgap values in the literature vary from 1.66 eV to 1.75 eV [START_REF] Lunz | Temperature dependence of the energy gap of zinc-blende CdSe and Cd 1-x Znx Se epitaxial layers[END_REF]]; [START_REF] Islam | Structural and optical properties of CdSexTe1 -x thin films grown by electron beam evaporation[END_REF]; [START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF]]; [START_REF] Gullu | Temperature effects on optical characteristics of CdSe thin films[END_REF]]; [START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]]; [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]]. In the following, we use the values given by Gullu et al [START_REF] Gullu | Temperature effects on optical characteristics of CdSe thin films[END_REF]] at both low temperature (E CdSe g = 1.75 eV, 0 K) and room temperature (E CdSe g = 1.70 eV, 300 K). The composition dependence of the bandgap is usually fitted with a second-order polynomial bowing formula written as

E g (x Se ) = E CdT e g (1 -x Se ) + E CdSe g x Se -bx Se (1 -x Se ), (4.5)
where E CdT e g and E CdSe g are the bandgaps of CdTe and CdSe, respectively, at a given temperature and b is the bowing factor. We assume, that despite Se diffusion, the Se content in the CdSeTe region is unchanged. The outlier experimental data point x=0.3 (open circles in Figure 4.9), is not taken into account in the fit. Consistent results are obtained at 296 K when comparing the bandgap extracted from CL measurements and optical indices with a bowing factor b = 0.73 and 0.71 respectively. It agrees with experimental bowing parameters from literature that range from 0.59 to 0.97 depending on the deposition technique and characterization methods [START_REF] Islam | Structural and optical properties of CdSexTe1 -x thin films grown by electron beam evaporation[END_REF]; [START_REF] Wei | First-principles calculation of band offsets, optical bowings, and defects in CdS, CdSe, CdTe, and their alloys[END_REF]]; [START_REF] Lingg | Structural and electronic properties of CdTe1-xSex films and their application in solar cells[END_REF]; [START_REF] Campbell | Comparison of cathodoluminescence and photoluminescence of CdSeTe films grown on Si by molecular beam epitaxy[END_REF] [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF], the red ones using CL spectra at RT, and dark blue values are determined using the bandgap values at 0 K extracted from the fit of experimental data with the Varshni law (equation 4.4). The bowing factor b is explicited for each fit.

Bandgap E g (T,x Se )

Since the parameters of the Varshni law and the bowing factor show little dependence on the Se concentration and temperature, respectively, we propose a set of equations to determine the bandgap E g (T, x Se ) (eV) of CdSeTe semiconductor alloys for any concentration (from x se =0 to x se =0.4) and temperature:

E g (0, x Se ) = 1.589(1 -x Se ) + 1.75x Se -0.67x Se (1 -x Se ) (4.6) E g (T, x Se ) = E g (0, x Se ) - 5 × 10 -4 T 2 T + 200 (4.7)
These equations provide a straightforward way to link the bandgap to the selenium concentration at any temperature between 0 K and 300 K.

Conclusion of the section

An in-depth analysis has been performed with high resolution CL at several temperatures on beveled CdSe x Te 1-x samples with nominal Se content ranging from 0 to 0.4. CL peak energy maps highlight two distinct regions corresponding to CdTe and CdSeTe, and the high-luminescence part of the panchromatic maps corresponds to the CdSeTe region. A 10-fold enhancement between CdTe and CdSe 0.4 Te 0.6 is observed. This higher radiative efficiency in this specific region is attributed to bulk passivation by Se. We have demonstrated that the RT CL spectra can be accurately fitted with the generalized Planck law, resulting in a bandgap estimation for each sample. CL maps acquired at multiple temperatures for each sample allow to link the bandgap (or the peak energy) to the Se concentration at any temperature.

Impact of the Se concentration on the radiative efficiency 4.4.1 CL luminscence vs. Se content relation

We have previously discussed the CL intensity evolution and the peak energy position for all samples. We reported an increase in the luminescence response and red-shift of the CL peak position with higher selenium content. To get insights into the beneficial impact of the selenium concentration on the radiative efficiency, we plotted the evolution of the CL intensity at RT as a function of the CL peak position (maximum peak) in Figure 4.10(b). Based on 256 × 256-pixels hyperspectral maps (Figure 4.5), the Figure 4.10(a) shows the scatter plot of the total number of pixels having the same (except for a small user defined interval ) peak properties: the peak energy and CL intensity for the sample CdSe 0.4 Te 0.6 . We are limiting the energy range between 1.4 eV (E peak (x Se =0.4)) and 1.5 eV (E peak (x Se =0)). Two striking features are evidenced in Figure 4.10(b). First, the data from different samples are superimposed. It means that the same luminescence intensity is found for the same Se concentration in the four samples, despite the different nominal concentrations and diffusion processes. Second, the relation between the luminescence intensity and the Se concentration can be determined quantitatively. A 10-fold increase is found between CdTe (1.50 eV) and CdSe 0.4 Te 0.6 (1.4 eV). The relation between the CL intensity and the peak energy can be fitted with a phenomenological exponential law:

I CL (E) = a • exp(-E/b) (4.8)
where b = 53 meV.

The more Se in the layer, the more significant defect passivation. Figure 4.10 attests of a significant and direct correlation between Se content and high CL luminescence. Regardless of the sample, grain size, or the position on the bevel, CL intensity is characterized by the amount of Se in the device.

Insights on defect passivation with Se

The mechanism behind the reduction of non radiative recombination in the bulk of CdSeTe is still largely unexplored. In this part, we present a qualitative investigation in an attempt to better understand how the selenium passivation works by trying to connect the evolution of the radiative efficiency, the V oc and lifetime performances of cells and the presence of defects.

Investigation of defects

From LT CL acquisitions, we extracted the CL spectra by dissociating the CdSeTe area from the CdTe one for each sample. The CL peak energy values are extracted from spectra averaged over two distinct areas : a 20×256 pixels CdTe area and a 20×256 pixels CdSeTe area considered as homogeneous. The top line in Figure 4.11 shows a sharp excitonic peak attributed to the CdTe band-edge transition with a low-energy band defects (x=0). As the Se concentration increases, we notice a second defect broad band emission at a lower energy (∼1.25 eV) and only for the lowest temperature for samples with Se content x=0.1, 0.2 and 0.3, which disappears for x=0.4.

On the other hand, the bottom line displays the excitonic peak attributed to the CdSeTe band edge transition (redshifted) with a single low-energy band defect. We compare the energy difference between the broadband defect and the sharp excitonic peak ∆E for both regions. Note that we consider the higher energy of the defect emission when there are two broad peaks. From Figure 4.12, it can be seen that ∆E is larger for the CdSeTe part (bottom graph) than for the CdTe part (top graph) (0.2-0.28 eV compared to 0.17 eV, respectively). Overall, ∆E CdT e remains constant regardless of the Se content and the temperature. However, as the Se content increases (the excitonic peak energy decreases), ∆E CdSexTe 1-x also increases, from 0.20 eV to 0.28 eV at 15 K for x=0.1 and x=0.4, respectively. This implies that the sub-bandgap peak energy decreases even faster with increasing concentration. According to Fiducia et al., this indicates that the defect emission in the CdSeTe material originates from deeper donor and acceptor states than in the CdTe material. They suggest that either these deeper defects do not exist in CdTe or they are present but do not undergo radiative recombination as in CdSeTe [START_REF] Fiducia | Selenium passivates grain boundaries in alloyed CdTe solar cells[END_REF]]. Looking at the top line of Figure 4.11 for samples with Se concentration from x=0.1 to x=0.3, the second hypothesis seems to hold as we distinguish two types of defect emission, while the CL spectra behavior of the CdSe 0.4 Te 0.6 sample tends to confirm the first suggestion.

(a) (c) (b) (d) (e) (f) (h) (g) (i)

CdSeTe area

CdTe area To analyze the spatial distribution of each band of the CL spectra at 20 K over the hyperspectral maps, we attributed a color to each band: blue for the defects, red for the CdSeTe exciton and green for the CdTe exciton. The CL spectra integrated over the whole surface for each sample are plotted in Figure 4.13(a-e) with the corresponding spectral ranges highlighted. The integrated intensities are calculated over a defined spectral region for each pixel and then normalized. We isolate the two excitonic peaks and integrate over the entire defect spectral range without distinguishing which is from the CdSeTe or the CdTe region. The CL maps of the defect-related transition are plotted in Figure 4.13(f-j). The CL maps of the excitonic emissions have been superimposed and plotted in Figure 4.13(k-o) and the scale has been adjusted for each map and for each excitonic emission. This Figure 4.13 has already been presented in the PhD manuscript of T. Bidaud [START_REF] Bidaud | Characterization of nanomaterials by cathodoluminescence for photovoltaic applications[END_REF]].

We observe that the defect band intensity within the film increases from x= 0.1 to x= 0.3 and then decreases for x=0.4. Its repartition is quite homogeneous but slightly more intense in the CdSeTe area except for the sample x=0.3. Nevertheless, the highest defect band intensity is observed at the interface between CdTe and CdSeTe for each sample. While it is confirmed that the CL intensity increases with Se content, we find no clear correlation between the high radiative efficiency and the local variation of the defect distribution.

Radiative efficiency as a function of Se

Lifetime and V oc measurements were performed on similar CdSeTe/CdTe samples and presented in Ref [START_REF] Ablekim | Exceeding 200 ns Lifetimes in Polycrystalline CdTe Solar Cells[END_REF]]. Figure 4.14(a) reports that the V oc is correlated with the measured carrier lifetime. It confirms that the defect passivation observed at high Se content is consistent with the high V oc . However, solar cells with 40% Se composition have a slightly lower V oc , although they have a higher CL intensity. This observation leads us to believe that there are more complex mechanisms behind the passivation of defects by selenium. The activation energy of defects could be obtained from the CL intensity as a function of the temperature. For each sample, we integrate the CL intensity of the 200 brightest pixels in the defect emission spectral range for CdTe and CdSeTe regions, respectively, and plot the average as a function of the temperature (see Figure 4.14(b)). Unfortunately, we do not observe a distinct constant intensity for both regions at low temperatures, which would indicate that only radiative recombinations is occuring. Therefore, no accurate activation energy values can be extracted from fitting these data. We note that regardless of the regions, the sample x=0.3 exhibits the highest intensity, which would indicate a higher passivation.

(b)

ar cell with 30% Se composition. The green line shows a single exponential dTe solar cells ith Se composition ranging from 0% to 40% The straight

CdTe

CdSe Te We have extensively discussed the impact of Se on the luminescence response of the samples, and in the following we will focus on the Se gradient behavior. Indeed, Figure 4.13 shows a sharp transition between the CdTe and CdSeTe regions, and displays that some grains are both green and red which can be explained by carriers diffusion or by a slight Se diffusion within a grain. These observations and the unusual CdSe 0.3 Te 0.7 sample lead us to look at the exact distribution of Se in the CdSeTe/CdTe samples and how it diffuses into the layer.

Conclusion of the section

In this section, we have demonstrated a strong correlation between the CL intensity and the selenium composition. It is shown that the radiative efficiency improves with Se content and this is independent of the structural properties of the sample. We have evidenced a 10-fold increase of the radiative efficiency of CdSeTe from x= 0 to x= 0.4. Investigations on the defects did not allow us to identify a correlation between the radiative efficiency and the defects distribution. An increase in V oc and in the minority carriers lifetimes is observed with increasing Se content from x= 0 to x= 0.3, but a slight decrease is reported for x= 0.4, despite higher CL intensities. These observations point out the complex role of Se passivation.

Gradient of Se 4.5.1 Gradient of Se concentration in the depth

The depth profiles of the peak energy integrated along the bevels widths are shown in Fig- Se content (%)

x =0.1

x =0.3 x =0. 2Se content (%) The peak energy rises from the MZO/CdSeTe interface until reaching a plateau whose associated energy matches CdTe. Putting once again the CdSe 0.3 Te 0.7 sample appart due to its higher inhomogeneity, a steep gradient at the interface between the CdSeTe and CdTe layers is evidenced for all samples. The thickness of the interdiffusion region is of the order of a few hundreds of nanometers. Using the Fick's second law, the profiles can be fitted with an error function of the form erf (z/δ) [START_REF] Colegrove | Se diffusion in CdTe thin films for photovoltaics[END_REF]] and lead to the characteristic diffusion lengths δ given in Table 4.4. However, this approximate assessment assumes a linear relation between the Se concentration and the peak energy, and the microscopic diffusion effects are smoothed by the integration along the bevel width.

Detailed maps of the Se content can be calculated using the relations between the bandgap energy and the Se concentration, equations 4.6 and 4.7. At room temperature, with a 10 meV shift between the peak energy (E peak ) and the bandgap, we obtain for x Se < 0.4 (see Figure 4.9): . For a nominal Se content of 0.1, we find an average of x Se = 9% in the CdSeTe region of CL maps, confirmed by the 12% found by SEM/EDS. For a nominal Se content of 0.2, a selenium concentration lower than expected is found, with an average of 14-15% in the CL map, corroborated by an estimation of 17% by SEM/EDS and 10-14% by STEM/EDS. The CdSeTe region of the CdSe 0.3 Te 0.7 sample is more heterogeneous, with an average x Se around 20% for both CL and SEM/EDS estimation, and a higher concentration exhibited locally in the CL map (Figure 4.15(d)). The estimated Se content are globally always below the nominal concentration. Overall, the absolute value of x Se is difficult to estimate precisely, but the high spectral resolution of CL maps translates into an estimation of x Se with an accuracy of less than 0.5% for low Se concentration, and can be used to study the diffusion of selenium at the nanoscale.

E peak (x Se ) = 1.51(1 -x Se ) + 1.67x Se -0.67x Se (1 -x Se ), ( 4
Maps in Figure 4. 15(b-d) exhibit concentration inhomogeneities at the grain level in the interdiffusion region. In the CdTe region, a small concentration of Se (about 1%) is found at the grain boundaries. It could be the signature of higher diffusion coefficients of Se at grain boundaries as compared to grain interiors, with a potential impact on the transport and passivation properties.

Se diffusion along grain boundaries

Selenium diffusion is mainly driven by the CdCl 2 annealing and appears to be favoured along GBs [Munshi et al. 2018b]; [START_REF] Fiducia | Selenium passivates grain boundaries in alloyed CdTe solar cells[END_REF]]; [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]. From Se content maps, we can quantify the Se composition gradient at different positions, followed by an estimation of the diffusion length in the bulk and along grain boundaries. We illustrate this analysis for the CdSe 0.2 Te 0.8 sample in Figure 4.16. The map of x Se exhibits a clear increase of the Se content in the GBs of CdTe close to the interdiffusion region (Fig. 4.16(b)). Figure 4.16(c) shows an increase of 1% across a GB. In-depth profiles of x Se have been extracted from different linescans (colored lines in Fig. 4.16(a)), in order to investigate the decrease of the selenium concentration away from the Se-rich region, in grain interiors (Fig. 4.16(d)), along GBs (Fig. 4.16(e)), and along a straight line (Fig. 4.16(f)). The linescans are integrated over a width of 150 nm, and the pixel position is translated into the depth z.

The profiles are fitted with a 1D diffusion model [START_REF] Colegrove | Se diffusion in CdTe thin films for photovoltaics[END_REF]] which assumes a constant reservoir of Se elements at the CdSeTe/CdTe interface, and we determine the diffusion coefficient D of Se atoms during the annealing process performed at 420°C during t = 10 min:

C(t, z) = C 0 2 erf c( z -h 2 √ Dt ), (4.10)
where erf c is the complementary error function, C 0 the Se concentration and h is the abscissa of the transition zone between CdSeTe and CdTe. The results are shown in Figure 4.17 For the intragrain diffusion (orange dashed line in Figure 4.16(d)), care was taken to extract a profile perpendicular to the plan surface, in a large grain. A bulk diffusion coefficient of D= 1.2 × 10 -13 cm 2 /s is found. The same fitting procedure leads to a diffusion coefficient six times larger along grain boundaries, D= 7.8 × 10 -13 cm 2 /s. This result confirms that the diffusion of Se is mainly driven by grain boundaries, and provides a quantification of the diffusion coefficients. The use of a vertical linescan (dashed green line) across the depth of the CdTe layer is a simpler way to extract profiles of x Se , and results in D= 1.1 × 10 -12 cm 2 /s, close to the value obtained at GBs.

The same analysis was performed on straight linescans for each sample with nominal compositional value x= 0.1, 0.2, 0.3, in three or four distinct areas. The results of the diffusion coefficients are reported in Fig. 4.17. D increases with the concentration of Se in the CdSeTe reservoir, from D= 0.5 × 10 -12 cm 2 /s for x=0.1 to D= 5 × 10 -12 cm 2 /s for x=0.3. The outlier green point of the CdSe 0.3 Te 0.7 sample, refers to a small area of the bevel where the transition CdSeTe/CdTe is very steep in contrary to the rest of the sample. The vertical linescans along which the diffusion coefficients are calculated are shown in Appendix A.10 where the corresponding values are also reported. The diffusion length δ expressed in equation 4.11 and provides a distance of how far the concentration has propagated by diffusion in time t.

δ = 2 √ Dt (4.11)
where D is the temperature-dependent diffusion coefficient. For the sample with nominal Se content x= 0.1, the diffusion length of 400 nm is of the order of magnitude of the grain size. As for the sample with x= 0.3, its diffusion length of about 1 µm is less than the total thickness of the material.

We have shown that selenium passivates both grain interiors and GBs of alloyed CdSe x Te 1-x material. However, the mechanisms behind the selenium passivation in the bulk and in GB remain to be determined. For example, the conversion of deep defects into shallower ones compared to band-egdes is a possible explanation [Shah et al. 2021b]. Another viable alternative proposes that Se interacts directly with harmful defects (such as Te antisite defects of cadmium vacancy). For passivation in GBs, we did not see any significant segregation of Se to the GBs with STEM-EDS (see discussion in Appendix A.2.2 and Figure A.9) for the CdSe 0.2 Te 0.8 sample [START_REF] Fiducia | Selenium passivates grain boundaries in alloyed CdTe solar cells[END_REF]].

Conclusion of the section

Maps of selenium concentration were plotted using equation 4.9. Estimation of the Se content was proposed and confirmed by SEM-EDS and STEM-EDS. While the determination of the absolute value is complex, thanks to the high-spectral resolution of the CL maps we reported an estimation of the Se content with an accuracy of 0.5% for low concentration. We presented a quantitative study of the Se gradient in the CdTe thin film, emphasizing on the concentration inhomogeneities at the grain boundaries within the interdiffusion region. The diffusion coefficients of Se atoms in both GI and GBs were quantified thanks to a 1D analytical model (eq. 4.10). We have evidenced that the diffusion of Se during the critical CdCl 2 annealing is mainly driven by GBs. Furthermore, it attests that Se reduces detrimental non-radiative recombination at GB, which explains the superior performance of CdSe x Te 1-x solar cells.

Investigation of high quality CdSe x Te 1-x devices

Thanks to a collaboration between Arizona State University (ASU), Colorado State University (CSU) and C2N, we had access to high-performance CdSeTe-based devices. Four different cell architectures are investigated in this study and described in Figure 4.18(a). We compare CdSeTe/CdTe and CdSeTe solar cells with two types of back contacts. A baseline back contact structure made of Te and a passivating back contact stack consisting of alumina, a p-type boron-doped hydrogenated amorphous silicon capped off with an indium tin oxide/silver. The corresponding sketches are plotted in Figure 4.18(a). A CdSeTe/CdTe/Te control sample (19934) is introduced in addition to the previously described devices.

In actual devices, the non-radiative recombination losses are quantified by the external radiative efficiency (ERE). This means that in ideal case, the ERE is limited by the external luminescence escaping in the free space. Note that even in the ideal case the maximum ERE value is not always 100% depending on the optics of the sample [START_REF] Ganapati | The voltage boost enabled by luminescence extraction in solar cells[END_REF]]. As a baseline, Green et al. reported the ERE of record photovoltaic devices: 35.7% for GaAs, 2.4% for CIGS and 0.008% for CdTe [START_REF] Green | Pushing to the Limit: Radiative Efficiencies of Recent Mainstream and Emerging Solar Cells[END_REF]]. Figure 4.18(b) presents the corresponding state-of-the-art external radiative efficiencies measured at ASU. The ERE measurements were obtained with calibrated PL measurements. The samples were excited with a low-power 532 nm diode-pumped solid-state laser chopped at a frequency of 105-110 Hz, in parallel with a 1sun-equivalent white light-emitting-diode bias light. The emission was collected with a 715 nm absorptive long-pass filter. The ERE reaches very high values: almost 3% for the best ones [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]]. However, quantitative photoluminescence measurements have revealed that most of the emission comes from a sub-bandgap peak. In this section, we present the preliminary results of CL characterization, performed on these five devices to gain insights on this defect emission peak, in the hope of attributing it to a specific region of the samples (like GB or the passivated back interface). To conduct an in-depth analysis, we performed on each sample a bevel using a FIB (see the scheme in Figure 4.4(b)), and characterized them by cathodoluminescence. The usual 750 nm 1000 nm spectral range is not suited to investigate on the defect emission which occurs in the NIR range. An adapted collection system is used, with a -80°C cooled InGaAs camera, a Kymera spectrometer and a grating 150/1250 blazed at 1250 nm. Since the sub-bandgap emission spectra are larger than the spectral dispersion of the grating, we need to acquire two maps at the same position on the sample, but with the grating centered at two distinct wavelengths (near 900 nm and 1200 nm). We performed the experiments at 20 K, to increase the signal to noise ratio.

The two spectra collected on the same position are joined together to form only one spectrum at each pixel of the CL map.

CL spectra averaged over the entire surface of the five bevels are plotted in Figure 4.19 on the same graph to facilitate the comparison between samples. The control sample (dashed green line) is less luminescent compared to the others, emphasizing their high-quality material. The excitonic peak linked to CdSeTe is the one with the highest CL signal around 850 nm for each sample, the other peaks are attributed to defect emission. Depending on the structure of the absorber the defect peaks energy is different. For CdSeTe cells only the defect peak energy is centered around 1180 nm, while for CdSeTe/CdTe devices its energy is around 1050 nm and other contributions are visible between 850 nm and 930 nm. Note that the peak at 910 nm is common for each sample. Interestingly, the 1809-10L sample that displays the highest ERE is less luminescent (apart from the control sample). Another counter-intuitive observation is that for samples with the same absorber structure, the passivated one (solid orange or dark blue curves) displays a lower CL intensity. Similar data treatments as in Section 4.4 and Figure 4.13 were performed for the defect study. CL maps of the defect-related transition were plotted, and as observed with the previous batch of samples, this emission does not appear to be associated with any particular area within the bevel. A more detailed analysis highlights the complexity of the collected signal, which is highly position dependent. Figure 4.20(a) shows a LT CL map of the passivated CdSeTe/CdTe sample restricted to the interdiffusion region of the sample 1809-5L.

The CL spectra of four different pixels indicated with the colored arrows are plotted in Figure 4.20(b). Compared to the averaged CL spectrum, the excitonic peak at 820 nm is easily distinguishable from the defect emission. The heterogeneity of the defect signal as a function of its position is pointed out with contributions that might come from the CdTe, CdSeTe or the interdiffusion region. Moreover, the CL intensity of the defect peaks are almost comparable to the excitonic peak.

Conclusion of the section

We have presented the preliminary characterization results of high-quality CdSeTe-based devices with state-of-the-art ERE. To glean insights on a sub-bandgap peak, we have performed low temperature CL measurements and collected the signal with an InGaAs detector. Two CL maps were acquired on the same area but centered at different wavelength to get the full emission of the defect in the infrared for each sample. Although we can't attribute the whole defect emission to a specific area of the sample, we have emphasized the complexity of the signal, strongly dependent on the structure of the solar cells but also on the position on the sample. Further study is required to better understand the mechanisms behind the high performances of these samples.

Summary

In this chapter, we report on the characterization of polycrystalline CdSe x Te 1-x materials by high-resolution cathodoluminescence. The addition of selenium to a CdTe thin film significantly improves the performance of solar cells. This is partly due to the bandgap bowing, but this alone is not sufficient to explain all the reported improvements. A systematic study with different Se concentrations (CdSe x Te 1-x , x from 0 to 0.4) allows us to quantify the passivation role of Se.

High resolution CL maps were performed on the five samples for nine different temperatures at the same area of interest. The bandgap is extracted from spectra averaged on the CdSeTe region, and we provide a set of equations to link the bandgap and the Se content at any temperature.

We demonstrate a ten-fold luminescence enhancement between x= 0 (CdTe) and x= 0.4 and correlate high CL luminescence with Se content, confirming of its passivation role. We evidence that the radiative efficiency improves with the Se content regardless of the structural properties of the sample. We study the impact of defect passivation and radiative efficiency on macroscopic measurements such as lifetime and V oc .

Finally, high-resolution CL is used to map the Se concentration probed by the luminescence peak position. We conducted a detailed analysis to quantify the diffusion coefficient of Se atoms both in the grain interior and at the grain boundaries during the thermal annealing for the sample with nominal Se content x= 0.2. We evidenced that Se atoms channel through the GBs to diffuse into the CdTe film, with a diffusion coefficient more than six times higher at the GB than in the grain bulk, which also supports the passivation of the GBs. From this analysis we deduced a robust and simpler way to extract the diffusion coefficient of Se atoms from all other samples.

Preliminary CL experiments in the IR and at low temperature were performed to evaluate the effect of different structure, back contacts or passivation layers on CdSeTe-based devices. Although remarkable performances and competitive prices place CdTe cells at the forefront of thin film panel commercialization, this technology still suffers from high voltage deficit. Numerous studies are conducted on this hot topic and present several back contact types, or introduce new layers to ease the separation of hole-electron pairs and the transport of minority carriers. However, to date the V oc of the record efficiency CdTe cell is still below 900 mV, lower than in GaAs cells, even though GaAs has a smaller bandgap. This V oc deficit value, about 500 mV, deeply contrasts with the ones from mature thin film devices.

In this chapter, we first introduce the voltage deficit issue for CdTe and CdSeTe cells, then we discuss on solutions proposed by the scientific community to overcome it. In Section 5.2, we present our experimental investigation of different back contacts on CdSeTe/CdTe thin films. Since our final goal is to reduce the thickness of CdSeTe/CdTe cells, highly reflective contacts are studied. The effect of a p-type transparent conductive oxide contact (Cu 2 O) is extensively described and the impact of a passivating layer on the cells performance is also analyzed. Lastly, in Section 5.3, the performances of 800 nm-thick CdSeTe/CdTe cells are presented and compared to standard devices (3-4 µm). Based on optical measurements we simulate, using the transfer matrix method, the absorption spectrum of both types of cells (standard and thin cells), trying to remain as consistent as possible with the measurements.

Factors limiting the efficiency of CdTe solar cells

Voltage deficit in polycrystalline CdTe and CdSeTe solar cells

Despite multiple materials adjustment, like CdS window layer switched with MZO or group V dopants used instead of Cu, the V oc of CdTe solar cells remains low and still limits the efficiency (32% is achievable according to the Shockley-Queisser limit). The V oc of record-efficiency devices is limited to 900 mV and still suffers from a large V oc deficit E g /q-V oc , where E g is the absorber's bandgap. Compared to other materials in Figure 5.1, we observe that the CdTe is well behind its theoretical limit and exceeds 500 mV. In comparison, GaAs displays one of the lowest V oc deficit equal to 280 mV. Keeping the same FF and J sc of the CdTe record device but considering reaching the same V oc deficit as the GaAs would lead its efficiency to be improved from 21% to 29%. Note that with Se alloying, both the bandgap and the voltage deficit are lowered compared to CdTe device [START_REF] Bothwell | Characterization of thin CdTe solar cells with a CdSeTe front layer[END_REF]. It is justified by its important passivation role in the bulk as well as in grain boundaries as reported in Chapter 4. The dotted line next to the CdSeTe marker in Figure 5.1 stands for the band gap incertitude of the record device (evaluated between 1.4 eV and 1.45 eV)) (V) [Green et al. 2022a]. The dashed red line represents the limiting value of the bandgap voltage offset V oc deficit of a solar cell using the detailed balance model. It indicates the highest voltage a solar cell of a given band gap can achieve [START_REF] King | Band gap-voltage offset and energy production in next-generation multijunction solar cells[END_REF]; [START_REF] Brandt | Determining interface properties limiting open-circuit voltage in heterojunction solar cells[END_REF] The prediction of 29% efficiency is evaluated in a simplistic way, and although attractive, appears unrealistic. Therefore, more extensive and realistic simulations have demonstrated that CdTe devices with efficiencies above 24% could be achieved by improving the interfaces and reducing detrimental back contact recombination [START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grain-boundary recombination for 25% efficient CdTe solar cells[END_REF]; [START_REF] Ablekim | Interface Engineering for 25% CdTe Solar Cells[END_REF]]. Previous models did not include surface and interface recombination because of the expected dominance of bulk recombination [Sites et al. 2007]. However, as carrier lifetime is enhanced and bulk material properties continue to improve [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1V barrier[END_REF], defective back and front interfaces become in turn efficiency-limiting factors. Front surface recombination velocities for polycrystalline CdTe is evaluated around 10 5 cm/s [START_REF] Reese | Intrinsic surface passivation of CdTe[END_REF].

Origin of the voltage loss

Recently, Onno et al. have presented a comprehensive voltage loss analysis of state-of-theart CdSeTe cells [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]]. The thermodynamic voltage limit V oc, ideal , the internal voltage iV oc and the external voltage V oc are differentiated to evaluate what is the origin of the loss, which depends on the solar cell fabrication choices such as doping, passivation or back contacts. We report a small part of their analysis for two different CdSeTe-based solar cells. In Figure 5.2, we observe that the V oc, ideal is lower for the As-doped sample than for the Cudoped. The V oc, ideal is calculated in the detailed balance limit [START_REF] Shockley | Detailed Balance Limit of Efficiency of p-n Junction Solar Cells[END_REF], and this discrepancy is justified by sub-bandgap absorption in the As-doped sample that increases the J sc but decreases V oc, ideal according to equation 5.1. This V oc, ideal reduction is amplified in the presence of a highly reflective back contacts.

V oc,ideal = k B T q ln a(λ)Φ exc (λ)dλ a(λ)Φ BB (λ, T )dλ + 1 (5.1)
where k B is the Boltzmann constant, T is the sample temperature, q is the elementary charge, a(λ) is the sample's absorptance, Φ exc and Φ BB are the excitation and the blackbody photon current, respectively. V oc, ideal is then compared to iV oc and V oc . Based on Ross' study in 1967, the internal open-circuit voltage is correlated to the external radiative efficiency [START_REF] Ross | Some thermodynamics of photochemical systems[END_REF]]; [START_REF] Rau | Reciprocity relation between photovoltaic quantum efficiency and electroluminescent emission of solar cells[END_REF]] as follows:

iV oc = V oc,ideal - k B T q |ln(ERE)| (5.2)
Hence, the non-radiative recombinations reduce the quasi-Fermi level splitting (QFLS = q×iV oc ) from the theoretical limit. From Figure 5.2, we observe that sample 2 exhibits a very high iV oc , linked to its excellent ERE measured by calibrated photoluminescence measurements. Its passivating back contact appears to be beneficial and results in a decrease of non radiative losses. Note that this particular sample is similar to the ones presented in the previous chapter in section . Finally, the study emphasizes on the selectivity losses linked to resistances encountered by charge carriers. This last loss is responsible of the difference between the iV oc and the V oc :

V oc = S oc iV oc (5.3) where S oc is the selectivity metric of the device. While sample 1 has a decent selectivity, this is the primary identified reason of the poor V oc of sample 2. The passivated sample exhibits very high ERE but poor selectivity (orange markers), the Cu-doped sample shows high selectivity but high non-radiative recombination (blue markers). Onno et al. highlight in their study that for open-circuit voltage to exceed 1 V, the back contacts should be both selective and passivating, coupled with high quality absorber [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]]. [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF].

Glass

Impact of interfaces and back contacts

Many parameters are responsible for the V oc deficit, its dependence on material properties has been outlined in Chapters 1 and 4, and here we highlight the role of three other factors.

Minority carrier lifetime improvement: Carrier lifetime of several hundreds of ns is necessary to boost the CdTe cells efficiency, and recently, numerous studies have shown a significant gain in lifetime reinforced by the Se alloy [Munshi et al. 2017]; [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF]]; [START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]. Very long lifetimes exceeding 200 ns for CdSeTe devices have been reported [START_REF] Ablekim | Exceeding 200 ns Lifetimes in Polycrystalline CdTe Solar Cells[END_REF]]. Nevertheless, Burst et al demonstrate that even with a high minority lifetime of hundreds ns, the V oc could still be limited to less than 900 mV because of high surface recombination velocities, which is proportional to the density of traps per unit area [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1V barrier[END_REF]]. To minimize the detrimental impact of high surface recombinations and interfaces, surface passivation is essential.

Front interface passivation: The emitter/absorber interface presents high carrier recombination, and several models show the necessity to carefully design the band alignment [START_REF] Bittau | Analysis and optimisation of the glass / TCO / MZO stack for thin film CdTe solar cells[END_REF]. MZO is progressively replacing the CdS window layer due to its high bandgap and because elemental ratio of the ternary alloy can be varied offering the possibility to tune the bandgap. The emitter/absorber conduction-band offset can be higher with MZO than with CdS, and while generating a small positive barrier at the front interface a larger built-in potential is formed. The hole density is then reduced, suppressing detrimental interface recombinations [START_REF] Song | Emitter/absorber interface of CdTe solar cells[END_REF]. Various experiments have shown the positive impact of a well engineered front interface [START_REF] Kephart | Band alignment of front contact layers for high-efficiency CdTe solar cells[END_REF]; [START_REF] Ablekim | Tailoring MgZnO/CdSeTe Interfaces for Photovoltaics[END_REF]; [START_REF] Bittau | Analysis and optimisation of the glass / TCO / MZO stack for thin film CdTe solar cells[END_REF]. For example, with a CdTe doping of 2 × 10 16 cm -3 , a change of two orders of magnitude of the interface recombination velocity (10 5 and 10 3 cm/s) can generate up to a 5% difference in the efficiency of the device [START_REF] Ablekim | Interface Engineering for 25% CdTe Solar Cells[END_REF]]. Nevertheless, to reach an optimal band alignment the main factor limiting the performances remains the low doping of the absorber [START_REF] Ablekim | Interface Engineering for 25% CdTe Solar Cells[END_REF] compared to lifetime or surface recombination. These models show that multiple parameters such as the conduction band offset, the emitter and absorber doping, the distribution of defects or the carrier lifetime, are strongly interconnected and each of them must be considered for the fabrication of new and even more efficient devices.

Back interface passivation and contact: CdTe has a high electron affinity of χ = 4.5 eV, which is not suitable to form easily an ohmic contact. To ease the holes extraction, many types of back contact were attempted over the years [START_REF] Kumar | Physics and chemistry of CdTe/CdS thin film heterojunction photovoltaic devices: Fundamental and critical aspects[END_REF]; [START_REF] Hall | Back contacts materials used in thin film CdTe solar cells-A review[END_REF]. While no ideal solutions were found, multiple alternatives exist and we present in the following the outcomes that contribute to the development of efficient back contacts.

To avoid large hole energy barrier and less resistance, the best contacts are metals with a high work function. However, metal contacts directly applied to the absorber tend to form Schottky barriers (Φ SB ) (even Au with a work function of 5.3 eV). Figure 5.3 shows the band alignment for a Schottky contact formed between CdTe and a metal contact. The contact barrier Φ SB at the CdTe/metal interface is given by the difference between the CdTe valence band edge and the Fermi energy in the metal (Φ SB = E g /q + χ -Φ m ). Such barrier for holes results in a limitation of hole collection and in V oc and FF losses, visible with a roll over effect in JV curve, that will be further discussed in Section 5.2.

The electrons tend to diffuse toward the back because of the band shape, which increases recombinations. To ease the formation of a quasi-ohmic back contact, Cu is an acceptor widely used. The hole density is increased due to Cu diffusion in CdTe, the barrier is reduced and holes are extracted more easily. Cu can be directly deposited or diffused through a chemical CuCl treatment, the latter giving more stable devices [START_REF] Ablekim | Thin-Film Solar Cells with 19% Efficiency by Thermal Evaporation of CdSe and CdTe[END_REF]; [START_REF] Artegiani | Analysis of a novel CuCl2 back contact process for improved stability in CdTe solar cells[END_REF]; [Bastola et al. 2019b]. However, copper diffusion in CdTe is a source of damages in the long term [START_REF] Corwine | Copper inclusion and migration from the back contact in CdTe solar cells[END_REF]. Another strategy used to address the hole collection issue is to add a back buffer layer: an interstitial layer between the absorber and the rear contact. The

Φ SB E c E v E F E VAC χ Φ m
CdTe Metal [START_REF] Kumar | Physics and chemistry of CdTe/CdS thin film heterojunction photovoltaic devices: Fundamental and critical aspects[END_REF].

goal is to achieve a band bending limiting interface recombinations [START_REF] Liyanage | The Role of Back Buffer Layers and Absorber Properties for >25% Efficient CdTe Solar Cells[END_REF]. For example, metal telluride such as ZnTe:Cu [START_REF] Li | Controlled activation of ZnTe:Cu contacted CdTe solar cells using rapid thermal processing[END_REF]; [START_REF] Wolden | The roles of ZnTe buffer layers on CdTe solar cell performance[END_REF]] presents high performances. A major role is also given to the chemical treatment of the CdTe rear surface by etching [Bastola et al. 2019a], leading to a beneficial Te-enrichment [START_REF] Song | Te layer to reduce the CdTe back-contact barrier[END_REF]]. The band diagram in Figure 5.4(a) sketches the possible impact of copper at the rear for CdTe solar cells where the back contact barrier is both lowered and narrowed. With Cu, the conduction band bending repels electrons at the interface which decreases the non-radiative recombinations. Similar scheme is shown in Figure 5. 4(b) with ZnTe as a buffer layer, Cu appears to favor a bending, lowering interface recombination as described by Liyanage et al [START_REF] Liyanage | The Role of Back Buffer Layers and Absorber Properties for >25% Efficient CdTe Solar Cells[END_REF]].

While numerical model offers a path to 28% efficiency CdTe devices [START_REF] Duenow | Back-surface recombination, electron reflectors, and paths to 28% efficiency for thin-film photovoltaics: A CdTe case study[END_REF], the formation of a chemically stable and a barrier-free back contact without recombination remains an ongoing challenge. In the next part, examples of surface passivation carried out on CdTe or CdSeTe cells are detailed. [START_REF] Corwine | Copper inclusion and migration from the back contact in CdTe solar cells[END_REF]]; [START_REF] Li | Controlled activation of ZnTe:Cu contacted CdTe solar cells using rapid thermal processing[END_REF]

Notable results on passivation layers for CdTe solar cells

Among back buffer layers, surface passivation with an alumina layer in a double heterostructure Al 2 O 3 /CdSeTe/Al 2 O 3 resulted in very high carrier lifetime, 430 ns in 2018 [START_REF] Kephart | Sputter-Deposited Oxides for Interface Passivation of CdTe Photovoltaics[END_REF], and 750 ns [START_REF] Kuciauskas | Radiative Efficiency and Charge-Carrier Lifetimes and Diffusion Length in Polycrystalline CdSeTe Heterostructures[END_REF]] in 2020. Moreover, the interface recombination velocity at alumina/absorber interface has been evaluated below 100 cm/s [Kuciauskas et al. 2018].

Noteworthy results on surface passivation by Al 2 O 3 for CdTe or CdSeTe cells are reported in Table 5.1. High carrier lifetime and high PL intensity have been reported when Al 2 O 3 is added, confirming its passivation role. However, little improvement has been observed concerning its direct effect on the V oc . Further studies are required, but it is acknowledged that the passivation layer must be thin enough to allow the carriers to tunnel through. Moreover, its positive impact seems to be related to chemical passivation which reduces defects (at the interface or in the bulk) rather than to a reduction in the band bending at the back side. Indeed, better results were observed in the majority of the studies after the CdCl 2 treatment. Two other significant results combining together copper and alumina layers are worth mentioning. A p-type and metallic Cu-Al 2 O 3 layer was created as a back contact for CdTe solar cells. Cu diffusion was controlled through the Al 2 O 3 layer deposited by Atomic Layer Deposition (ALD), allowing an efficiency gain of almost 2% compared to standard devices [START_REF] Lin | A novel p-type and metallic dual-functional Cu-Al2O3 ultra-thin layer as the back electrode enabling high performance of thin film solar cells[END_REF]. Finally, a Cu x AlO y solution was spin-coated at the CdTe rear side to passivate the surface. Increased PL intensity, carriers lifetime longer than 650 ns and increased V oc (+ 20 mV) strongly denote the positive impact of such a device [START_REF] Jamarkattel | Improving CdSeTe Devices With a Back Buffer Layer of CuxAlOy[END_REF]]. Nevertheless, the continuity of the film is not confirmed and further investigations are needed.

Reducing CdSeTe absorber thickness: aim and challenges of this work

As described in the introduction chapter, reducing the CdTe thickness would be beneficial in terms of cost, and to decrease the relatively high bandgap-voltage offset. It usually comes at the expense of light absorption, and requires an efficient light-trapping strategy with a highly reflective mirror such as silver to preserve the J sc . Thus, we add a new constraint for the back contact: it must be selective, passivating, and highly reflective.

In the following, we explore alternative back contacts for CdSeTe/CdTe solar cells compatible with highly reflective and nanostructured back mirrors for different absorber thicknesses. We start by studying several back contacts for standard CdSeTe/CdTe devices and the effect of passivation by Al 2 O 3 . A p-type transparent conductive oxide (TCO) layer is investigated as an alternative and efforts to overcome low efficiency are developed. Afterwards, experimental results on thinner cells are presented and compared to standard devices and simulations.

Conclusion of the section

The main limitation of CdTe-based solar cells is attributed to the low V oc . Recent models predict efficiencies as high as 28%, but to achieve such performance experimentally, back contacts should be both passivating and selective. Many cells have front/back passivation layers that increase their intrinsic properties, yet the positive impact on the figures of merit FF and V oc is not always observed. Moreover, there is no consensus in the community to explain the mechanisms at stake behind such performance with oxide passivation layers.

Back contacts for CdSeTe/CdTe solar cells

In this section, we report on the fabrication of CdSeTe/CdTe cells. The sample stack was provided by the NREL, and we developed the back contact deposition at the C2N. First, we present the detailed steps of the CdSeTe/CdTe device stack deposited at the NREL. Then, we present the performances of the cells with different metal back contacts tested with and without a passivation layer in between. Finally, an extensive study of Cu 2 O contacts is presented.

Solar cells fabrication process

Layer stack fabrication for standard and thin CdSeTe solar cells

This study is conducted with CdSeTe/CdTe solar cells with two distinct thicknesses of the absorber layer. We consider a standard CdSeTe/CdTe device with a 3450 nm thick active layer vs. a thinner one of 800 nm. The semiconductor layer and the solar cell layers stack were fabricated at the NREL. The layers with target thickness are described in Table 5.2. Note that it also underlines where the process steps differ for the two stacks (otherwise, the fabrication is strictly similar). The cells are fabricated in a superstrate configuration (the light is incident through the glass) using vapor transport deposition. The steps are shown in Figure 5.5.

-A soda-lime glass (commercial Pilkington TEC12D) is coated by chemical vapor deposition with fluorine-doped tin oxide, SnO 2 :F (FTO) and an intrinsic buffer i-SnO 2 .

-An MZO layer is then deposited by radio frequency sputtering using a mixed-oxide target with a weight ratio of 4wt% and 96wt% for MgO and ZnO, respectively. It is deposited at room-temperature under a flow ratio of 1% O 2 :Ar (see Figure 5.5(a)).

-The CdSe layer is thermally evaporated, immediately followed by the CdTe without having to break the vacuum in between. The substrate temperature is kept at approximately 450℃ for both evaporated materials during the deposition (see Figure 5.5(b)).

-The samples are CdCl 2 treated at 450℃ for 10 min in an environment of 400 Torr He (see Figure 5.5(b)).

-Lastly, the samples are dipped into a 0.1 mmol CuCl 2 solution for 3 min and annealed in air for 25 min at the temperature specified in Table 5.2 to ease the diffusion. The layer stacks provided by the NREL are 3.8x3.8 cm After the CdCl 2 annealing, the CdSe/CdTe layers form a graded CdSe x Te 1-x alloy (see Figure 5.5(b)) which will differ for standard and thin films. The SEM image in Figure 5.5(b), highlights the grain composition of CdTe. This particular stack gives the best efficiencies achieved at NREL to date [START_REF] Ablekim | Thin-Film Solar Cells with 19% Efficiency by Thermal Evaporation of CdSe and CdTe[END_REF]]. Nevertheless CdCl 2 and CuCl 2 annealing temperatures required some adjustments for standard and thin films to reach the best performance. The typical CdCl 2 temperature for standard CdTe solar cells is around 450℃, but multiple temperatures were tried for thinner cells. Figure 5.6 shows JV curves measured at NREL for several cells after they deposited gold back contacts. Figure 5.6(a) highlights that some inhomogeneities between cells can occur during the same process batch for standard CdSeTe/CdTe samples. Although the J sc remains identical for all the devices, the difference is particularly noticeable for the V oc and the FF. For thin CdSeTe/CdTe devices, CdCl 2 annealing temperature was held lower than the typical ones for thicker CdTe stacks to match the amount of CdCl 2 that diffuses to the front interface. From Figure 5.6, we clearly see that the lower the CdCl 2 temperatures, the lower the J sc . The low current is likely due to incomplete interdiffusion of CdSe with CdTe, and an incomplete passivation of bulk defects. From 370℃ to 380°C only the J sc varies. Note that for both types of sample (standard and thin ones), the CuCl 2 step was not yet wholly optimized either and the temperature was held at 200℃.

Back contacts fabrication

Before depositing back contacts, a gentle HCl etch (1:10) is conducted for 1.5 minutes to remove surface oxides without substantially affecting the CdTe. Metallic back contacts are deposited through a physical mask by e-beam evaporation. They consist of an array of seven metal squares regularly spaced, with a size of 0.3 x 0.3 cm 2 , as pictured in Figure 5.5(c). The CdSeTe/CdTe samples provided by the NREL offer a surface of 3.8x3.8 cm 2 . To optimize space and expand the number of possible back contact tests, the contacts are arranged horizontally, and each line corresponds to a different type of back contacts. No mechanical scribing is done to isolate each cell, and CdTe is left between the devices. The latter may result in a small collection outside the gold contact area (probably no more than 1-2 mA/cm 2 ).

The impact of a passivation layer is also discussed in this section. The Al 2 O 3 passivation layer is deposited by ALD. A horizontal band on the sample is isolated by photolithography, leaving the rest of the sample protected by the resin. This is followed by the Al 2 O 3 deposition at low temperature (120℃). This choice of temperature is justified by two reasons. First, our recommended annealing/deposition temperatures should not go above 200 ℃, because the CuCl 2 treatment at the back side is performed at 200 ℃ at NREL. Too high temperatures could induce detrimental copper diffusion in the sample. Secondly, high ALD temperature would have a lasting impact on the resin, and it would be hard to remove it afterwards. A metallic back contact evaporation follows the Al 2 O 3 deposition. The resin is removed from the CdTe surface in a remover-PG bath heated at 80℃.

In the following, standard CdSeTe/CdTe devices will be referred as A201014Ax-Axx and thin ones as A201016Ax-Axx, x corresponds to the number of the sample and xx is the number of the cells, numbered in ascending order from right to left in the horizontal line.

Back contacts for CdSeTe/CdTe solar cells

CdSeTe/CdTe samples with two nominal thicknesses have been fabricated and optimized at the NREL: one standard device with nominal thickness as of 150/3300 nm and a thinner one of 150/650 nm. The metallic contacts are deposited by e-beam evaporation at the C2N, with or without Al 2 O 3 deposited by ALD. A band on the sample presents a set of 7 cells with the same type of back contacts.

Highly reflective back contacts: Au and Ag

To compensate for the absorption loss in thinner CdSeTe/CdTe devices, highly reflective materials are considered at the backside. However, because of CdTe high electron affinity and low doping, replacing gold contacts while keeping high performances is not trivial. First, we investigate and compare Au and Ag back contacts on standard CdSeTe/CdTe (3450 nm-thick) solar cells.

Current-Voltage characterization for CdSeTe solar cells

In Figure 5.7, the JV curves (under one sun illumination and in the dark) of the best Cd-SeTe/CdTe solar cells fabricated with Au and Ag back contacts are plotted (among seven cells for each type of contacts). In Table 5.3, the results obtained at C2N are compared to equivalent devices performed at NREL and to CdTe devices fabricated at Colorado State University (CSU) with direct gold and silver back contacts, respectively.

Standard device with Au contact reveals promising performances even though the JV trend displays a strong rollover (JV curve bending around 0.8 V in the dark and under illumination) and a crossover (the curve under illumination crosses over the curve in the dark). It denotes the formation of a barrier at the CdTe/metal interface, suggesting a lower fill factor than typical values. Niemegeers et al. demonstrate in their early work, an analytical theory linking the back contact barrier height Φ b (see Figure 5.3) to the rollover in thin film CdTe solar cells. They also argue that the crossover might be caused by the surface recombination of electrons at the back contact [START_REF] Niemegeers | Effects of the Au / CdTe back contact on IV and CV characteristics of Au / CdTe / CdS / TCO solar cells Effects of the Au / CdTe back contact on IV and CV characteristics of Au / CdTe / CdS / TCO solar cells[END_REF]]. The silver contact, studied for its valuable reflective properties, has a significant J sc almost as good as the gold contact. Nonetheless, V oc and FF are deteriorated, and the curve displays an important S-shape, lowering the performance of the cell overall. Demtsu et al. show that the rollover is more pronounced, and FF is significantly reduced as the barrier height increases [Demtsu et al. 2006b]. This is what we observe for gold and silver back contacts. We estimate the barrier height for each metal contact using the following equation [Demtsu et al. 2006b]:

J bo = qv r N v exp(-qΦ b /(kT )) (5.4)
where J bo is the saturation current (the current value where the rollover occurs) and estimated graphically at about 6 mA/cm 2 and 0.1 mA/cm 2 for gold and silver, respectively. v r =3×10 7 cm/s is the Richardson velocity, N v the effective density of states in the valence band (1.8 × 10 19 cm -3 , [START_REF] Chen | The Influence of Conduction Band Offset on CdTe Solar Cells[END_REF]), q is the electronic charge, kT stands for the thermal energy.

We calculate a barrier height of 0.54 eV for Au and 0.64 eV for Ag. Regardless of the back contact type, FF values are responsible for the efficiency drop when comparing our cells to those of the NREL or the CSU (Table 5 [START_REF] Pandey | Metal Oxides as buffer layers in polycristalline CdTe thin-film solar cells[END_REF]] (CSU-CdTe) and painted [Demtsu et al. 2006a] (CSU-CdS-CdTe). The J sc was extracted from the EQE measurements (see Figure 5.11).

The spectral response and collection efficiency of both cells are investigated, see Figure 5.8. The external quantum efficiency (EQE) is measured in standard conditions (0 V) and under different bias voltages between -0.4 V and +0.4 V. The ratio of the EQE measured at short-circuit condition and at a bias -0.4 V, i.e. EQE(-0.4V)/EQE(0V) is very close to 1 over the whole wavelength range for the Au contact. A good collection efficiency is achieved. Nevertheless, at forward bias the ratio EQE(0.4V)/EQE(0V) is closer to 0.9 with no significant change over the whole wavelength range. This 10% drop corroborates what we observe in Figure 5.7, indeed there is slight slope visible on the yellow JV curve between 0 V and 0.4 V, which might be linked to shunt resistance or to a decrease in collection efficiency. Ag contacts are more impacted by the collection efficiency barrier. At positive bias and from 0.2 V, we see its significant impact through a degraded EQE(0.2V)/EQE(0V) ratio close to 0.8. More importantly, the collection efficiency is improved at reverse bias: an EQE similar to the Au contact is recovered at -0.4 V, which explains the J sc loss between Au and Ag back contacts.

Light soaking impact on solar cell performance

Nearly all the polycrystalline thin film PV technologies undergo changes under long illumination exposure, and CdTe cells are no exception. Studies on the effect of light soaking for CdTe solar cells and modules demonstrate instabilities in the performances [START_REF] Del Cueto | Long-term Transient and Metastable Effects in Cadmium Telluride Photovoltaic Modules[END_REF]]; [START_REF] Gostein | Light soaking effects on photovoltaic modules: Overview and literature review[END_REF]; [START_REF] Silverman | Performance stabilization of CdTe PV modules using bias and light[END_REF]]. Nevertheless, when the performance of CdTe cells improves under light soaking, it has been associated with Cu ions migration [START_REF] Guo | A Comprehensive Study of Light Soaking Effect in CdTe Solar Cells[END_REF]. A similar study with CdSeTe/CdTe has shown significant improvements under light soaking. Jamarkattel et al. demonstrate that with a 120 nm-thick CdSe layer and after 16 hours under illumination, the efficiency rises from 13.5 to 16.5%, showing more intense PL emission after light soaking [START_REF] Jamarkattel | Reduced Recombination and Improved Performance of CdSe/CdTe Solar Cells due to Cu Migration Induced by Light Soaking[END_REF]. It also demonstrates a strong correlation between the CdSe thickness and the effect on light soaking: with 150 nm-thick CdSe the performances are degraded. .9(a) shows light and dark IV curves for several light soaking times for CdSeTe/CdTe solar cells with a gold back contact. The main observation is that longer illumination seems to attenuate the rollover. The results conducted on seven cells are summarized in the box plot Figure 5.9(b). From 0 min to 10 min exposure, the J sc and FF values slightly increase and seem to stabilize after longer illumination. On the other hand, the V oc decreases until an exposure of 10 min. A higher temperature during exposure might explain this behavior, even though no dramatic temperature rise was observed. The improvement in efficiency remains relatively small (e.g., ∆Eff =0.5%); nonetheless it demonstrates the positive impact of a longer exposure before measurement.

The temperature of the cells is rising when they are soaked in light. It might impact negatively on the performance and as the temperature was not monitored precisely, no longer tests were conducted in our case. However, collaborators from NREL have tested more extensive light soaking experiments. Longer times were tested (≈ 2 hours) with temperatures either maintained at 25℃ or done without intentional heating/cooling (resulting in ≈ 60℃) before JV measurements. They confirm that with an open-circuit light soak of about 5 min slight improvements in device performance are sometime observed. For longer soaking time, they do not notice significant differences between the two temperatures.

Au contacts have shown promising results, with correct J sc and V oc . The low FF due to the strong rollover effect may be an indication of poor contact during the measurement. As for Ag contacts, a good J sc is found but the V oc and FF values suffer from a bad collection efficiency. Light soaking has proven to increase slightly the cells performance.

Role of Al 2 O 3 passivation layers

In addition to the previous back contact experiments, the impact of a passivation layer is studied. Several thicknesses of Al 2 O 3 passivation layer deposited by ALD have been examined in order to investigate the trade-off between the passivation and the hole collection by tunnel effect. The performances are compared with Au and Ag back contacts only. The best JV curves under one sun with different thicknesses of Al 2 O 3 are compared in Figure 5.10(a) and 5.10(b) with Au and Ag metallic back contacts, respectively. A summary of the performances measured for seven cells (except for the cells with 1 nm Al 2 O 3 + Au contact, for which only three cells were characterized) and the best cell parameters are shown in Figure 5.10(c-f). The J sc values are extracted from EQE measurements. Typical EQE curves obtained with our CdSeTe/CdTe cells are plotted in Figure 5.11.

Au metallic back contact

First, 1.5 nm of Al 2 O 3 is deposited at the rear side of the stack. The V oc values are slightly improved, and the data dispersion is reduced. However, the FF and J sc suffer from this additional layer, contributing significantly to decreasing the efficiency. This is likely linked to a high hole collection barrier generated at the interface. The Al 2 O 3 layer thickness was thus reduced to 1 nm.

Hence, with a very thin layer of alumina, JV curves similar to the reference are obtained with values as high as the cells with gold contact only. Interestingly, the rollover is more important when there is a passivation layer in between which indicates that the back contact barrier might be higher than without the latter. It appears that as the alumina thickness is higher than 1 nm, the electrical parameters start to degrade, and the passivation layer is detrimental to the proper operation of the cell. The carrier transport by tunneling is ensured for an Al 2 O 3 thickness equal or less than 1 nm. While the Al 2 O 3 passivates the back interface (more details in section 5.3.2) and reduces detrimental recombination, only a thin layer is suitable for a decent current extraction.

Ag metallic back contact

1 nm of Al 2 O 3 with Ag contacts results in a strong efficiency drop (less than 2%). Still, the V oc undergoes a significant increase of almost ≈ 150 mV compared to the device without any passivation layer. However, looking at the JV curve shape, where the FF is hardly quantifiable, the V oc values must be considered with care. Regarding J sc values coming from JV curves, we observe they are capped at 7 mA/cm 2 . However, these results are inconsistent with the ones extracted from EQE measurements. In Figure 5.11, even though the EQE slightly decreases for the Al 2 O 3 +Ag contact (blue curve) compared to others, no significant drop is visible. The corresponding J sc are estimated at about 25 mA/cm 2 , close to the values found for contacts without a passivation layer. The EQE measurements are performed at much lower illumination intensities, which may explain the higher collection efficiency and the different electrical behavior. This assumption will be further discussed in subsection 5.2.4. As for Au contacts, the rollover is more pronounced with Al 2 O 3 , but this time the FF is strongly impacted as well. Contrary to our expectations and to what can be suggested in the literature, the passivation layer does not improve the Ag back contact. J sc -V oc characterization In Figure 5.12, we construct a pseudo-JV curve (black curve) in comparison with the JV curve measured directly under a 1-sun light (yellow dots) for an Au contact with or without a passivation layer (Figure 5.12(a) & 5.12(b) respectively)). Each black circle corresponds to a couple (V oc , J sc ). The incident light intensity is varied on a cell, and V oc and J sc are measured at each illumination level.

These data points are plotted with a vertical shift of the one sun value of J sc (see equation 5.5). We fit these points with a 2-diodes model (equation 1.4), which gives rise to the black curve. The pseudo-JV curve is a convenient way to study the intrinsic behavior by putting aside series resistances and barrier limitations. It is unaffected by the voltage drop due to series resistance (no current is drawn from the cell). Hence, comparing the two curves at a given current, the voltage drop is related to the series resistance of the device. pseudoJ = J sc @1sun -J sc (5.5) From Figure 5.12(a), we can extract roughly R s ≈ 8 Ω.cm 2 with an Au contact, far from the R s ≈ 34 Ω.cm 2 obtained from the dark JV fit. To extract a more accurate R s value and evaluate the potential impact of the passivation layer, we go even further in the analysis. First, the pseudo-JV data are fitted, then the JV under 1 sun illumination is fitted with the same parameters but varying only the series resistance (and n 2 for better precision). The parameters are listed in Table 5.4. We obtain convincing fits for both devices. However, the rollover effect is not taken into account in the JV sun fit.

The same series resistance obtained for both types of contacts demonstrates that the passivation layer does not induce detrimental resistance losses at the interface. A pseudo-FF of 78% can be evaluated from Figure 5.12(a). In comparison, the real FF is 57%. The 21% in difference in FF is attributed to the resistive loss and the contact barrier at the CdTe/metal interface. The pseudo-FF obtained with the passivated device reaches 80%.

J 01 (mA/cm 2 ) n 1 J 02 (mA/cm 2 ) n 2 R s (Ω.cm 2 )
Au contact only, pseudo-JV A CdSeTe/CdTe cell with a very thin layer of Al 2 O 3 below the Au contact keeps the same performance than without, and even allows a slightly better V oc . The hole collection is done by tunneling. On the other hand, it deteriorates the efficiency of cells with Ag contacts. A significant difference between the J sc determined with JV curve or EQE is observed for the latter. A pseudo-JV characterization highlights that without the resistive losses and the contact barrier the FF could be as high as 80%.

P-type transparent conductive oxide as a back contact : Cu 2 O

In section 5.1, we have highlighted the necessity of highly selective back contacts to overcome the 1 V barrier of CdTe-based solar cells. To continue to take advantage of the high reflectivity property of gold or silver, these selective contacts should be transparent. In the ideal case, materials such as p-type semiconductors with a wide bandgap or p-type transparent conductive oxides (TCO) meet these requirements. However, having a large gap and highly-doped semiconductor is very complex and there are very few intrinsically p-type TCO [START_REF] Zhang | P-type transparent conducting oxides[END_REF]]; [START_REF] Banerjee | Recent developments in the emerging field of crystalline p-type transparent conducting oxide thin films[END_REF]]; [START_REF] Kawazoe | P-type electrical conduction in transparent thin films of CuAlO2[END_REF]]. Among these few p-type TCO, NiO and Cu 2 O are suitable candidates.

In the following, we use the Cu 2 O as a stable, non-toxic and potentially selective back contact. This promising material acts as a barrier for the generated electrons (electron reflector similar to ZnTe back buffer) decreasing the recombinations at the interface resulting in high V oc . Moreover, it could lead to higher voltage for lower absorber layer thickness [START_REF] Pettersson | The influence of absorber thickness on Cu(In,Ga)Se2 solar cells with different buffer layers[END_REF]; [START_REF] Liu | Numerical simulation of ultra-thin CdTe solar cells with a buffer layer of MoOx in the backwall configuration[END_REF]]. The only experimental study using a Cu 2 O back contact was conducted by Türck et al., and they obtained efficiency higher than 15% [START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF]. Besides, numerical modeling has yielded promising results (efficiency of almost 20%) despite a CdS window layer and hole mobility of 25 cm 2 /(Vs) [START_REF] Tinedert | Numerical modelling and optimization of CdS/CdTe solar cell with incorporation of Cu2O HT-EBL layer[END_REF]]. Here, in collaboration with D. Munoz-Rojas and A. Sekkat, we use a high-quality Cu 2 O material deposited by spatial atomic layer deposition (SALD) at atmospheric-pressure at the Laboratoire des Matériaux et du Génie Physique (LMGP). It is chosen for its attractive parameters: it has a high hole conductivity of 92 cm 2 /(Vs) and wide optical bandgap, which can be tuned between 2.08 eV and 2.5 eV, depending on the deposition temperature [START_REF] Frouin | Ultrathin CdSeTe solar cells with optimized rear side contact[END_REF]]. Prior to Cu 2 O deposition, several steps were handled. The process is described in Figure 5.13(a). A gold layer was deposited on the TCO to ease the connexion of the front contact during measurement. A SiO x layer was sputtered on the front contact and on a small part of the CdSeTe/CdTe rear side to prevent any residual deposition of Cu 2 O on TCO. The influence of the Al 2 O 3 layer is also studied in this particular case. Hence, 1 nm was deposited by ALD on half of the sample for its passivation potential and to prevent detrimental copper diffusion in the active layer as the Cu 2 O deposition reached temperature of around 260℃. Note that this temperature, while very low for such a deposition, exceeds by 60℃ the recommended one to prevent fast copper diffusion in CdTe. The Cu 2 O layer was deposited by SALD at the LMGP after a rapid deoxidation in dilute HCl (1:10) for 1 min. Cu 2 O deposition and SiO x are removed from the front contact by chemical etching and reactive ion etching, respectively. Metallic back contacts are deposited afterwards. Two different architectures are compared in the following, one with Cu 2 O deposited directly on the CdTe surface another with 1 nm Al 2 O 3 .

Cu 2 O deposition

Solar cells performance characterization

The characterization of cells with a p-type oxide contact is shown in Figure 5.14. The two architectures are compared through IV and EQE with Au contacts (a,c) and Ag contacts (b,d). With a dramatically low J sc , those first results suggest a transport problem linked to a relatively high resistivity of the Cu 2 O layer and/or a barrier at the CdTe/Cu 2 O interface. Unlike the reference cells with a standard gold contact at the rear side, no crossover is observed.

Overall, the passivation layer seems to have a rather detrimental effect on the performances, even though the s-shape is attenuated. The J sc appears to be the most impacted parameter as shown in Figures 5.14(c,d) with a surprising shape from 800 nm to 900 nm. However, it is hard to conclude on its influence on the FF or the V oc since different behavior can be observed depending on the metal back contact (Au or Ag). Once again, as observed in part 5.2.3 for Ag contact with a passivation layer, the J sc extracted from the JV or the EQE measurements does not match regardless of the cells measured. Figure 5.15 compares the J sc extracted from EQE and JV curves for different contacts. With 1 nm of alumina, we observe a high disparity in J sc values (from EQE, plain bars and light colors) depending on the position of the cells on the sample (P1 to P3 for the Au contact, visible in Figure 5.13(c)). This suggests a heterogeneity in the Cu 2 O or in the CdTe layer with an optimum in the middle of the sample. Note that the part of the sample with the passivation layer would also be more sensitive to a non-homogeneous Cu 2 O thickness deposition. This subsection emphasizes two major issues in the devices with a Cu 2 O back contact. On the one hand the performances obtained with this Cu 2 O back contact are dramatically inferior to a CdTe device with a classical back contact. On the other hand, the J sc difference between EQE and JV data have not been justified so far. These two outcomes will be addressed in the following.

Attempt to improve CdSeTe solar cells with p-type oxide as back contact

Two main approaches are considered to improve the efficiency of CdSeTe/CdTe devices with Cu 2 O as back contact: the annealing of the devices and the thickness reduction of the Cu 2 O layer.

Effect of the annealing

Türck et al. have deposited Cu 2 O by reactive sputtering on CdTe/CdS cells, using gold as metallic back contact [START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF] . They observed an improvement in the device performance after a 150 ℃ annealing step conducted for 5 min as shown in Figure 5.16. Their s-kink disappears after treatment. They attribute this behavior to an improved contact linked to changes in the doping levels.

Voltage (V) Current (mA/cm²) We tried a similar annealing treatment for our devices. The sample with Cu 2 O and Au-Ag metallic back contacts was annealed on a hot plate at 150 ℃ in air for 5 min. Corresponding JV curves and EQE are plotted in Figure 5.17.

J sc , V oc , FF and R s before and after the annealing treatment are shown in Figure 5.18 as a function of the cells position on the sample. Part of the sample with cells with an alumina passivation layer is highlighted in light green. After the annealing, we observe no clear improvements. Figures 5.17(a,b) show a slight modification and the kink for annealed devices is attenuated. However the J sc is still reduced. This J sc decrease is confirmed by the EQE measurements (Figures 5.17(c,d)). Regardless of the metal back contact used, the resonances in the interval [800 -900] nm are more marked.

Figure 5.18 confirms a J sc reduction after the annealing for most of the cells. Interestingly, after the annealing, the FF is constantly enhanced. Leaving asides the position P3, the FF is balanced at around 30% with gold back contacts. It remains more dispersed with silver contacts. No evident trend can be extracted for V oc however, the annealing treatment tends to decrease it overall. It is worth noting that the V oc is less impacted by the type of metal contact used with an oxide in between. Series resistance, evaluated by fitting with a 2-diodes model dark JV curves and plotted in logscale, are very scattered and sometimes displays dramatically high values. Lastly, the passivation layer does not seem to make a difference even with a supplementary annealing step. Nevertheless, higher FF values are obtained with passivated devices even though V oc and J sc are much lower compared to unpassivated cells.

Finally, the device at position P3 stands out from the crowd with an unexpected behavior with relatively high J sc and V oc but low FF, unlike the cells in positions P1 and P2. The low FF could be explained by the corresponding high R s and the surprising results by the low homogeneity of the Cu 2 O layer, as already mentioned in subsection 5.2.4.

CdSeTe/CdTe with a Cu 2 O oxide p-type contact have not resulted in high efficiency devices. Besides, no drastic improvements arised from an annealing treatment in air. A 130 nm-thick Cu 2 O layer might be too thick. While its carrier mobility is relatively high for such a material, its series resistance might be excessive. In the following, we reduce the Cu 2 O layer thickness and study once again the influence of an additional annealing step.

Reducing Cu 2 O thickness and annealing impact

For the sake of a good homogeneity of the layer at the rear side of CdSeTe stack, the Cu 2 O was physically etched by reactive ion etching using CHF 3 at -350 V bias, for 20 min. A layer of approximately 50 nm is left at the back of the sample. As seen in Figure 5.19, a deep green color has replaced the purple shade at the CdTe surface after the etching. Figure 5.19(b) shows a SEM image of the surface of the sample after 6 min of etching. There is a visible surface modification compared to Figure 5.19(a), as if the Cu 2 O had melted. After the layer reduction, gold contacts are deposited at the back side. It is followed by a 5 min annealing step on a hotplate at 300 ℃ in the air. Solid lines in Figure 5.20 display similar behavior as plotted in Figure 5.17, even with a thinner Cu 2 O layer. The drop in J sc might be linked to the structural degradation during the Cu 2 O dry etching. Although J sc remains low, adding a 300℃ annealing step allows a recovery of decent FF and V oc values: a global increase of about 350 mV for the V oc is achieved, (see Table 5.5). No improvement due to the passivation layer is observed.

To assess a possible degradation of the cells due to the high annealing temperature, another set of measurements was conducted two weeks later. Two cells were shunted, and the J sc , V oc of the remaining were slightly reduced. Interestingly, half of the cells without Al 2 O 3 did no longer work, while the passivated ones were operating, with the same V oc as before. This passivation layer might preserve the cells from detrimental copper diffusion resulting in shunted cells.

Devices performances summary

Table 5.5 summarizes the performances of CdSeTe/CdTe cells with a Cu 2 O oxide layer and a gold back contact. It shows the results before and after the Cu 2 O etching and the annealing steps. While the devices with a p-type oxide back contact are not as efficient as reported by Türck et al. [START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF], we still report high J sc . Annealing slightly reduces the shortcircuit current, but corrects the s-shape, increasing at the same time the FF. The cells with a 50 nm thick Cu 2 O have poor performances, but a recovery of the V oc is observed after a 300 ℃ annealing.

Sample name and description J sc (mA/cm 2 ) IV J sc (mA/cm From these first results, we present some suggestions for future experimental work:

-Depositing a thinner Cu 2 O layer (less than 50 nm) by SALD to assure a decent J sc . Surface treatment just before the oxide deposition could be beneficial.

-An annealing step is required to recover a usual IV shape and a good V oc . Temperatures in the range [150 -200] ℃ shall be preferred to prevent damages. The annealing time has also to be defined.

-A Cu 2 O/Ag back contact remains a relevant alternative since similar results are obtained regardless of the metal back contact in Figure 5.18.

External Quantum Efficiency assessment

Inconsistent characterization results have been observed between EQE and JV curve for several configurations such as with Cu 2 O, and Ag+Al 2 O 3 back contacts (subsections 5.2. 3 and 5.2.4). While EQEs show a high absorption, poor J sc are extracted from corresponding JV curves. To glean insights on this disparity we conducted experiments on devices with Cu 2 O layer at the back side specifically.

PseudoJV & JV under monochromatic, low intensity illumination

As the intensity of the monochromator light used for EQE is much lower compared to the one from the one-sun simulator, one hypothesis is that our CdSeTe/CdTe cells behave in a different way depending on the light intensity. To test this hypothesis, we measured IV characteristics with several neutral optical densities (OD) between the light source and the cells for Cu 2 O+Au back contact. Figure 5.21(a) shows the JV curves at different illuminations. We observe that the s-shape becomes less pronounced as the power decreases, and that the current at reverse bias is almost the same at V=0 V for very low illumination (green curve). Figure 5.21(b) shows the pseudoJ sc vs. V oc compared to JV curve under one-sun illumination. Interestingly, the slope near V=0 V is flat confirming a better collection efficiency at low illumination. Using a monochromator, IV curves are measured under specific wavelengths. Results are displayed in Figure 5.21(c). These JV curves show that the current is saturated from V=-0.4 V until V=-0.3 V, however the contact barrier around V=0.4 V is so important that it influences the collection around 0 V. This also confirms that at higher intensity (under 1 sun illumination), we never saturate the current, even at high reverse bias. The wavelength dependence of the JV curve is justified by the varying light power. Thus, the collection efficiency is modified during the measurement and is strongly related to the incident power. The resonances behavior between 750 nm and 850 nm can be justified by the spectral response of the xenon lamp of the monochromator.

We also observe that the collection efficiency is dependent on the back contact. When adding a passivation layer or using a silver contact, it increases the barrier and with a low V oc the carrier collection issue is even more visible. The reverse applied voltage should be higher to expect to collect all photogenerated charges. Although we still have to conduct experiments to better understand the mechanisms at stake, we evidenced that the low J sc from JV curves is due to a low collection efficiency of charges which depends on the density of photogenerated charges.

EQE and beam size variation

Further investigation led us to examine the influence of the monochromator beam size on the EQE data. 5.22(a) and (b) show that higher EQE are measured with a large beam that covers the entire cell than with a small focused beam, and regardless of the presence of the passivation. Annealed samples exhibit similar outcomes. For comparison, EQE measured with only a Au back contact shows almost no dependency on the beam size or displays a slightly improved curve for a smaller beam (Figure 5.22(c)). Finally, in Figure 5.22(d), we plot EQEs at several positions on the cells (blue dashed curves) and compare it to the EQE obtained with a large beam (solid black line). The significant dependence on the beam position suggests inhomogeneously distributed defects.

Conclusion of the section

In this section, the performance of standard CdSeTe/CdTe cells with a Cu 2 O back contact was reported. It is one of the few p-type TCO and was chosen to act as a selective contact layer. Cu 2 O deposited by SALD have a hole conductivity of 92 cm 2 /Vs. Cells with asdeposited Cu 2 O have poor performances, and the Al 2 O 3 layer deteriorates them further. Regardless of the type of the metal contact (Au or Ag), the EQE reveals decent J sc , but still very dependent on the position of the cells on the sample. An annealing step at 150°C on the same cells decreases systematically the J sc , but leads to a small improvement of the FF. Again the performance is related to the place of the cells on the sample. After reducing the p-type TCO thickness, the J sc is negatively impacted which might be due the deterioration of the Cu 2 O layer by reactive ion etching. However, a 300°C annealing helps to recover a higher V oc . The Al 2 O 3 layer prevent the shunting of the cells after annealing. A study was conducted to explain the different responses of the cells for IV and EQE measurements. This highlights the dependence of some cells on the illumination.

Analysis of 800 nm-thick CdSeTe solar cells

Returning to the primary goal of reducing the CdSeTe/CdTe absorber thickness, we carry on the back contact study on thinner 800 nm CdSeTe/CdTe cells. In this section we present our characterization results and compare the performance with standard devices. To glean more insight into the material properties, photoluminescence data are used. Afterwards, the spectral response and optical measurements of standard and thinner samples are compared with optical simulations. They provide a reliable basis for refining the optical models.

Solar cells performance characterization

JV curves and EQE of the best thin CdSeTe/CdTe solar cells fabricated with Au and Ag back contacts are plotted in Figure 5.23. The JV curve in Figure 5.23(a) with a Au contact reveals encouraging performances with V oc = 0.67 V and FF= 66%, values comparable to standard devices, with no rollover (see Figure 5.7), and only a small a crossover. The EQE results in a high and homogeneous J sc value around 28 mA/cm 2 . Dark JV curve for a device with Ag back contact (dashed grey curve) has an unexpected shape with negative values. Nevertheless, correct J sc (evaluated from JV under 1 sun illumination) and FF are obtained overall since no kink is visible. In contrast, the V oc is even more deteriorated than the one of standard CdSeTe/CdTe cells, ∼150 mV compared to ∼550 mV, respectively. The EQE for Ag contacts shows artifacts at short wavelengths with EQE higher than 1, and at long wavelengths with an unexpected signal above 1000 nm. The EQE shape suggests a strongly illumination dependence, which confirms a complicated charge collection due to the contact barrier. Devices statistics are given in Figure 5.24, calculated for seven devices for each group of cells. Note that the J sc are extracted from EQE except for the cells with Ag contact only where the values come from the JV data due to inconsistent EQE data. Table 5.6 summarizes the performances of the best thin cells obtained at the C2N with different metallic back contact and compares them to the cells obtained at the NREL and to CdTe cells with the same absorber thickness found in the literature. Fewer disparities are observed between our devices and the ones from NREL than for standard CdSeTe/CdTe cells. We have the same J sc , but V oc and FF values are still below what we expected. Our best cell with an Au contact has increased J sc compared to CSU due to the 150 nm-thick CdSeTe layer, which lower the bandgap. Although it still suffers from low V oc , the two resulting efficiencies are comparable. [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF].

Once again, we study the impact of an Al 2 O 3 passivation layer between the active layer and the back contact. We limit its thickness to 1 nm. Results are summarized in Figure 5.24. JV curves in Figures 5.24(a) and (b) from cells with a passivation layer now display a s-shape, even more pronounced with the silver contact. It results in a drastic reduction of the FF, detrimental to the efficiency. However, Al 2 O 3 greatly improves the V oc for cells with Ag back contacts as observed with standard devices. Note that the increased J sc with the passivation layer might only be related to its type of extraction (EQE vs. IV). Adding an alumina layer with a silver back contact removes the EQE artifact visible in Figure 5.23(b) (blue and grey curves, respectively). As for standard devices, for Al 2 O 3 +Ag back contact, we observe the same gap concerning the J sc value between EQE and IV. Globally, thin CdSeTe/CdTe appears to be more negatively impacted by the passivation layer than the standard one.

JV curves under 1-sun and EQE for Au back contacts comparing the two different thicknesses of the absorber layer are plotted in Figure 5.25(a) and (b), respectively. We do not have a drastic J sc drop for thinner cells (light blue curve), which encourages decreasing the active layer even further. However, we observe a systematic decrease in the V oc . Remarkable EQE is obtained for thinner cells. As expected the absorption of the thicker sample (Figure 5.25(b)) slightly outperforms the thinner one at high wavelengths from 700 nm to 900 nm. The transmission measurements shown in Figure 5.25(c) support the EQE results. The intensity drop occurs near the CdSeTe bandgap for both devices. The progressive decrease of the transmission signal in the long wavelength for the thin absorber highlights the role of the absorber thickness. The two peaks at 400 nm and 500 nm are unexpected attributed to artifacts.

Overall, the performance of both thicknesses with Au contacts are comparable, the J sc are similar but lower V oc is observed for thin CdSeTe/CdTe solar cells. Unlike thicker cells, the JV curves of thinner CdSeTe/CdTe devices exhibit no rollover nor crossover. On the one hand, these results demonstrate that reducing the absorber layer by more than three times its usual thickness with good quality materials does not introduce significant losses in the absorption nor in FF or V oc deficit. On the other hand, this reveals an heterogeneity in the CdSeTe/CdTe layers provided by the NREL.

The difference between the performances of the C2N and the NREL cells (see Tables 5.6 and5.3) may be explained by aging effects, since several weeks/months have passed before the back deposition at the C2N. The difference could also originate from the significant heterogeneities observed in these samples. 800 nm-thick CdSeTe/CdTe cells have shown promising efficiency with Au back contact. They are more sensitive to the addition of the Al 2 O 3 passivation layer, and the efficiency is reduced. Remarkable EQE is obtained demonstrating a high quality material.

Photoluminescence response

To investigate on the difference between thin and standard CdSeTe/CdTe samples, and gain insights on the charge recombination in the absorber layers, we performed photoluminescence measurements at the back side on both samples. The resulting PL spectra are plotted in Figure 5.26.

Thin and standard film comparison

The PL peak is located at 825 nm (1.50 eV) and 880 nm (1.41 eV) for standard and thin Cd-SeTe/CdTe cells, respectively. Those values are consistent with the CdTe and CdSeTe bandgaps and confirm the presence of CdSeTe [START_REF] Halliday | A photoluminescence study of polycrystalline thin-film CdTe/CdS solar cells[END_REF]]; [START_REF] Kraft | Comprehensive photoluminescence study of chlorine activated polycrystalline cadmium telluride layers[END_REF]]. The thin sample displays a broader full-width half maximum than the standard layer. The shifted blue curve of thin CdSeTe/CdTe layer testifies that most radiative recombinations occur in the CdSeTe region of the stack, while they occur mostly in CdTe for the standard device. It demonstrates a satisfying interface between the CdTe and the CdSeTe layers, thus a successful incorporation of CdSeTe. Moreover, the two devices display the same PL intensity, suggesting a high-quality material for the thin CdSeTe/CdTe sample. The laser beam at 532 nm is absorbed in the CdTe layer at a depth of approximately 150 nm. The spectral shift that occurs for the thin (650 nm CdTe) but not for the standard (3300 nm CdTe) devices gives us a rough approximation of the diffusion length, restricted to ∼3.1 µm.

Surface passivation

We also investigate on the surface passivation impact. Several studies have observed a significant PL intensity increase (by several orders of magnitude) when an alumina layer is added at the back of CdTe samples [Munshi et al. 2018a]; [START_REF] Danielson | Sputtered Aluminum Oxide and p+ Amorphous Silicon Back-Contact for Improved Hole Extraction in Polycrystalline CdSexTe1-x and CdTe Photovoltaics[END_REF]. Note that the PL measurements reported in these studies were conducted at the front side of the samples. Under our experimental conditions, the PL measurements indicate a fairly negligible passivation effect. This might be explained by the order of the experimental steps: in our case the alumina was deposited after the CdCl 2 annealing treatment.

Spectral responses and simulation comparison

Meanwhile, one of our next goal is to reduce further the active layer thickness to 200 nm and develop a light trapping structure to compensate absorption loss. To conduct such a study, we need to ascertain that a simulation model can reproduce experimental results. Spectral response and optical measurements were used to assist our calculations.

We use numerical computation based on the transfer matrix method (TMM). This enables the calculation of the reflection and transmission coefficients of a stack of thin planar layers, and the absorption in each layer. The refractive indices used for optical simulations of CdSeTe/CdTe solar cells are plotted in Figure A.11 in the Appendix A.3.1. As we have very few information about the commercial soda-lime TEC12D coated with FTO, we introduced a slightly modified stack model with corresponding optical indices. Our sample is in a superstrate configuration, and the incident light comes from a medium with a refractive index similar to glass: n glass = 1.5. Then, 30 nm SiO 2 are followed by 500 nm of fluorine-doped tin oxide (FTO). The FTO indices are taken from data gathered on a TEC7 substrate used for perovskite with corresponding thickness [START_REF] Raoult | Iterative method for optical modelling of perovskite-based tandem solar cells[END_REF]. The refractive index of the MgZnO buffer layer is approximated by the one of the ZnO layer. CdTe and CdSe x Te 1-x data are provided by the NREL and from reference [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]. Refractive index of Au is taken from [P. B. [START_REF] Johnson | Optical Constant of the Nobel Metals[END_REF]]. The corresponding scheme is plotted in Figure 5.27.

We launched a first trial of calculations using only a CdTe layer in our simulated stack and compared it to the EQE for both samples. In Figure 5.28, we compare the EQE data to the simulated absorption in the active layer. Our model offers already a convincing fit for standard CdSeTe/CdTe plotted in Figure 5.28(a). The EQE of the standard solar cell exhibits less pronounced resonances due to the spectral resolution of the measurements and the small roughness of the Cd(Se)Te layers, but similar averaged absorption is observed. Nonetheless, the inset displays the EQE obtained through FTPS measurements with a higher spectral resolution and resonances are visible in the long wavelength part of the spectrum. The high signal-tonoise ratio makes it possible to distinguish some resonances. Because of an inaccurate optical alignment, EQE measured with FTPS are not exploited: overall they display a much lower J sc compared to EQE measured with a monochromator. Although there is a good match for standard CdSeTe/CdTe device, the influence of the Se alloy is clearly visible for the 800 nm-thick sample (Figure 5.28(b)). Indeed, if the CdTe layer alone is considered, the simulated absorption curve drops at a wavelength of about 820 nm, at least 40 nm below the EQE. Then, absorption in the active layer of stacks with 650 nm of CdTe and 150 nm of CdSe x Te 1-x with varying composition in Se are compared in Figure 5.28(c). All the other parameters are kept constant. It can be seen that the higher the Se concentration, the better the fit. More specifically, the absorption in the active layer is very similar for the two highest Se concentrations, where the differences are only visible at short wavelengths. The alloy CdSe 0.424 Te 0.576 will be chosen hereafter. Again to improve the match, the relative thicknesses of CdTe and CdSeTe are varied while keeping the total thickness equal to 800 nm (Figure 5.28(d)). The best agreement between experimental EQE and numerical simulation is obtained when introducing a 300 nmthick CdSe x Te 1-x layer, with x se = 0.424. The simulated CdSeTe layer is twice thicker than the nominal thickness of the CdSe layer. This is due to the interdiffusion of Se on the CdTe layer, and we emphasize that this thickness of CdSe 0.424 Te 0.576 is consistent with the initial amount of Se elements. To keep the fit consistent, we consider the same CdSeTe thickness for the standard sample.

Using a FTIR spectrometer with a Si detector, transmission data (dashed curves) are compared to calculated and optimized stack in Figure 5.29. Figure 5.29(a) shows the simulated transmission for a 3.45 µm thick CdTe only and the one when CdSeTe is added. Unlike EQE for standard device, it highlights the key role of the CdSeTe layer as the transmission is slightly reduced at 850 nm (red curve) and fits better to the experimental data. For the thinner device, we compare the transmission of a bilayer CdSeTe/CdTe device to a 800 nm-thick CdSeTe device (see Figure 5.29(b)). The transmission peak expected at 900 nm is not observed in experiments. Although the transmission of the device with a CdSeTe layer only (solid blue curve) matches better to the measurements near its bandgap (∼ 880 nm), Figure 5.28(d) invalidates the assumption of such a stack. This observation and the nonetheless good match of the bilayer to EQE measurements suggest a good interdiffusion between the CdSe and CdTe layers. Unlike in Chapter 4, where the interdiffusion region was quite small between the CdSeTe and the CdTe region, here we may have a larger gradient since the fabrication process was different. We have a layer with an unknown Se concentration that varies with depth, and we demonstrate a good modelisation of the device response with a uniform layer of CdSe x Te 1-x with x = 0.424. Additional measurements by EDX, or CL would give us access to the gradient as it was done in the previous chapter, however this good agreement is sufficient for the following work.

Optimized simulated results are plotted in Figure 5.29(c & d) for the standard and thin devices, respectively, and compared to their corresponding EQE. The excellent match between measurements and simulations suggests a high quantum yield of the active layer. The absorption calculated in the CdTe layer only is plotted for both devices (solid purple and blue line in Figure 5.29(c), (d) respectively). Comparing this result to EQE measurement, we show that the CdSeTe is the layer that mainly dominates the absorption even though it represents only one twentieth of the active layer. It accounts for almost 75% of the current collection. For the thinner device, it represents 78%. A detailed current density loss analysis for CdSeTe/CdTe sample will be given in the next chapter.

Although, there is a satisfying match between the simulation and the experimental results, there is still room for improvement. Beyond 920 nm wavelength, the simulation gradually shifts from measurements due to imperfect optical indices. To improve our model, we could refine the refractive indices of the CdSe x Te 1-x , with a broader range of Se concentration (x), adjusting the concept developed for refractive indices of CIGS [START_REF] Carron | Refractive indices of layers and optical simulations of Cu(In,Ga)Se2 solar cells[END_REF]; [START_REF] Minoura | Optical constants of Cu(In, Ga)Se2 for arbitrary Cu and Ga compositions[END_REF]].

Conclusion of the section

Once tested on standard CdSeTe/CdTe devices, similar back contact depositions were carried out on thinner cells. Cells with Au contacts have a remarkable EQE, but performances are degraded when a very thin alumina layer is added. Cells with Ag contacts display poor results. The alumina layer tends to increase the V oc but this interpretation is taken with care. For both samples, EQE measurements are compared to absorption spectra computed with the transfer matrix method. While for the standard cells, simulating a 3450 nmthick CdTe fits nicely with experimental results, for the 800 nm-thick sample the best match between numerical simulation and EQE is obtained with a 300 nm-thick CdSe x Te 1-x layer, with x se =0.424 in addition to a 500 nm-thick CdTe. The excellent match between measurements and simulations suggests a high quantum yield of the active layer, and high material quality was confirmed by photoluminescence.

Summary

In this chapter, we evaluated the performance of CdSeTe/CdTe cells with alternative back contacts. We discussed on the high V oc deficit of polycrystalline materials that hinders their efficiencies. We summarized the factors impacting the V oc , and presented solutions developed in the community to overcome the 1 V barrier. To reach experimentally good results, the back contacts should be both passivating and selective.

In Section 5.2, we deposited two types of metallic back contacts with high reflective properties on standard CdSeTe/CdTe samples. This constraint meets the need for light-trapping to increase the optical path for thinner cells. We compared the results obtained with and without an Al 2 O 3 passivation layer. Promising results were obtained for cells with Au contacts that could be improved with light soaking, in comparison to rather poor performances with Ag contacts. We have shown that a very thin layer of alumina in between Au contact and CdTe, if it does not dramatically improve the performance, it enables a slighlty better V oc . Preliminary results with a p-type Cu 2 O layer have left interesting perspectives: with much thinner layer and a controlled annealing in air, good figure of merits are expected regardless of the metal contact deposited on Cu 2 O. We also discussed the low J sc attributed to a low collection efficiency of charges, which depends on the density of photogenerated charges.

In Section 5.3, the same back contacts are deposited on thinner CdSeTe/CdTe samples. Thinner cells show very high performance with Au contacts due to a high quality material confirmed by photoluminescence. Numerical fits of EQE measurements for both samples have allowed to build a robust optical model for CdSeTe/CdTe cells, and to simulate finely the absorption spectrum of such cells. Woodhouse demonstrated that with 1 µm and 18% efficiency, the requirement of Te material of 74 metric tonnes (MT) per GW could be lowered to 17 MT/GW [START_REF] Woodhouse | Perspectives on the pathways for cadmium telluride photovoltaic module manufacturers to address expected increases in the price for tellurium[END_REF], while Zweibel shows for 200 nm-thick device and an efficiency of 15% this number could reach 4.4 MT/GW [START_REF] Zweibel | The impact of tellurium supply on Cadmium telluride photovoltaics[END_REF]]. In an era where sustainability should be emphasised and widely recognised in technology application, these predictions fully support the development of ultrathin CdTe devices. Ultrathin solar cells with thicknesses reduced by an order of magnitude or more have the unique potential to efficiently convert solar energy into electricity while saving scarce material or element, and exhibit remarkable stability to degradation [START_REF] Rimmaudo | Superior stability of ultra thin CdTe solar cells with simple Cu/Au back contact[END_REF]. For CdTe cells, such a reduction would result in thicknesses in the 200 nm range. However, as with ultrathin GaAs cells, the efficient absorption of the solar spectrum is challenged by incomplete photon absorption due to the very thin absorber. In this chapter, we study numerically ultrathin CdTe solar cells with nanostructured back contacts, in order to find a realistic route towards an efficient light-trapping. After a rapid overview of thin and ultrathin CdTe solar cells, we briefly describe the detrimental effect of the CdTe thickness reduction for the J sc in Section 6.2. Then, a light-trapping strategy with periodic patterning is implemented with CdTe. We propose several rear side architectures based on multi-resonant absorption, in accordance with the requirements for the process of CdTe back contacts. This is followed by a further analysis of the resonance mechanisms involved (Section 6.3). In Section 6.4, the fabrication steps of ultrathin CdTe cells with nanostructured back contacts is proposed. Finally, due to the already low price of CdTe cells and modules, the economic viability of our approach is discussed. Remarkably, all the experimental results lie below the single-pass absorption, as observed for CIGS solar cells [START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. This can be partly justified by the predominance of the CdS layer in theses results (few results with MZO) that generated parasitic absorption at short wavelength. [START_REF] Gupta | High efficiency ultra-thin sputtered CdTe solar cells[END_REF]], b. and h. Bothwell et al. [Bothwell et al. 2020], c. and d. Paudel et al. [Paudel et al. 2014], e. [START_REF] Salavei | Influence of CdTe thickness on structural and electrical properties of CdTe/CdS solar cells[END_REF], f. Amin et al. [START_REF] Amin | Highly efficient 1 µm thick CdTe solar cells with textured TCOs[END_REF], g. and i. [START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF], j. [START_REF] Paudel | Ultrathin CdS/CdTe solar cells by sputtering[END_REF], k. [START_REF] Luk | Surface plasmon polariton enhanced ultrathin nano-structured CdTe solar cell[END_REF].

Overview of thin and ultrathin CdTe solar cells

Only few devices investigated light-trapping strategy in CdTe solar cells. Rashwan et al. reported by nanostructuring both the top (glass) and the back side of the device a theoretical J th = 25.3 mA/cm 2 for a 250 nm-thick CdTe solar cell [START_REF] Rashwan | Optical modeling of periodic nanostructures in ultra-thin CdTe solar cells with an electron reflector layer[END_REF]; [START_REF] Rashwan | Optimization of the Light Trapping Nano Structures in CdTe Thin Film Solar Cells[END_REF]. On the other hand, Amin et al. reported with a 600 nm-thick device an experimental J sc =22.6 mA/cm 2 with the top TCO textured to increase the light confinement [START_REF] Amin | Highly efficient 1 µm thick CdTe solar cells with textured TCOs[END_REF]]. Finally, a strong absorption enhancement is reported for a 110 nm-thick CdTe nanocylinders coupled with Ag mirror ( J th = 23.4 mA/cm 2 ) via the excitation of surface plasmon polariton [START_REF] Luk | Surface plasmon polariton enhanced ultrathin nano-structured CdTe solar cell[END_REF]].

Paudel et al. have sputtered CdTe samples with different thicknesses from 2.1 µm to 250 nm. Interestingly, they reported no significant changes on back-contact recombination with the thickness, with unchanged collection of photogenerated carriers [START_REF] Paudel | Ultrathin CdS/CdTe solar cells by sputtering[END_REF]. This implies that with high-quality materials the implementation of a light trapping strategy would have an effective and essential role for higher photocurrent with thinner devices. In the following, the effect of the thickness reduction of CdTe solar cells will be further detailed, and the impact of the back side nanopatterning highlighted.

Light-trapping in CdTe solar cells

Effect of thickness reduction in planar CdTe solar cells

We first investigate the impact of thinning the absorber in standard planar CdTe solar cells. In Figure 6. 2(a), two examples from the literature illustrate the performance of CdTe solar cells after the absorber layer reduction [START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF]; [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]. In contrast to GaAs solar cells, when the CdTe layer is reduced it generates a drop in the V oc and the FF in addition to the J sc . Nevertheless, Bothwell et al. confirm the predominant role of the J sc in the decrease of the efficiency [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]]. Figure 6. 2(b) shows the theoretical J sc as a function of CdTe thickness with a gold back contact with the J sc measurements from Refs [START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF]; [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]. The parameters of the simulation are described in Section 5.3.3 of Chapter 5 and the corresponding stack is given in Figure 6. 3(b). This plot can be divided into two distinct trends. Above 1 µm, the short-circuit current is only slightly affected by the thickness variation. Below 1 µm, however, there is a consistent decrease in J sc , which is highly dependent on the CdTe thickness. The J sc values from the literature follow the same trend with a shift. [START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF]; [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]]. The J sc is written under each marker in mA/cm 2 . (b) Calculated J sc in CdTe solar cells with increasing absorber thickness. To gain more insight into the absorption behavior when the active layer is reduced, theoretical calculations for different CdTe layer thicknesses are plotted in Figure 6. 3(a). The detailed stack of the cell used for the simulation is shown on the left (see Figure 6. 3(b)). Considering a standard 3450 nm-thick CdTe layer, a thin 800 nm device decreases moderately the J th from 28.81 to 27.03 mA/cm 2 , but more dramatically down to 21.23 mA/cm 2 when considering an ultrathin 200 nm-thick CdTe layer. As with ultrathin GaAs cells, it is primarily the long wavelength range that is affected.

To get a clearer picture of light absorption in the whole cell, the absorption of each layer is plotted for a 200 nm-thick and a 3450 nm-thick CdTe device in Figure 6.4. The different EQE and optical simulation of CdTe and CdSeTe/CdTe solar cells in Chapter 5 have shown sub-bandgap absorption. Even though the bandgap of CdTe is closer to 830 nm, we consider the CdTe absorption until 900 nm. In the following, the current density of CdTe and CdSeTe will be calculated on the spectral interval 300 to 900 nm, assuming a perfect collection of photogenerated charges. The resulting J th is an upper limit value. Two types of typical back contacts are compared: Te/Ni and Au. Tellurium is a material used as a back buffer for record devices reported in the literature [Munshi et al. 2018c]. As for the gold contact, it has the highest work function and is therefore the most suitable contact from the electrical point of view (but the reflectivity is lower than Ag). Overall, the CdTe layer, regardless of its thickness, shows strong absorption in the 350-500 nm wavelength range. Only a small fraction is lost by parasitic absorption in the MZO layer (approximated with ZnO in our simulation), much less than with a CdS layer at the front [Munshi et al. 2018c]; [Munshi et al. 2017]; [START_REF] Wang | Exploring window buffer layer technology to enhance CdTe solar cell performance[END_REF]; [START_REF] Kephart | Optimization of the front contact to minimize short-circuit curent losses in CdTe thin-film solar cells[END_REF]]. In the case of a standard CdTe absorber thickness, very little difference is visible between the two types of back contacts. Nonetheless, with the gold contact and in the infrared region, we observe a slight increase in absorption. This is related to resonances generated by the increased reflection at the CdTe back interface as well as less parasitic absorption in the contact layer. A corresponding increase in the theoretical short-circuit current of ∆J th =0.8 mA/cm 2 between the two contacts is observed.

On the other hand, the absorption in ultrathin CdTe is strongly influenced by the type of back contact. The J th drops to 16.85 mA/cm 2 with Te/Ni contact, while reflective Au contact allows a better absorption in the 600-800 nm range and displays a higher J th of 21.23 mA/cm 2 . Nevertheless, it shows that a very thin CdTe layer requires an efficient light trapping strategy to better absorb the light coming from the red part of the spectrum [Goffard et al. 2017]; [START_REF] Gouillart | Development of ultrathin Cu (In,Ga) Se2 -based solar cells with reflective back contacts[END_REF]]; [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF].

While Au is a decent reflector, Ag has the highest reflectivity among metals in the visible-NIR spectral range and provides better results. Figure 6. 5(a) shows the improved absorption in very thin CdTe compared to the Au contact. J th are reported in Figure 6. 5(b), showing ∆J th =0.8 mA/cm 2 between the two metal contacts. Note that while the parasitic absorption in silver is always smaller than in gold, the change in contact has no visible effect on the J th for thicker films. We also note the large parasitic absorption coming from the FTO layer, where the dJ loss is about 3 mA/cm 2 overall (2.45 mA/cm 2 for thicker cells). This phenomenon will be further investigated in section 6. We have demonstrated the strong correlation between the decrease in J sc and the thickness reduction of CdTe solar cells. This confirms the need to develop a novel and efficient light-trapping strategy with a nanostructured grating and a highly reflective and flat mirror to maintain a proper J sc . In the following, we present three different backside architectures for a 200 nm-thick ultrathin CdTe solar cell.

As for ultrathin GaAs cells, the reduction of the CdTe absorber thickness critically affects the J sc , and typical back contacts are not suitable compared to Au. However, the J th drop between standard and ultrathin devices also observed with Au contacts proves the importance of increasing the optical path length.

Nanostructured back mirror

As discussed in Chapter 5, CdTe back contacts are a major challenge. In this respect, we try to keep the back contacts realistic in our simulations. We adapt the highly efficient method introduced with the 200 nm-thick GaAs solar cells, where periodic nanostructured layers have demonstrated their ability to increase the solar cell efficiency [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]; [START_REF] Chen | Ultrathin and nanowire-based GaAs solar cells[END_REF]]. We propose novel structures with transparent and nanostructured layers made of dielectric, Cu 2 O, Al 2 O 3 and p-doped amorphous silicon a-Si:H layers followed by a reflective silver layer. Al 2 O 3 and a-Si:H layers are used in the stacking of high-performance cells [Munshi et al. 2018a]; [START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF]; [START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]]. The CdTe active layer will be kept flat (no nanostructuring within the layer) to avoid damaging the absorber and potential electronic degradation such as increased surface recombination [START_REF] Misra | Cadmium Selective Etching in CdTe Solar Cells Produces Detrimental Narrow-Gap Te in Grain Boundaries[END_REF]]. We extensively develop the initial architecture optimization and results. We then compare it with three alternative back structures.

Rear side architecture with a transparent dielectric layer

For our first set of simulations, the structure is based on the same architecture originally developed with ultrathin GaAs: Ag nanopillars embedded in an optically transparent layer with a given refractive index n. As described in Chapter 2 for ultrathin GaAs cells, several parameters are varied in the simulation to optimize the back structure. These variables are highlighted in red in the simplified scheme of an ultrathin CdTe solar cell with a nanostructured back mirror used for the calculation (see Figure 6.6(a)). Again, J th is used as a figure of merit in the optical simulation to characterize the performance of the cells. The parameters are optimized by maximizing J th in the absorber layer. The refractive index of each layer is plotted in The best parameters found for this back side geometry of an ultrathin CdTe solar cell are summarized in Table 6.1. Interestingly, this particular structure is hardly affected by the refractive index value of the dielectric layer in which the Ag nanopillars are embedded. Simulations with a large set of parameters and where the refractive index is swept highlight this conclusion (see Appendix A.4.1,Figure A.12(b)). The best currents are observed in a range of small periods of the nanostructures and for width around 250 nm. The maximum current is obtained at a refractive index of 1.9, close to the value of TiO 2 . To illustrate our optimized light trapping design, we compare CdTe solar cells with a nanostructured mirror to the same structure with a flat Au mirror. The corresponding absorption spectra are plotted in Figure 6 Since the standard CdTe film thickness is reduced by more than ten in the simulation, the light-trapping effect is evidenced in the long wavelength range (λ > 600 nm). This results in a gain of 3.5 mA/cm 2 for J th compared to the Au planar mirror alone. To explain these improvements, we first have the effect of the highly reflective silver layer. Then, a large number of overlapping resonances (peaks in the absorption spectra) contribute to further absorption enhancement. The black curve in Figure 6.6(b) represents the total absorption for the nanostructured cell and reveals that there is still room for improvement. In addition, the poor electrical contact between Ag and CdTe has been shown in the previous chapter. Therefore, alternative back side architectures need to integrate commonly used back contacts or passivation layers for CdTe devices. Examples are investigated in the following.

Alternative rear structures

For a 200 nm-thick CdTe cell, three different designs are proposed to address these challenges and are illustrated in Figure 6.7.

-Structure A: A direct application of the first example, as the mirror is nanostructured rather than planar. The backside textured layer consists of a copper oxide layer with Ag nanopillars. The Cu 2 O is considered to be optically transparent with a refractive index of 2.262. The height, width and period of the nanostructures are optimized.

-Structure B: 20 nm of alumina is added between the absorber layer and the back contact. Ag nanopillars are embedded in a dielectric layer, whose refractive index is optimized to 1.3. This value can be achieved with a porous SiO 2 . Again, the height, width and period are optimized to maximize J th .

-Structure C: The silver metallic mirror is no longer nanostructured and in direct contact with the active layer. a-Si:H nanopillars are embedded in a dielectric layer whose refractive index has been optimized to 1.9. A planar Ag mirror is then added. The h and h n parameters are optimized. This structure is a direct extension of our best ultrathin GaAs model to CdTe, where this back architecture has proven its efficiency.

These three structures are attractive because they combine the characteristics required for good electrical contacts for CdTe cells. Moreover, the know-how developed for GaAs in the C2N, makes their fabrication entirely achievable. For each structure, we performed numerous simulations with various geometrical parameters to maximize the short-circuit current. In the following, we only show the best results for each structure. All the optimized parameter values are summarized in Table 6.2. The gain in J th compared to a device with a planar Au mirror, J gain , is also mentioned, and the parasitic absorption occurring in the back contact is referred to as dJ. Although the contact between the p-type TCO Cu 2 O and CdTe needs refinement (see Chapter 5), results in the literature make it an interesting choice for structure A [START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF]]. The use of the Cu 2 O optical index offers a substantial increase in J th compared to the Au contact alone. As for the initial structure with TiO 2 , the nanostructuration leads to a multi-resonant absorption spectrum with several sharp resonance peaks. Nevertheless, the J gain is less important than for the first structure (24.23 versus 24.7 mA/cm 2 ). A possible explanation could be related to the refractive index value of Cu 2 O, which is too close to that of CdTe. Moreover, the absorption spectrum in Ag shows large optical losses. It could be lowered by adding an interstitial dielectric layer before it.

In structure B, we add an interstitial Al 2 O 3 layer, which is typically used as a passivation layer. The silver mirror remains structured and not planar. A higher J th is obtained in the CdTe layer and the parasitic loss in the Ag layer is reduced. However, the J gain , is lower than expected. On the one hand, less resonances are visible for the structure B compared to A in the range of 600-800 nm. On the other hand, although the parasitic absorption in the silver mirror is reduced beyond 800 nm, the total absorption is also reduced.

The use of a-Si:H as a contact layer for CdTe cells has led to promising results in the recent literature [Munshi et al. 2018a], and we find that structure C has the best J gain . In this configuration, the metal mirror is no longer in contact with the active layer. The overlapping resonances provide a constant absorption close to 0.85 in the 700-800 nm range. With a flat mirror on the back, parasitic absorption in Ag is strongly reduced in the region of interest. However, as for GaAs cells, detrimental parasitic absorption in the amorphous silicon layer reduces the collected current in the absorber layer between 500 nm and 600 nm. Overall, improving the back side design increases the gain in J th and decreases the parasitic absorption losses.

These first results demonstrate the value of simulation for the optimization of ultrathin CdTe cells. It paves the way for the development of new and efficient devices.

Alternative architectures for reducing parasitic losses

In this part, we take a closer look at the best structure obtained in terms of current collection in the active layer. We address the loss induced by the FTO layer by simulation, since minimizing its loss would lead to an almost perfect absorption in the 650-800 nm region (green curve in the graph 6.7(c)). Figures 6.5(b) and 6.7(c) highlight the significant absorption loss in the SnO 2 :F (FTO) layer. For the structure C, the collected current in the active layer below the bandgap region is J th =25 mA/cm 2 , while the loss in the FTO is equal to the non negligible value of about 3 mA/cm 2 . Figure 6.8(a) shows the calculated absorption and transmission of the SiO 2 +FTO stack corrected by the light transmission through bare glass. No precise refractive indices are available for the commercialized glass TEC12D coated with SiO 2 and FTO used for CdTe cells. It could justify that the optical properties we use lead to an overestimation of the losses in the FTO compared to the reality. Indeed, other groups have reported much lower values for the losses in the FTO, in the order of 0.6 mA/cm 2 [START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]; [START_REF] Bothwell | Characterization of thin CdTe solar cells with a CdSeTe front layer[END_REF]]. Lower extinction coefficient over the visible spectral range (400-800 nm) have also been reported below than 0.01 [Compaan et al. n.d.] 300 400 500 600 700 800 900 1000 0.000 If we consider the structure C without any loss in the FTO layer, the collected current in the 200 nm-thick CdTe is optically equivalent to a J th of 28 mA/cm 2 (we assume that the part of the absorption lost in the FTO is now completely collected by the CdTe layer). Kosyachenko et al. estimate that the FTO layer at a thickness of 500 nm reduces the short-circuit current by 4-5% (much less than ITO, which can reduce the current by up to 25% for the same thickness) [START_REF] Kosyachenko | Quantitative assessment of optical losses in thin-film CdS/CdTe solar cells[END_REF]]. In our case for a J th =28 mA/cm 2 the parasitic absorption in the FTO layer represents more than 10%. To reduce the losses in the FTO, we divide in our simulation model the extinction coefficient of the FTO by 5 (shown in Figure 6.8(b)).

The original structure C is slightly modified as we adjust the extinction coefficient of the FTO layer. The parameters are re-optimized and we obtain the new calculated absorption spectrum in Figure 6.9(a). In the active layer we find J th = 26.41 mA/cm 2 and the loss in FTO dJ= 0.64 mA/cm 2 . The J th value is one step closer to the J sc of record standard CdTe cells, evaluated at 30.25 mA/cm 2 and corresponds to an increase of 25% over solar cells with flat Au back contact. Note that if we consider a 200 nm-thick CdSeTe absorber and the same structure, the collected current reaches 28.26 mA/cm 2 (see Figure 6.9(b)).

Conclusion of the section

Standard CdTe cells with thickness of several micrometers efficiently absorb a large portion of the visible spectrum, but their J sc is strongly reduced when the active layer is reduced by a factor of 10. This absorption loss justifies the implementation of a lighttrapping strategy to increase the optical path length and make the ultrathin CdTe cells attractive. In this section, we have proposed four different back mirror designs for a 200 nm-thick CdTe cell. The variety of the solutions is justified by the demanding criteria of a good contact with CdTe films. Structure C offers the best performance, J th = 25 mA/cm 2 , with a strong overlapping of the resonances leading to a constant absorption of 85% over the 700-800 nm range. Reducing the parasitic absorption in the FTO layer and considering a 200 nm-thick CdSeTe cell (with a lower bandgap), an efficient light absorption is obtained, resulting in J th = 28.3 mA/cm 2 .

Analysis of the multi-resonant light-trapping in CdSeTe solar cells

In this section, we attempt to identify the resonances observed in optimized ultrathin CdTe solar cells. The theoretical calculation of the resonances are carried out thanks to H.L. Chen who developed and computed a code to find both Fabry-Perot and guided-mode resonances. The detailed equations involved in the identification of the peaks are extensively explained in his manuscript [START_REF] Chen | Ultrathin and nanowire-based GaAs solar cells[END_REF]]. We conduct this study on the improved structure C, with less parasitic absorption in the FTO (see Figure 6.9). In Figure 6 To gain more insights in the mechanisms of optical resonances, the specular reflectance R is calculated. The total absorption variation with geometrical parameters is highlighted by plotting 1-R as a function of the wavelength and for CdTe thicknesses from 30 nm to 600 nm in Figure 6.11(a), and 1-R as a function of the wavelength and for grating periods from 100 nm to 600 nm in Figure 6.11(b). In this case, the width-to-period ratio (50%) is kept constant. Figure 6.11(a) confirms the high absorption quality of CdTe cells, showing that for absorber thicknesses between 150 nm and 300 nm the specular reflectance is generally high over the visible range. However, for thinner CdTe layers, it drops significantly.

With Figure 6.11(b), we immediately observe the strong dependence of the absorption with the grating period. The best absorptions are obtained with a period centered around 400 nm. Outside this range, the absorption is reduced because of a lack of resonance. In the short wavelength range, up to λ= 600 nm, 1-R is globally constant with the grating period. Moreover, the absorption spectrum is similar to the one with an Au back contact only, the peaks are more or less at the same wavelength. In Chapter 2, two types of resonances were presented: Fabry-Perot and guided modes resonances. With their typical field distribution, we can identify these resonances inside the cells. For our simulations, we consider a multi-layer cavity with an averaged refractive index (equation 2.6), the incident medium is glass where n=1.5 and the substrate is silver.

λ = 805 nm k in H in E in k in H in E in (a) (b) (c) (d) (e) (f) (g)
Figure 6.12: Cross-sections of the electric-field intensity (a-f), or magnetic-field intensity (g) calculated by RCWA in a 200 nm-thick CdTe device (optimized structure C) for some resonance wavelengths.

We associate the short wavelength peaks to FP resonances. For the solar cell with Au contact only, and based on the two equations 2.5 and 2.6, we find the resonant wavelengths: 396 nm (FP 9 ), 470 nm (FP 7 ) and 527 nm (FP 6 ) close to the absorption peaks in Figure 6.10 (yellow line). For the cell with the nanostructured a-Si:H/TiO 2 grating, FP resonances are slightly shifted and we obtain 460 nm (FP 10 ) and 514 nm (FP 9 ). We show in Figure 6.12(a,b) the cross-section maps of electromagnetic field intensities for these resonant wavelengths. They are calculated using RCWA. For resonances A and B, the electric field varies mainly along the vertical direction, as expected for FP resonances. At longer wavelengths and for a period comprised between 300 nm and 500 nm, (1-R) depends on the grating parameter. Resonances are attributed to light-coupling to guided-modes and they show a lateral structuring of the light (Figure 6.12(c-g)). It is possible to link the shape of the field and the number of intensity maxima in x and z as a function of the mode type and the coupling order [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]]; [START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF]]. These resonances are identified as TE q (m 1 ,m 2 ) and TM q (m 1 ,m 2 ), where q is the order of vertical oscillations and m 1 ,m 2 the orders of diffraction. For a given guided mode, we calculate the resonance wavelengths for each grating period. The results are plotted in Figure 6.11(b) and are in excellent agreement with the peaks computed by RCWAs.

Conclusion of the section

We have tried to identify some resonances in a 200 nm-thick CdTe cell, with an a-Si/TiO 2 grating followed by a planar Ag mirror. Based on the theoretical model developed in H.L. Chen's thesis we associated Fabry-Perot and guided modes to most peaks visible in the CdTe absorption spectrum.

Experimental perspectives for ultrathin CdTe cells 6.4.1 Description of fabrication steps

We have numerically demonstrated that a 200 nm thick device could have a competitive shortcircuit current. In this section, we present the fabrication steps for three structures, necessary to complete these nanostructured ultrathin CdTe devices. The back side architecture is inspired from our calculations, respectively structures C and B. We propose three lines of thought to achieve efficient cells: Model 1. The Ag mirror is necessary to enhance light trapping in ultrathin CdTe solar cell due to its high reflectivity, however we have demonstrated its poor performances as an electrical contact for CdTe solar cells in Chapter 5. Here, to overcome the tradeoff between high reflectivity and good ohmic contact, the process involves a passivation layer and localized gold contacts followed by an Ag metallic mirror. No other light-trapping strategy than the planar Ag mirror at the rear side is involved. Note that this was the only model subjected to experimental trials at the C2N.

Model 2. This model is inspired by the theoretical structure C, with an Al 2 O 3 passivation layer followed by a nanostructured dielectric layer (TiO 2 ) with embedded highly doped a-Si:H nanopillars. A planar Ag mirror is added at the end.

Model 3. This model is inspired by the theoretical structure B, with a passivation layer and a nanostructured SiO 2 /Ag mirror. This model should be less costly.

To begin with, we describe the steps to obtain a CdTe cell with localized back contacts embedded in a passivation layer (Model 1.). A planar silver mirror is added and assures a good reflectivity.

Step 1: Spin-coated PMMA is patterned by e-beam lithography. The size of the squared Au contacts are in the order of hundred of nanometers (d), while the pitch is ten times larger (about 1 µm). The pattern remains to be optimized. Step 4: Each layer is nanostructured. E-beam lithography will be preferred for proof-of-concept experiments with a-Si:H layer, while nanoimpression (low cost process) will be easily implemented for the dielectric layer [START_REF] Chen | Ultrathin and nanowire-based GaAs solar cells[END_REF]].

CdTe CdTe

Rear side nanostructuration

Step 5: The a-Si:H layer is planarized with TiO 2

Planarization with TiO2

CdTe TiO 2

Step 6: For both models, contacts are revealed using optical lithography/EBL and either wet chemical or reactive ion etching.

CdTe TiO2 SiO2

Contact revealing

CdTe

Step 7: An Ag mirror is evaporated on the contacts, completing the cells.

CdTe CdTe Ag Ag

Ag mirror deposition

Cost assessment

While no detailed economical study is presented here, we justify the value of reducing the use of Te in the CdTe photovoltaic fields. Although the need of reducing the CdTe thickness is unquestionable, the technologies and material needed to compensate the absorption losses might be costly. As CdTe modules efficiency is slightly lower than silicon ones, it is their very low levelized cost of energy that compensates this fact and turns them into a competitive technology. Moreover, the fast all in line fabrication of CdTe modules which is a real asset will be challenged if the absorber thickness is reduced. Hence, having the same efficiency but for more expensive production thwarts the thin-film pv modules objectives. The gain through reducing the semiconductor thickness for CdTe is not as straightforward as for GaAs which is an expensive material. The evaluation of the CdTe module costs aggregated by major process is shown in Figure 6.13. This graph confirms the low impact of CdTe material which represents only 15% of the price of all materials combined to build a module [START_REF] Smith | Photovoltaic ( PV ) Module Technologies : 2020 Benchmark Costs and Technology Evolution Framework Results[END_REF]]. Reducing the absorber thickness would result in a material cost of 0.0026 $/W instead of 0.026 $/W. For ultrathin CdTe cells, the process suggested in the last section would include several costly fabrication steps.

The most two expensive steps would be the localized gold contact evaporation with lithography and the silver mirror deposition. Based on report carried out for GaAs devices, those steps would add more than 9$/W to the original price [START_REF] Horowitz | A Techno-Economic Analysis and Cost Reduction Roadmap for III-V Solar Cells[END_REF]]. Nevertheless, several low cost alternatives could be developed. Back localized contacts could be a combinaison of Te/Ni (5000 times cheaper than gold) while silver mirror could be deposited using electroplating instead of evaporation. The latter would results in an additional cost of 0.02 $/W for the material part In the next years, CdTe manufacturing is set to continue to grow significantly. Note that the direct impact of the Te scarcity on future development is not fully understood and it is still hard to obtain a clear picture of the future demand. Unless proved otherwise, resources in their present state are unlikely to constrain CdTe technology to supply a large quantity of utility-scale PV systems. Nonetheless, Candelise et al. show in their study that Te scarcity could lead to a dramatic material price increase, and keeping a low production cost would be jeopardized. Scenarios were developed with material usage and module efficiency. The scarcity of key components takes on its full significance, and Te represents a significant part of the total cost (Figure 6.14) [START_REF] Candelise | Implications for CdTe and CIGS technologies production costs of indium and tellurium scarcity[END_REF] Although additional research is required to fully develop the ultrathin CdTe devices fabrication, we have shown as a first step, that with adapted processes no dramatic cost is added. Moreover, progress in the reduction of materials consumption is key for a sustainable development.

Conclusion of the section

In this section, perspectives for the fabrication of ultrathin CdTe devices are presented. The combination of Au localized back contacts spaced a few µm apart and a reflective Ag mirror has never been tested to our knowledge. The passivation layer and the Au contact would ease the carrier collection while the mirror could enhance the optical path. This technique would allow the development of the light-trapping of models 2 and 3. This would open the way to excellent ultrathin cells with similar performances as the standard ones. Such designs implemented with high quality CdSeTe material would lead for the best ultrathin cells, to J sc =28 mA/cm 2 . Considering the high V oc and FF values of the record device, 887 mV and 78.5% respectively, the efficiency of a 200 nm-thick CdSeTe solar cell could reach 19.5%, without significantly increasing the cost production.

Summary

Due to its high absorption coefficient, very few studies have been conducted on the impact of thickness reduction of CdTe absorber.

In Section 6.2, we have shown that below 1 µm, when the CdTe absorber thickness is reduced, the J th is directly impacted and is decreased drastically: 6 mA/cm 2 is lost between 1 µm and 200 nm. Based on the optical light-trapping strategy implemented with ultrathin GaAs, we have simulated the absorption in a 200 nm-thick CdTe cell with a periodic patterning at the backside. We presented different back mirror/contact architectures, with material properties that make good contacts experimentally. Thanks to a broadband multi-resonant absorption we were able to simulate a J th = 26.41 mA/cm 2 for a 200 nm-thick CdTe cell and J th = 28.26 mA/cm 2 for a 200 nm-thick CdSeTe device. The in-depth analysis of resonances of the best performing CdTe cell is presented in Section 6.3. Finally, Section 6.4 describes encouraging perspectives for the fabrication of ultrathin CdTe cells, with path to a 19.5%-efficient CdSeTe cells. [START_REF] Gupta | High efficiency ultra-thin sputtered CdTe solar cells[END_REF]], b. and h. Bothwell et al. [Bothwell et al. 2020], c. and d. Paudel et al. [Paudel et al. 2014], e. [START_REF] Salavei | Influence of CdTe thickness on structural and electrical properties of CdTe/CdS solar cells[END_REF], f. Amin et al. [START_REF] Amin | Highly efficient 1 µm thick CdTe solar cells with textured TCOs[END_REF], g. and i. [START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF], j. [START_REF] Paudel | Ultrathin CdS/CdTe solar cells by sputtering[END_REF], k. [START_REF] Luk | Surface plasmon polariton enhanced ultrathin nano-structured CdTe solar cell[END_REF]. The blue and the purple stars are the J sc obtained with a CdTe absorber and a CdSeTe absorber, respectively.

Conclusion and perspectives

Ultrathin solar cells, with thicknesses of about 10 times less than conventional cells, represent a promising option to relieve the limitations on scarce material, save fabrication time and cost and ease the carrier collection in defective absorber. They open the way towards more efficient and low-cost photovoltaic devices that may result in competitive advantages in the PV industry. To achieve such a breakthrough, efficient light-trapping strategies within the absorber are required. Importantly, they must meet many constraints related to the process of actual devices as well as the efficient transport and collection of charge carriers. In this thesis, I reported several strategies for the design and the fabrication of ultrathin solar cells made of GaAs and CdTe.

Ultrathin GaAs solar cells

In 2019, our team demonstrated a record device with 19.9% efficiency using a 205 nm thick GaAs absorber (Section 2.1.2). Broadband multi-resonant absorption was achieved with a nanostructured TiO 2 /Ag back mirror. In order to improve the performance of this first generation of ultrathin GaAs solar cells, a comprehensive analysis of the absorption losses was performed in Chapter 2. From these optical simulations, two possible ways were explored to get one step closer to the Lambertian scattering model in terms of short-circuit current.

The nanostructured back mirror was identified as one of the main sources of parasitic absorption and to address this issue we modified the backside of the original ultrathin stack in the first approach. I introduced a novel back architecture, using a 2-dimensional nanopatterned diffraction grating made of a high refractive index semiconductor a-Si:H with a relatively high bandgap and a transparent mesoporous SiO 2 dielectric with refractive index n= 1.2. The use of the SiO 2 as spacer in between the grating and the silver mirror allows the reduction of the losses in the flat silver mirror. A very high broadband multi-resonant absorption was achieved in the spectral range 700 nm to 920 nm leading to a theoretical short-circuit current density J th of 28.2 mA/cm 2 . Despite some parasitic absorption in the a-Si:H, the net J sc increase compared to the original structure corresponds to 2.63 mA/cm 2 .

In the second approach, to further reduce the parasitic absorption in the back surface field layer and in the mirror, as well as to increase the V oc , I investigated the use of a heterostructure composed of 200 nm n-GaAs emitter / 50 nm p-AlGaAs base and using a similar light-trapping scheme this time based on an AlGaAs 2-dimensional nanopatterned diffraction grating (replacing the a-Si:H), a transparent mesoporous SiO 2 spacer and a flat mirror. Assuming a perfect collection of the photo-generated carriers also inside the base (250 nm-thick device) this new architecture can reach a J th of about 28.9 mA/cm 2 .

Subsequently, I carried out in Chapter 3 the fabrication of ultrathin GaAs solar cells with the novel back architecture (with an a-Si:H diffraction grating), using the first-generation III-V epitaxial layers grown at the Fraunhofer ISE. I reported the successful fabrication of solar cells with several innovative steps such as the nanostructuration of an a-Si:H layer by e-beam lithography and dry etching and the planarization with a very low index material (n= 1.2) made of mesoporous SiO 2 . I demonstrated functional devices with state-of-the-art V oc and FF, and EQE exhibiting a multi-resonant behavior. To validate this approach, an additional optimization of the fabrication process would be required, to get close to the expected efficiency of 24%.

Conclusion and perspectives

CdSeTe/CdTe polycrystalline devices

In the second part of this thesis work, I investigated CdSeTe/CdTe solar cells and focused on the different factors that hinder the successful development of ultrathin CdTe devices.

CdTe-based solar cells have significantly benefited from the introduction of selenium. This results in an increase of the photogenerated current due to the bandgap lowering, in an increase of the carrier lifetime and a reduction of the V oc deficit. In Chapter 4, I conducted a systematic study on five CdSeTe/CdTe samples with different concentrations of selenium by cathodoluminescence (CL). High resolution CL maps at both room and low temperature were acquired in bevels, allowing an in-depth analysis of the material properties. From the analysis of the hyperspectral maps I have obtained several important results. First, the bandgap of CdSe x Te 1-x semiconductor alloys was determined for any temperature and Se concentration between 0 and 0.4. Then, a 10-fold enhancement of the luminescence between CdTe and CdSe 0.4 Te 0.6 was evidenced, and the higher radiative efficiency is attributed to the passivation role of Se. Regardless of the sample, the grain size, or the position on the bevel, the CL intensity only depends on the amount of Se in the device. Finally, maps of Se concentration were plotted allowing an estimation of the Se content with an accuracy of 0.5%. Most of the Se diffusion in the CdTe layer occurs during the CdCl 2 post-deposition treatment and is mainly driven by the grain boundaries. A quantification of the diffusion coefficient of Se atoms in both grain interior and grain boundaries was obtained.

In Chapter 5, I explored alternative back contacts for CdSeTe/CdTe solar cells grown at the National Renewable Energy Laboratory taking into account the necessity for them to be ohmic, passivating and optically reflective. The performance of both standard (3-4 µm) and thin (800 nm) devices has been experimentally investigated. Promising results were obtained for cells with Au contacts with efficiencies around 12.5%, in comparison to rather poor performances with Ag contacts due to low V oc and FF. A very thin alumina passivation layer with Au contact enables a slightly higher V oc , yet is detrimental with Ag contact. I demonstrated preliminary results using one of the few p-type TCO available: Cu 2 O deposited by Spatial Atomic Layer Deposition. Although poor performance was shown, improved figure of merits are expected with thinner layer and low temperature annealing. Except for the Cu 2 O, the same back contacts have been deposited on thinner (800 nm) CdSeTe/CdTe samples. High quality material confirmed by photoluminescence is responsible for the best performance obtained with Au contacts.

Prior to the analysis of ultrathin CdTe-based solar cells, numerical fits of EQE measurements led to a robust optical model for CdSeTe/CdTe cells.

Ultrathin CdTe-based solar cells

Light trapping strategies similar to GaAs have been proposed and investigated for ultrathin (200 m) CdTe solar cells. Care was taken to present different back architectures with solutions that would comply with the demanding criteria of a good back contact: ohmic contact, high optical reflectivity and compatibility with the solar cell fabrication process. I presented an optimized structure with an a-Si:H/TiO 2 diffraction grating associated to a flat mirror that results in a theoretical short-circuit current of 25.7 mA/cm 2 for a CdTe device and of 28.1 mA/cm 2 for a CdSeTe device. I briefly evaluated the mechanisms of optical resonances by varying the parameters of the grating structure and the thickness of the absorber. Only a few studies report on ultrathin CdTe solar cells and our work not only overcomes experimental devices that lie below the single-pass absorption, but also exceeds the double-pass absorption limit.

Perspectives

Ultrathin solar cells embody a new class of devices with thicknesses one order of magnitude smaller than their standard counterpart. In this work, I have presented advances toward ultra-thin and highly-efficient GaAs and CdTe solar cells.

I have demonstrated the feasibility of an ultrathin GaAs solar cell with a nanopatterned diffraction grating made of a-Si:H/SiO 2 . Fine-tuning of the fabrication process, in particular the spacer thickness, should allow to overcome the J sc of the original ultrathin GaAs solar cell in the short term. Combining this change with our effort to increase the cell dimension would result in a higher V oc , and a higher efficiency overall. The nanopatterning of the amorphous silicon layer by nanoimprint lithography instead of e-beam lithography would be beneficial to fabricate large-area cells at lower cost.

The new stack developed in Chapter 2 based on a heterojunction has been grown at the Fraunhofer ISE and received during the writing of this manuscript. The processing of these new ultrathin samples offer new perspectives with higher V oc enabled by the heterostructure and the enhanced photon recycling.

The J sc limit of 30 mA/cm 2 has yet to be exceeded and in the long term the development of more complex patterns could be implemented at the backside of the ultrathin device. For instance, hexagonal arrangements or three-dimensional nanostructures could help to further increase the number of resonances. This strategy could lead to higher efficiencies or even thinner (less than 100 nm) and performant cells allowing the introduction of advanced concepts involving hot carriers. Ultrathin GaAs solar cells may be of interest for space applications due to their higher intrinsic radiation tolerance, but further studies are needed to evaluate the performance of our samples after irradiation (EOL).

Prior to the introduction of light trapping strategies to ultrathin CdTe-based devices, I characterized high-quality CdSeTe material by cathodoluminescence.

A quantitative study of the selenium passivation role has been demonstrated but we could go one step further and determine the absolute value of the radiative efficiency in the layer thanks to an absolute calibration of the setup. To this end, a calibration diode in GaAs was fabricated. The photon flux emitted by the diode was first assessed using a calibrated hyperspectral imager and then compared to the electroluminescence signal collected in the CL setup. A first conversion ratio of 1.5 × 10 5 photons.cts -1 was estimated. Repeating similar steps for other calibration diodes emitting at different wavelengths would allow to build a robust model for the absolute calibration of the CL. The processing of preliminary measurements in the infra-red region and at low temperature performed to evaluate the effect of different structures, back contacts or passivation layers on CdSeTe-based devices are still in progress.

The implementation of passivating and selective back contacts remains critical for the CdTe and would benefit from systematic measurement assessment: the voltage loss analysis presented in the beginning of Chapter 5 could be performed on the standard and thin CdSeTe/CdTe solar cells fabricated by the NREL (Chapter 5) to better understand the reason behind their efficiency limitation.

The optical simulations predict a bright path for ultrathin CdTe-based solar cells. Devices with 10 times less material displaying an efficiency of almost 20% encourage us to pursue this approach.

where R sqf is the sheet resistance of the metal grid fingers and the n-GaAs contact layer highly doped. Considering the equivalent resistance of both material, we restrict the sheet resistance only to the metal (gold), and R sqf =ρ f /t, we take ρ f =2.35 × 10 -6 .cm 2 . The front contacts have a gold thickness of 220 nm. The last one is due to the contact resistance and calculated by the simplified equation [START_REF] Liu | Optimization of grid design for solar cells[END_REF]; [START_REF] Schroder | Solar Cell Contact Resistance-A Review[END_REF]: .5) where the resistivity between the metallic contact and the emitting layer is ρ contact and taken equal to 2.16 × 10 -4 .cm 2 . This value was experimentally determined by transmission line measurement (TLM) method.

P contact = ρ contact J mp V mp S w f (A

A.1.2 Thick front contact fabrication

The thickness of the front metal contacts is targeted between 1 µm and 2 µm. To ensure a good lift-off, the ratio between the resist and the metal is estimated at 3 for 1, which requires a process with thick resist of about 7 µm. Two different fabrications have been used to address this challenge. On the one hand, a low temperature process using the nLof-2070 resist has been developed and optimized by L. Nguyen and is extensively described in her manuscript. On the other hand, I developed a bilayer process using LOR30B and AZ5214 resists. In this section, the latter will be emphasized. Figure A.2(f) illustrates the design of the front electrode lithography mask: four 0.5x0.5 cm 2 patterns which are composed of 5 µm wide fingers connected to 30 µm wide busbars. In practice, a width of 5 µm almost reaches the minimum resolution that can ensure the photolithography technique for lift-off and wet etching. This process has first been tested on silicon wafer, then on a transferred GaAs sample on glass, and finally optimized for ultrathin GaAs layer transferred on glass. It first requires a 5 min dehydratation at 180°C on a hotplate, to prevent cracks in the resin during the developing. The resin LOR30B is enducted by spin-coating (60s, speed:1800 rpm, acceleration:500 rpm/s), the thickness reached is about 7 µm. The sample is annealed for 5 min at 180°C (the timing takes into account the glass thickness), and after this step some bead edges might appear. Then the second resist AZ5214 is spin-coated: (30s, speed:4000 rpm, acceleration:2000 rpm/s), resulting in 1.4 µm thickness. One again the resist is annealed at 125°C on a hotplate for 2 min. To remove potential bead edges and ensure a good contact between the sample and the mask during the exposure step, the whole sample is protected by a physical mask except the borders. It is UV-insolated for 2 min, and then developed in AZ400K:H 2 O (1:3) for 3 min. Then the sample is exposed to UV with the mask (hard contact) for 5s. The resist is reticulated for 2 min at 125°C, and then the whole sample undergoes a flood exposure of 30s. The patterns are developed in the developer MIF826 for about 1 min. A.2 Chapter 4

A.2.1 Generalized Planck's law fit

In Chapter 4, we fit the CL spectra with the generalized Planck law. To attest to the correct assumption made on the absorption coefficient α c =30000 cm -1 , we plot in Figure A.4(a) the absorption coefficients calculated from the refractive indices found in the literature, [START_REF] Adachi | Optical properties of CdTe: Experiment and modeling[END_REF]]; [START_REF] Treharne | Optical design and fabrication of fully sputtered CdTe/CdS solar cells[END_REF]] and provided by the NREL [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]]. The values taken by alpha at an energy of 1.7 eV (200 meV above the bandgap) are specified in the legend. In Figure A.4(b), the absorption coefficient for several concentration in selenium is plotted (from NREL), and shows very little variation between the samples. The approximation made on α c is justified. [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF] and found in the literature [START_REF] Adachi | Optical properties of CdTe: Experiment and modeling[END_REF]; [START_REF] Treharne | Optical design and fabrication of fully sputtered CdTe/CdS solar cells[END_REF]]. (b) Absorption coefficient of CdSe x Te 1-x calculated with refractive indices provided by the NREL depending on the Se concentration [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]. The graph on the right is a zoomed plot corresponding to the red rectangle, the absorption coefficient α is plotted between 1 eV and 2.5 eV.

Figure A.5 compares the absorption coefficients provided by the NREL (solid lines) and the ones modelled by a parabolic band approximation (dashed lines) (equations 4.2 and 4.3), with α c =30000 cm -1 and E 0 = 1.7 eV for several Se concentrations. The values taken by E g and γ depend on the selenium concentration, and we use the values extracted after the generalized Planck law and the absorption model fit, which are summarized in Table A.1. Note that when we compare the two absorption coefficients curve for different Se concentration, the Se content is not exactly the same. For example, in Figure A.5(a), the absorption coefficient provided by the NREL is for a selenium concentration of x=0.083 and the absorption coefficient calculated is for a nominal concentration of x =0.1.

Se content x se = 0 x se = 0.1 x se = 0.2 x se = 0.3 x se = 0.4 E g (eV)

1.501 [START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]] (solid lines) compared to absorption coefficient calculated with equations 4.2 and 4.3 (dashed line), the inset is in logscale for (a) x=0. 1, (b) x=0.2, (c) x=0.3, and (d) x=0.4. CL spectra plotted as solid lines in Figure A.6, are fitted with the generalized Planck law (see Chapter 1) and an absorption model (following the method described in reference [Chen et al. 2021b]; [Chen et al. 2021a]), to accurately determine the CdSe x Te 1-x main electronic parameters. Above the bandgap, the absorption coefficient α is modelled by a parabolic band approximation defined by a value α c = 30000 cm -1 at an energy E 0 = 200 meV above the bandgap.

A.2.2.2 STEM-EDS

A very thin lamella of the CdSe 0.2 Te 0.8 sample was extracted using with a Focused Ion Beam (FIB). STEM-EDS was conducted on the CdSe 0.2 Te 0.8 sample. Transmission electron microscopy (TEM) images were acquired at 200 kV. The system equipped with an EDX detector was used to investigate elemental distribution maps along the lamella. Figure A.8 shows a crosssectional STEM image of the Glass/FTO/MZO/CdSe 0.2 Te 0.8 /CdTe:Cu/ device, along with the respective EDS elemental maps of Cd, Se, and Te. EDS elemental maps evidences the higher Se content near the MZO front interface. From the scan plot we conclude that the atomic concentration of Se is in the order of 5-7%, hence between 10 and 14% of Se at the front (substituting to Te element). Se content (%)

A.2.2.3 Diffusion coefficient of Se atoms in CdSeTe layers

x =0.1 x =0.3 x =0.2 It is used in Chapter 6 as a reference for the implementation of a light-trapping strategy in this material. Figure A.12(b) shows maps of J th in the active layer as a function of the Ag nanopillars width and the grating period for several refractive indices (from 1.2 to 2 by step of 0.1) taken by the dielectric layer (optically transparent). The other parameters, h (the thickness of Ag above the nanostructure) and h n , the height of the nanostructures are kept constant, equal to 200 nm and 130 nm respectively. par une faible tension en circuit ouvert et une faible efficacité radiative externe. C'est en partie justifié par leur nature polycristalline complexe et le défi que représente un bon contact arrière. Dans cette thèse, nous présentons différentes stratégies visant à optimiser, tant sur le plan optique qu'électrique, la conception et la fabrication de cellules solaires ultrafines à base de GaAs et de CdTe. Le manuscrit est structuré en six chapitres.

Tout d'abord, nous détaillons dans le chapitre 1 le contexte de cette thèse et les objectifs poursuivis. Les progrès toujours plus importants dans le domaine des cellules solaires ultrafines en GaAs nous ont poussés à développer de nouvelles architectures pour atteindre des rendements toujours plus élevés. Les principales limitations du premier lot de cellules solaires ultrafines en GaAs développées par notre équipe sont d'abord étudiées théoriquement. Pour dépasser les 20%, nous proposons d'abord de modifier l'architecture de la face arrière afin d'améliorer l'absorption de la lumière puis d'utiliser une hétérojonction comme absorbeur (chapitre 2). Les premiers résultats expérimentaux d'une cellule solaire de GaAs ultrafine avec un nouveau design en face arrière sont présentés dans le chapitre 3. Le deuxième objectif de cette thèse est d'appliquer ces nouveaux concepts de piégeage optique au CdTe. Cependant, les problèmes de passivation et de contact arrière sont des obstacles à la réalisation de cellules solaires ultrafines performantes en CdTe. Ainsi, avant de s'intéresser au CdTe ultrafin, nous avons réalisé dans un premier temps une analyse structurelle du matériau CdSeTe par cathodoluminescence dans le chapitre 4. Dans un second temps différents contacts arrière ont été expérimentés sur des échantillons de plusieurs épaisseurs afin d'améliorer la tension en circuit ouvert (chapitre 5). Enfin, dans le chapitre 6, nous montrons les résultats des simulations optiques sur des cellules solaires ultrafines en CdTe avec un piegeage optique optimisé.

Cellules solaires ultrafines en GaAs

La structure et les performances de la première génération de cellules ultrafines en GaAs développées par les membres de l'équipe ont été briévemment rappelé (Figure B.1(a)). Afin d'améliorer les performances de cette première génération de cellules solaires ultraminces en GaAs, une analyse complète des pertes par absorption à l'aide de la méthode RCWA (Rigorous coupled wave analysis) a été réalisée dans le Chapitre 2. À partir de ces simulations optiques, deux stratégies ont été explorées afin d'améliorer la densité de courant de court-circuit (J sc ). Le miroir arrière nanostructuré a été identifié comme l'une des principales sources d'absorption parasite et notre première approche s'est axée sur la modification de la face arrière des cellules ultrafines. J'ai introduit une nouvelle architecture arrière, en utilisant un réseau de diffraction bidimensionnel nanopatterné composé d'un semi-conducteur à indice de réfraction élevé (n(λ=800 nm)=3.87) et une bande interdite relativement élevée (1.77 eV) a-Si:H et d'un diélectrique transparent et mésoporeux le SiO 2 d'indice de réfraction n= 1,2. L'utilisation du SiO 2 comme "espaceur" entre le réseau et le miroir d'argent permet de réduire les pertes dans le miroir d'argent plan. Une absorption multi-résonante à large bande très élevée a été obtenue dans la gamme spectrale de 700 nm à 920 nm, conduisant à une densité de courant de court-circuit théorique de 28,2 mA/cm 2 . Malgré une certaine absorption parasite dans l'a-Si:H, l'augmentation nette de J sc par rapport à la structure originale correspond à 2,63 mA/cm Dans une seconde approche, pour réduire l'absorption parasite dans la couche de passivation arrière et dans le miroir, ainsi que pour augmenter la tension de circuit ouvert (V oc ), j'ai étudié une hétérostructure composée d'une couche émittrice de 200 nm n-GaAs et d'une base de 50 nm p-AlGaAs. En supposant une collecte parfaite des porteurs photogénérés à l'intérieur de la base (dispositif de 250 nm d'épaisseur), la nouvelle structure modélisée peut atteindre une desnité de courant de court-circuit d'environ 28,9 mA/cm

Films fins polycristallin de CdSeTe/CdTe

Dans la deuxième partie de ce travail de thèse, nous avons étudié les cellules solaires en Cd-SeTe/CdTe. Tout d'abord, nous avons réalisé une analyse quantitative sur la structure des films minces CdSeTe/CdTe par cathodoluminescence. Les cellules solaires à base de CdTe ont largement bénéficié de l'introduction du sélénium dans la couche de l'absorbeur. Cela a permis d'augmenter le J sc de manière significative en diminuant le gap en énergie ainsi que d'augmenter la durée de vie des porteurs tout en gardant un V oc similaire aux cellules de CdTe (et donc une diminution du déficit de V oc ). Dans le chapitre 4, une étude systémique sur cinq échantillons de CdSeTe/CdTe avec différentes concentrations de sélénium a été menée par cathodoluminescence et nous a permis de quantifier le rôle de passivation du Se. Des biseaux ont été préparés pour l'analyse en profondeur des efficacités radiatives et des propriétés optiques au sein des films. Des mesures de CL hyperspectrales hautes résolutions ont été acquises à neuf températures différentes dans la même zone d'intérêt.

L'énergie du gap a été extraite à l'aide des spectres moyennés sur la région CdSeTe et j'ai fourni un ensemble d'équations pour relier l'energie du gap avec la proportion nominale en Se quelque soit la température (Figure B.3(a)). J'ai démontré que l'efficacité radiative est décuplée entre x= 0 (CdTe) et x= 0,4 et nous avons corrélé la luminescence CL élevée avec la teneur en Se, confirmant son rôle de passivation (Figure B.3(b)). Nous montrons que l'efficacité radiative s'améliore avec la teneur en Se, quelles que soient les propriétés structurelles de l'échantillon.

Des cartographies de concentration en Se ont été tracées grâce à la position du pic de luminescence de CL. J'ai réalisé une analyse détaillée pour quantifier le coefficient de diffusion des atomes de Se à la fois à l'intérieur des grains et aux joints de grains pendant le recuit au CdCl 2 pour l'échantillon avec une teneur nominale en Se x= 0,2 ). J'ai montré que que les atomes de Se diffusent majoritairement dans le CdTe par les joints de grains, avec un coefficient de diffusion plus de six fois plus élevé aux joints de grains qu'à l'intérieur des grains. Cela confirme également la passivation des joints de grains. De cette analyse, une méthode robuste et plus simple pour extraire le coefficient de diffusion des atomes de Se dans tous les autres échantillons a été implémentée: l'extraction de spectres le long d'une ligne parcourant le biseau suffit (Figure B.3(i)).

Des expériences CL préliminaires dans l'IR et à basse température ont été réalisées pour évaluer l'effet des différentes structures, des contacts arrière ou des couches de passivation sur les dispositifs à base de CdSeTe.

Dans le Chapitre 5, j'ai étudié plusieurs contacts arrière pour les cellules solaires Cd-SeTe/CdTe fabriquée au NREL, tout en tenant compte de la nécessité qu'ils soient ohmiques, passivants et optiquement réfléchissants. Les performances des dispositifs standard (3-4 µm)

Cellules solaires ultrafines en CdTe

Une stratégie de piégeage optique similaire à celle du pour les cellules solaire ultrafines en GaAs a été implémentée pour le CdTe (200 nm) dans le Chapitre 6. On a pris soin de présenter différentes architectures de piégeage optique avec des solutions qui répondent aux critères exigeants d'un bon contact arrière pour le CdTe: contact ohmique avec une réflectivité optique élevée et compatible avec le processus de fabrication. J'ai présenté une structure optimisée avec un réseau de diffraction a-Si:H/TiO 2 associé à un miroir plan qui permet d'obtenir un courant de court-circuit théorique de 25,7 mA/cm 2 pour un dispositif CdTe et de 28,1 mA/cm 2 pour un dispositif CdSeTe (

Conclusion

Dans cette thèse, les progrès réalisés sur les cellules solaires ultrafines en GaAs et CdTe on été présentés.

Les premiers résultats obtenus avec une cellule solaire ultrafine en GaAs avec un réseau de diffraction en a-Si:H/SiO 2 ont démontrés qu'une simple optimisation du processus de fabrication devrait permettre de dépasser le J sc des cellules ultrafines de première génération. De plus la fabrication de cellules solaires ultrafine basées sur une heterostructure permettra l'amélioration du V oc .

J'ai caractérisé les propriétés du CdSeTe par cathodoluminescence et une étude quantitative du rôle de passivation du sélénium a été réalisée, mais nous pourrions aller plus loin et déterminer la valeur absolue de l'efficacité radiative dans cette couche grâce à une calibration absolue du dispositif de cathodoluminescence. Dans ce but, une diode de calibration en GaAs a été fabriquée. Le flux de photons émis par la diode a d'abord été évalué à l'aide d'un imageur hyperspectral calibré, puis comparé au signal d'électroluminescence recueilli dans le dispositif de cathodoluminescence. Un premier rapport de conversion de 1.5 × 10 5 photons.cts -1 a été estimé. La répétition de ces étapes pour d'autres diodes d'étalonnage émettant à différentes longueurs d'onde permettrait de construire un modèle robuste pour la mise en place de la calibration absolue de la CL.

Enfin, les simulations optiques prédisent un brillant avenir pour les cellules solaires de CdTe ultrafines. Ces dispositifs contenant dix fois moins de matériau que les cellules standards et ayant un rendement de près de 20% nous encouragent à poursuivre cette approche.

Figure 1 :

 1 Figure 1: Estimated shares in the total energy consumption in 2009, 2019 and 2020 according to IEA. The share of modern renewable energies is highlighted. Figure reprinted from [REN21 2022].

1

 1 Operating principle . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.2 Characterization of solar cells . . . . . . . . . . . . . . . . . . . . . . . . 1.2.3 Overview of PV technologies . . . . . . . . . . . . . . . . . . . . . . . . 1.3 Overview of GaAs and CdTe solar cells . . . . . . . . . . . . . . . . . 1.3.1 GaAs and III-V solar cells . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3.2 CdTe thin films solar cells . . . . . . . . . . . . . . . . . . . . . . . . . 1.4 Aim of the thesis: towards ultrathin solar cells . . . . . . . . . . . .

Figure 1 . 1 :

 11 Figure 1.1: (a) Solar PV global capacity and annual additions between 2011 and 2021, according to IEA. Figure reprinted from [REN21 2022]. (b) Evolution of the module price as a function of the cumulative PV capacity. Reprinted from [Fraunhofer 2023].

Figure 1 . 2 :

 12 Figure 1.2: Schematic band diagram of semiconductor with bandgap E g with incoming photons of different energies.

Figure 1 . 4 :

 14 Figure 1.4: External Quantum Efficiency of a thin CdTe solar cell. The filled areas indicate the different EQE losses. Figure reprinted from[START_REF] Bothwell | Performance Analysis of 0.4-1.2-µm CdTe Solar Cells[END_REF]

Figure 1 . 5 :

 15 Figure 1.5: Power conversion efficiency of several single-junction PV technologies comparing the best laboratory cells to the best laboratory modules. Figure inspired from [Fraunhofer 2023].

Figure 1 . 6 :

 16 Figure 1.6: (a) Scheme of a typical stack of a GaAs solar cell. (b) Calculated band diagram for an ultrathin GaAs solar cell at short-circuit under 1 sun illumination (V= 0 V). Conduction (E c ) and valence (E v ) band edges are shown together with the quasi-Fermi level for electrons (blue dashed line, E Fn ), and for holes (red dashed line, E Fv ). Panel (b) is reprinted from [Chen 2018].

Figure 1 . 7 :

 17 Figure 1.7: Theoretical usage of the AM1.5 solar spectrum by Si solar cells on the left and by a multijunction with the corresponding sketch on the right. Thermalization and spectral losses are minimized i the second case. The figure is adapted from [Yastrebova 2007].

Figure 1 . 8 :

 18 Figure 1.8: (a) Scheme of a typical stack of a CdTe solar cell. (b) Plot reporting the record CdTe-based solar cell efficiency over the years.

Figure 1 . 9 :

 19 Figure 1.9: Comparison of the energy payback time of several PV technologies. Adapted from[START_REF] Smith | Life cycle assessment of functional materials and devices: Opportunities, challenges, and current and future trends[END_REF]

  -trapping in semiconductors . . . . . . . . . . . . . . . . . . . . . . 22 2.1.2 State-of-the-art of ultrathin GaAs solar cells . . . . . . . . . . . . . . . 27 2.1.3 Electromagnetic simulations tools used in the thesis . . . . . . . . . . . 33 2.2 Optical losses and advanced light-trapping in 200 nm-thick GaAs solar cell . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36 2.2.1 Parasitic absorption losses in state-of-the-art 200 nm-thick GaAs solar cell . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36 2.2.2 Improved light-trapping scheme . . . . . . . . . . . . . . . . . . . . . . . 37 2.3 New optoelectronic design for ultrathin GaAs solar cell . . . . . . . 44 2.3.1 Heterojunction solar cell for enhanced V oc . . . . . . . . . . . . . . . . . 44 2.3.2 Enhanced optical absorption . . . . . . . . . . . . . . . . . . . . . . . . 44 2.3.3 Final considerations on the optoelectronic design . . . . . . . . . . . . . 48 2.3.4 Required accuracy in the experimental planarization of the nanograting 50 2.4 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52

Figure 2 . 1 :

 21 Figure 2.1: (a) Absorption coefficient α, determined from the refractive index of several semiconductor materials, at room temperature as a function of the wavelength. (b) Effect of the slab thickness on the GaAs single-(solid lines) and double-pass (dashed lines) absorptions (red line 2 µm, orange line 1 µm and green line 200 nm).

Figure 2 . 3 :

 23 Figure 2.3: Double-pass absorption scheme (reprinted from [Massiot et al. 2020])

Figure 2 . 4 :

 24 Figure 2.4: Lambertian scattering scheme for a maximal optical path enhancement factor F (reprinted from [Massiot et al. 2020])

Figure 2 . 5 :

 25 Figure 2.5: Schematics for vertical Fabry-Perot resonances in the active layer on a periodical nanostructured mirror (reprinted from [Massiot et al. 2020])

Figure 2 . 6 :

 26 Figure 2.6: Field passing through a grating (a) away from resonances and (b) at resonances (reprinted from [Rumpf 2020]). (c) Multiresonant absorption in the active layer with Fabry-Perot and diffracted light coupled into guided mode resonances (reprinted from[START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]

Figure 2 . 7 :

 27 Figure 2.7: State-of-the-art of ultrathin GaAs solar cells, reprinted and updated from[START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. Short-circuit current of ultrathin GaAs solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models previously described. The values mentioned in the manuscript are in red and labeled. a.[START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]], b.[START_REF] Liu | Surface plasmon enhanced GaAs thin film solar cells[END_REF], c. Sayre et al.[START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF]], d. Camarillo Abad et al. [Camarillo Abad et al. 2020], e. Camarillo Abad et al. [Camarillo Abad et al. 2022], f.[START_REF] Buencuerpo | Engineering the reciprocal space for ultrathin GaAs solar cells[END_REF], g.[START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF], h.van Eerden et al. [Eerden et al. 2020a], i.[START_REF] Lee | High Performance Ultrathin GaAs Solar Cells Enabled with Heterogeneously Integrated Dielectric Periodic Nanostructures[END_REF] , j. Chen et al.[START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF], k.[START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF], l.[START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF], m.[START_REF] Massiot | Metal Nanogrid for Broadband Multiresonant Light-Harvesting in Ultrathin GaAs Layers[END_REF]].

Figure 2 . 8 :

 28 Figure 2.8: Ultrathin solar cell, based on a 205 nm-thick GaAs absorber with a nanostructured Ag mirror. (a) JV curves under 1 sun illumination and (b) EQE, the black curve is the specular reflection 1-R. Figure reprinted from[START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]

Figure 2 . 9 :

 29 Figure 2.9: Scheme of the typical structure used with Reticolo to compute the electromagnetic field inside the device. It contains a stack of layers, with a two-dimensional periodic nanostructures. (a) Cross-section of the device structure, (b) perspective view of the nanostructured layer with p x and p y the period in the x-and y-direction respectively. An unit pattern of widths d x and d y is delimited with the dashed lines. (c) Top-view of the nanostructured layer with the axis orientation for an incident plane wave polarized in TM mode (green) and in TE mode (black).

Figure 2 . 10 :

 210 Figure 2.10: Sketch of the different work axis developed in Sections 2.2 and 2.3, based on the original ultrathin GaAs solar cell developed in the team.

  Figure2.12: Schematic of an ultrathin GaAs cell with a novel nanostructured back contact including a grating and a planar reflector on the rear side. The scheme on the right shows the details of the nanostructured layer with a-Si nanopillars embedded in a dielectric nanogrid. The red parameters: height of the grating h N , period p and width w of the nanopillars, refractive index and thickness under the grating of the dielectric n and h will be optimized in the numerical simulation.

Figure 2 . 13 :

 213 Figure 2.13: Complex refractive indices n+iκ of materials. (a) Real and (b) Imaginary part of the optical indices of materials used for the simulation of ultrathin GaAs cells. Ta 2 O 5[START_REF] Gao | Exploitation of multiple incidences spectrometric measurements for thin film reverse engineering[END_REF]], Al 0.51 In 0.49 P[START_REF] Schubert | Open-air printing of Cu2O thin films with high hole mobility for semitransparent solar harvesters[END_REF]], Al 0.4 Ga 0.6 As and MgF 2[START_REF] Palik | Handbook of Optical Constants of Solids Cited on page 39[END_REF]], GaAs[START_REF] Sturge | Band Alignment for Rectification and Tunneling Effects in Al2O3 Atomic-Layer-Deposited on Back Contact for CdTe Solar Cell[END_REF]];[START_REF] Palik | Handbook of Optical Constants of Solids Cited on page 39[END_REF], Ag[Jiang et al. 2016],a-Si:H.

  Total absorption, J tot = 31.4 mA/cm 2 New design, GaAs, J th = 28.2 mA/cm 2 Original design, GaAs, J th = 25.6 mA/cm 2 New design, Ag, dJ = 0.35 mA/cm 2 Original design, Ag, dJ= 2.7 mA/cm 2 a-Si:H, grating, dJ = 0.61 mA/cm2

Figure 2 .

 2 Figure 2.14: (a) Absorption spectrum in GaAs of the new design of ultrathin GaAs cells with a-Si:H/dielectric grating and a planar mirror (solid yellow line) compared to the absorption spectrum of the initial structure (dashed yellow line). The total absorption of the new structure is plotted in black. The absorption spectra in the grating and in the silver layer are plotted in purple and gray respectively. (b) Updated detailed loss analysis for J sc , the colors match those in (a).

Figure 2 . 15 :

 215 Figure 2.15: Dispersion diagram of the calculated total absorption (1-R) and the corresponding current density J tot calculated for all spectra. (a) Absorption spectra as a function of the dielectric layer refractive index. The horizontal black line marks the index value used in the optimized simulation. The vertical black dashed line indicates the bandgap of GaAs. (b) Absorption spectra as a function of the grating period. The horizontal black dashed line marks the optimized grating period p = 750 nm. These diagrams allow to analyze the resonances and their behavior when a certain parameter is varied.

Figure

  Figure 2.16(a) presents the total absorption as a function of the incidence angle for both TE and TM polarizations. At short wavelengths, little variation with the incidence angle is observed. It becomes more important above 600 nm, especially for the TM polarization. As for the initial structure, we see in Figure2.16(b) that with our improved back grating, the short-circuit current J th is above 27 mA/cm 2 up to an angle of 60°.

Figure 2 . 16 :

 216 Figure 2.16: (a) 2D-diagram of the total absorption as a function of the incident angle for TE and TM polarizations. (b) Polar plot of the theoretical short-circuit current calculated from the total absorption for different incident angles.

Figure 2 .

 2 Figure 2.17: (a) Heterostructure n-GaAs/p-AlGaAs with a rear side grating composed of Al x Ga 1-x As nanopillars in a transparent dielectric nanogrid. (b) Absorption spectra in the GaAs (solid line) and in the grating (dashed line) as a function of the Al content in the Al x Ga 1-x As nanopillars.

Figure 2 . 18 :

 218 Figure 2.18: Absorption spectra in the main absorber GaAs, in the base and in the "dead zone" referred as BSF for the following stack: n-GaAs/p-Al 0.4 Ga 0.6 As/Al 0.6 Ga 0.4 As BSF with a grating with Al 0.6 Ga 0.4 As nanopillars. The respective thickness of each layer is (a) 200/80/20 nm and (b) 200/50/50 nm: Configuration 1.

Figure 2 .Figure 2 .

 22 Figure 2.19: (a) Absorption spectra in the main absorber GaAs, for several Al content in the Al x Ga 1-x As base layer. (b)Configuration 2: absorption spectra in the main absorber, in the base and in the BSF for the following stack: n-GaAs/p-Al 0.6 Ga 0.4 As/BSF Al 0.6 Ga 0.4 As with a grating of Al 0.6 Ga 0.4 As nanopillars. The respective thickness of each layer is 200/50/50 nm.

2 ( 4 ,Figure 2 . 20 :

 24220 Figure 2.20: Absorption spectra in the main absorber GaAs, in the base and in the BSF for the following stack: (a) Configuration 3: n-GaAs/p-GaInP/BSF Al 0.8 Ga 0.2 As with a grating of Al 0.8 Ga 0.2 As nanopillars and (b) Configuration 4: n-GaAs/p-GaInP/BSF Al 0.6 Ga 0.4 As with a grating of Al 0.6 Ga 0.4 As nanopillars with respective thicknesses of 200/50/50 nm

Figure 2 . 21 :

 221 Figure 2.21: Theoretical short-circuit current J th in the main absorber (GaAs) as a function of the spacer thickness in between the grating and the planar mirror.

Figure 2 .

 2 Figure 2.22: Absorption spectrum in the active layer depending on the mirror distortion, with its corresponding J th

Figure 2 .

 2 Figure 2.23: State-of-the-art of ultrathin GaAs solar cells, reprinted and updated from[START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. Short-circuit current of ultrathin GaAs solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models previously described. The values mentioned in the manuscript are in red and labeled. a.[START_REF] Muthukumarasamy | Structural phase change and optical band gap bowing in hot wall deposited CdSexTe1-x thin films[END_REF]], b.[START_REF] Liu | Surface plasmon enhanced GaAs thin film solar cells[END_REF], c.[START_REF] Sayre | Ultra-thin GaAs solar cells with nanophotonic metal-dielectric diffraction gratings fabricated with displacement Talbot lithography[END_REF]], d. Camarillo Abad et al. [Camarillo Abad et al. 2020], e. Camarillo Abad et al.[Camarillo Abad et al. 2022], f.[START_REF] Buencuerpo | Engineering the reciprocal space for ultrathin GaAs solar cells[END_REF], g.[START_REF] Buencuerpo | Efficient light-trapping in ultrathin GaAs solar cells using quasi-random photonic crystals[END_REF], h.van Eerden et al. [Eerden et al. 2020a], i.[START_REF] Lee | High Performance Ultrathin GaAs Solar Cells Enabled with Heterogeneously Integrated Dielectric Periodic Nanostructures[END_REF] , j. Chen et al.[START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF], k.[START_REF] Kayes | 27.6% Conversion efficiency, a new record for single-junction solar cells under 1 sun illumination[END_REF], l.[START_REF] Micha | Insights into the light trapping mechanisms in 100 nm-thick nanopatterned GaAS solar cells[END_REF], m.[START_REF] Massiot | Metal Nanogrid for Broadband Multiresonant Light-Harvesting in Ultrathin GaAs Layers[END_REF]. The blue star is obtained by modifying the back architecture of the record cells, and the purple star by transforming the homojunction to a heterojunction.

1

 1 Optoelectronic improvement . . . . . . . . . . . . . . . . . . . . . . . . 54 3.1.1 Electronic characteristics and edges recombination . . . . . . . . . . . . 54 3.1.2 Metallic contacts losses analysis . . . . . . . . . . . . . . . . . . . . . . . 54 3.2 Fabrication process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56 3.3 Characterization of ultrathin GaAs solar cell . . . . . . . . . . . . . . 61 3.3.1 Preliminary results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61 3.3.2 Performance assessment . . . . . . . . . . . . . . . . . . . . . . . . . . . 63 3.3.3 Impact of the cell size and its front grid design . . . . . . . . . . . . . . 66 3.4 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

Figure 3 . 1 :

 31 Figure 3.1: Sketch of a front contacts design (a) for the record ultrathin GaAs cells and (c) after theoretical optimization. Plots of the different sources of losses: shading, series resistance in the emitter layer, series resistance in the fingers, contact resistance between the metal and the emitter layer under 1 sun illumination as a function of the finger spacing S. The total loss is the black curve. (b) For n-GaAs emitter with R sq = 230 /x 2 , L = 0.2 cm, w f = 6 µm, w b = 60 µm and (d) R sq = 230 /x 2 , L = 1 cm, w f =5 µm, w b = 30 µm

Figure 3 . 2 :

 32 Figure 3.2: Sketches of the main fabrication steps of ultrathin GaAs solar cells with a nanostructured grating and a planar back mirror. (a) Ti/Au localized back contacts deposition. (b) Nanostructuration of the amorphous silicon layer by electron beam lithography, and reactive ion etching. (c) Planarization by a silicon oxide sol-gel. (d) Silver mirror deposition. (e) Transfer on a glass substrate, and substrate removal by wet chemical etching. (f) Front Ni/Ge/Au contact deposition and mesa definition.

  Figure 3.3(c) shows a photograph of the sample after the EBL and RIE steps, and exhibits the diffraction effects on the visible light.

Figure 3 . 3 :

 33 Figure 3.3: Pictures at several intermediate steps of the fabricated samples. (a) Optical microscope image showing the Ti/Au contacts after the lift-off. The inset is a zoomed-in image of the contacts. (b) SEM image, tilted view of the nanostructured a-Si:H layer, forming a 2-dimensional diffraction grating of period 750 nm. The larger patterns are the metallic localized back contacts. The inset shows a zoomed-in SEM image of the grating. (c) Light diffraction effect visible after the nanostructuration of the a-Si:H layer.

Figure 3 . 4 :

 34 photolithography after a-Si nanopatterning

Figure 3 . 5 :

 35 Figure 3.5: (a) C2853b-12 sample with a nanostructured back side consisting of an a-Si:H/SiO 2 grating and a planar Ag mirror. (b) The mask pattern used for front contact grids and solar cells labeling.

Figure 3 . 6 :

 36 Figure 3.6: (a) Current-Voltage (JV) curves of an ultrathin GaAs solar cell with a light trapping structure at the back side under one-sun illumination (red line) and in the dark (black line). (b) Corresponding EQE, the equivalent current at 1 sun is given.Sample [C2853b-12]

Figure 3 . 7 :

 37 Figure 3.7: (a) J sc under solar simulator, the shading part represent 6% to 9% depending on the finger spacing and the cells size. (b) V oc , the outlier point of 1×1 mm 2 cell is V oc =0.52 V. (c) Fill factor. (d)Efficiency. Fitted parameters of sun JV curves using the two diode-model with diode ideality factors fixed to 1 and 2. Fit parameters extracted from the two-diode model fit under 1 sun illumination (e) J 02 and (f) R s . Sample[C2853b-12]

Figure 3 . 8 :

 38 Figure 3.8: EQE of the solar cells (a) on a GaAs substrate [C2853-2], (b) on a flat Ag mirror [C2853b-9], (c) on a nanostructured TiO 2 /Ag mirror (from H.L. Chen manuscript) (blue line) and on a nanostructured back side consisting of an a-Si:H/SiO 2 grating and a flat mirror (red line) [C2853b-12]. The current density is given for each structure. The corresponding sample stacks and the different light-trapping strategies are indicated.

Figure 3 .

 3  shows the EQE plots and the corresponding structure sketches based on a

Figure 3 . 9 :

 39 Figure3.9: Comparison of the EQE of a solar cell with a nanostructured a-Si:H/SiO 2 back side followed by a flat mirror and the calculated absorption in the GaAs, taking into account the real shape of the grating, and the fabrication constraints (seeTable 3.3). (a) with no DLARC, (b) with a DLARC. Equivalent currents at 1 sun (J sc for EQE, J th for absorption in the GaAs absorber) are given.

4 :Figure 3 . 10 :

 4310 Figure 3.10: Front contact grid designs (a) early design, with 3x3 mm 2 , 2x2 mm 2 , and 1x1 mm 2 cells. (b) Optimized design with 5x5 mm 2 cells. (c) Early design with 5x5 mm 2 cells.

Figure 3 .

 3 Figure 3.11: (a) J sc under solar simulator (b) V oc , (c) Fill factor, (d) Efficiency. Fit parameters extracted from the two diode-model fit under 1 sun illumination with diode idealities factors fixed to 1 and 2, (e) R s . The performances are extracted from ultrathin GaAs solar cells transferred on a planar Ag mirror, and the group characteristics are summarized inTable 3.4. Sample [C2853b-9] (ABC),[C2853b- 4] (D1),[C2853b-5] (D2),[C2853b-13] (D3)

  -of-the-art of CdSe x Te 1-x material for photovoltaics . . . . . . 72 4.2 Cathodoluminescence characterization of CdSe x Te 1-x layers . . . . . 74 4.2.1 CL characterization setup . . . . . . . . . . . . . . . . . . . . . . . . . . 74 4.2.2 CdSe x Te 1-x samples fabrication and experimental conditions . . . . . . . 77 4.2.3 Hyperspectral cathodoluminescence maps . . . . . . . . . . . . . . . . . 78 4.3 Selenium and temperature dependence of the bandgap . . . . . . . 81 4.3.1 Bandgap and Urbach tail determination at T=296 K . . . . . . . . . . . 81 4.4 Impact of the Se concentration on the radiative efficiency . . . . . . 86 4.4.1 CL luminscence vs. Se content relation . . . . . . . . . . . . . . . . . . 86 4.4.2 Insights on defect passivation with Se . . . . . . . . . . . . . . . . . . . 87 4.5 Gradient of Se . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 92 4.5.1 Gradient of Se concentration in the depth . . . . . . . . . . . . . . . . . 92 4.5.2 Se diffusion along grain boundaries . . . . . . . . . . . . . . . . . . . . . 93 4.6 Investigation of high quality CdSe x Te 1-x devices . . . . . . . . . . . . 96 4.7 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99

Figure 4 . 1 :

 41 Figure 4.1: (a) EQE of CdSe x Te 1-x /CdTe solar cells with Se composition ranging from 0% to 40%, reprinted from [Ablekim et al. 2021]. (b) A SIMS elemental map of the selenium distribution on the bevelled surface of the CdSeTe/CdTe sample, chlorine signal is in white to help delineate the grain boundaries. (c) A CL map of the panchromatic CL intensity, taken on the same area as the SIMS map. Figure reprinted from[START_REF] Fiducia | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]].

  CL measurements were performed using the Attolight Allalin 4027 Chronos quantitative cathodoluminescence microscope * .

Figure 4 . 2 :

 42 Figure 4.2: Schematics of the CL setup. Overview of the Attolight cathodoluminescence system including an electron microscope (red frame) and a spectrometer/ CCD camera corresponding to the detection system (blue frame). Gray dashed frames show the different steps of the sample image: 1. SEM image, 2. secondary image at the entrance slit, 3. final image on the focal plane of the spectrometer, with a CL spectrum at each pixel, 4. CL response at the zero-order without correction, 5. CL spectra when the spectral shift is corrected. Figure inspired from[START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF];[START_REF] Bidaud | Characterization of nanomaterials by cathodoluminescence for photovoltaic applications[END_REF]].

Figure 4 . 3 :

 43 Figure 4.3: Results of CASINO simulation for a 6 kV acceleration voltage and 10 nm beam radius in CdTe. Distribution of the dissipated energy projected on a plane that is perpendicular to the surface of the sample. Each contour plot indicates a region corresponding to a given percentage of the total dissipated energy.

  Figure 4.4(a) shows a sketch of a CdSeTe/CdTe cell, where the CdSeTe region is close to the front interface with MZO and the CdTe part is near the back contact of the cell.

Figure 4 . 4 :

 44 Figure 4.4: (a) Sketch of graded CdSe x Te 1-x sample: Glass/FTO/MZO/CdSeTe/CdTe/Au contact. (b) Sketch of CdSeTe/CdTe device structure and bevel, with a top view SEM image of CdSe x Te 1-x stack (x=30%), where the whole depth is exposed. The white dashed line indicates the CL measurement region.

Figure 4 . 5 :

 45 Figure 4.5: (a-e) Room-temperature panchromatic CL maps (256×256 pixels) of CdSe x Te 1-x samples with compositional x values of (a) x= 0, (b) x=0.1, (c) x=0.2, (d) x=0.3 and (e) x= 0.4. The grayscale bars are adjusted for each CL intensity maps. (f-j) Corresponding CL peak energy maps of (f) x= 0, (g) x=0.1, (h) x=0.2, (i) x=0.3 and (j) x= 0.4. White scale bar: 3 µm.

Figure 4 . 6 :

 46 Figure 4.6: Determination of the energy map from the CL hyperspectral cube. For each pixel the peak energy of the main luminescence peak is extracted.

Figure 4 .

 4 Figure 4.7: (a) Normalized CL spectra averaged over the homogeneous area shown as colored rectangle in Fig.4.5(a-e) at 296 K for the five samples (solid curves). The dashed curves correspond to the generalized Planck law fit, open circles on the CL spectra indicate the bandgaps determined from the fit. The same figure plotted in logscale is shown in Figure A.6 in Appendix A.2.1. (b) The absorption coefficient of CdTe calculated with the refractive index provided by the NREL (solid line) is compared to the absorption coefficient calculated with equations 4.2 and 4.3 (dashed line). The inset is in logscale. Similar plots for different Se contents are plotted in Figure A.5 in Appendix A.2.1

Figure 4 . 8 :Table 4 . 3 :

 4843 Figure 4.8: (a) Normalized CL spectra averaged over the homogeneous areas shown as colored rectangles in Figure 4.5(a-e) at 15 K for the five samples. (b) CL peak position values for several Se contents as a function of the temperature. The CdTe CL peak positions (triangles) were extracted over the CdTe area of the samples while the CdSeTe peak positions (circles) were extracted focusing on the area with the highest amount of Se. Corresponding colored dashed lines stand for the Varshni law fit associated to each Se content

Figure 4 . 9 :

 49 Figure 4.9: Evolution of the CdSe x Te 1-x bandgap as a function of the Se composition x. Data points were fitted with equation 4.5. Green values were extracted from optical data[START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF], the red ones using CL spectra at RT, and dark blue values are determined using the bandgap values at 0 K extracted from the fit of experimental data with the Varshni law(equation 4.4). The bowing factor b is explicited for each fit.

Figure 4 .Figure 4 .

 44 Figure 4.10: (a) Scatter plot of the maximum CL intensity as a function of the peak energy for the CdSe 0.4 Te 0.6 sample. (b) Room-temperature CL intensity as a function of the CL peak position combined from the four CdSe x Te 1-x samples with Se content ranging from 0.1 to 0.4. The mean CL intensity is calculated for each energy range and sample.

Figure 4 .Figure 4 . 12 :

 4412 Figure 4.11: CL spectra averaged on 20×256 pixels at several temperatures for five distinct Se concentrations. The CL spectra of the first rectangle are integrated on the CdTe area (a) x= 0, (b) x= 0.1, (c) x= 0.2, (d) x= 0.3 and (e)x= 0.4, while the CL spectra of the second rectangle are integrated on the CdSeTe area (f) x= 0.1, (g) x= 0.2, (h) x= 0.3 and (i)x= 0.4.

Figure 4 .

 4 Figure 4.13: (a-e) CL spectra averaged on the whole surface of the five CdSe x Te 1-x samples with different Se compositions at 20 K. (f-o) CL false color maps showing normalized luminescence for each spectrum band: the defect in blue, the CdSeTe-exciton in red, and the CdTe-exciton in green. The CdSeTe-exciton and the CdTe-exciton maps are superposed to catch at first sight the change of composition within the thin film. Normalization factors have been determined from the maximum integrated luminescence from each spectral bands. White scale bar: 2 µm. Figure reprinted from [Bidaud 2021].

Figure 4 .

 4 Figure 4.14: (a) The correlation of lifetime versus Voc for CdSeTe/CdTe solar cells with Se composition ranging from 0% to 40%. The dashed orange line is a straight line in linear-log plot to guide the eye. Figure reprinted from [Ablekim et al. 2021]. (b) Integrated CL intensity in the defect spectral range for the 200 brightest pixels of each sample as a function of 1/T in logscale. The top graph shows the values found in the CdTe region while the bottom graph shows the intensities evaluated in the CdSeTe region.

Figure 4 . 15 :

 415 Figure 4.15: (a) Depth profiles of the peak energy extracted from the room-temperature CL maps of the CdSe x Te 1-x bevels for various Se contents. (b-d) Maps of the Se concentration x Se extracted from the CL maps of the peak energy (Figs. 4.5(g-i)) and equations 4.6 and 4.7 for nominal compositions of the CdSeTe region of (b) x=0.1, (c) x=0.2 and (d) x=0.3. White scale bar: 3 µm.

  Characteristic diffusion length obtained by fitting the interdiffusion regions for each CdSex Te 1-x sample.Maps of the Se concentration are plotted in Figure4.15(b-d), where the common colorscale highlights the global increase of Se content from (b) to (d). The Se content was also determined in homogeneous regions by SEM/EDS (see Appendix, Figure A.7), and by STEM/EDS (see Appendix, Figure A.8) for the sample CdSe 0.2 Te 0.8

  .

Figure 4 . 16 :

 416 Figure 4.16: (a) RT panchromatic map of CdSe 0.2 Te 0.8 where the analyzed line profile are indicated by the colored dotted lines. (b) Corresponding Se concentration map emphasising a higher Se content in GB. (c) Se concentration profile across one of the grain boundaries (white dotted line in (b)). (d-f) Data and fits of selenium concentration in-depth profiles. The colors of the fit matches the indicated lines in (a).

Figure 4 .

 4 Figure 4.17: (a) Diffusion coefficients and (b) corresponding diffusion lengths for three CdSe x Te 1-x samples (where x=0.1, 0.2 and 0.3) annealed at 420°C during the CdCl 2 post deposition treatment. For CdSe 0.2 Te 0.8 distinction between GI and GB is emphasized. The round ticks in the diffusion length plot correspond to the characteristic diffusion length δ determined in the interdiffusion area.

Figure 4 .

 4 Figure 4.18: (a) Sketches of the four samples under study with their associated name. (b) Corresponding External Radiative Efficiency (ERE) and non radiative voltage loss (∆V oc = kT q ln|ERE|).

Figure 4 .

 4 Figure 4.19: LT CL spectra averaged over the whole bevel for each sample.

Figure 4 .

 4 Figure 4.20: (a) LT CL map of the sample 1809-5L, CdSeTe/CdTe based devices with a passivated back contact, restricted on the interdiffusion zone. (b) LT CL spectra centered on the sub-bangap emission peak as a function of the localisation on the maps.

1

 1 Voltage deficit in polycrystalline CdTe and CdSeTe solar cells . . . . . . 102 5.1.2 Impact of interfaces and back contacts . . . . . . . . . . . . . . . . . . . 104 5.1.3 Notable results on passivation layers for CdTe solar cells . . . . . . . . . 106 5.1.4 Reducing CdSeTe absorber thickness: aim and challenges of this work . 106 5.2 Back contacts for CdSeTe/CdTe solar cells . . . . . . . . . . . . . . . 108 5.2.1 Solar cells fabrication process . . . . . . . . . . . . . . . . . . . . . . . . 108 5.2.2 Highly reflective back contacts: Au and Ag . . . . . . . . . . . . . . . . 111 5.2.3 Role of Al 2 O 3 passivation layers . . . . . . . . . . . . . . . . . . . . . . 114 5.2.4 P-type transparent conductive oxide as a back contact : Cu 2 O . . . . . 117 5.3 Analysis of 800 nm-thick CdSeTe solar cells . . . . . . . . . . . . . . 129 5.3.1 Solar cells performance characterization . . . . . . . . . . . . . . . . . . 129 5.3.2 Photoluminescence response . . . . . . . . . . . . . . . . . . . . . . . . . 132 5.3.3 Spectral responses and simulation comparison . . . . . . . . . . . . . . 133 5.4 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137

Figure 5 . 1 :

 51 Figure 5.1: V oc deficit as a function of the bandgap for record cells[Green et al. 2022a]. The dashed red line represents the limiting value of the bandgap voltage offset V oc deficit of a solar cell using the detailed balance model. It indicates the highest voltage a solar cell of a given band gap can achieve[START_REF] King | Band gap-voltage offset and energy production in next-generation multijunction solar cells[END_REF];[START_REF] Brandt | Determining interface properties limiting open-circuit voltage in heterojunction solar cells[END_REF]

Figure 5 . 2 :

 52 Figure 5.2:Detailed voltage loss analysis for two types of CdSeTe-based devices. The V oc,ideal is extracted from EQE and iV oc is deduced from ERE measurements. The figure is adapted and simplified from[START_REF] Onno | Understanding what limits the voltage of polycrystalline CdSeTe solar cells[END_REF]].

Figure 5 . 3 :

 53 Figure 5.3: Band alignment diagram between the back contact and CdTe semiconductor, showing a Schottky barrier. χ is the electron affinity of the p-type CdTe, Φ m is the work function of the metal measured in vacuum, and Φ SB is the Schottky barrier height, defined as the potential difference between the Fermi energy (E f ) of the metal and the valence band edge. Figure inspired from[START_REF] Kumar | Physics and chemistry of CdTe/CdS thin film heterojunction photovoltaic devices: Fundamental and critical aspects[END_REF]].

Figure 5 . 4 :

 54 Figure 5.4: CdTe solar-cell band diagram for the back contact. They show the potential impact of Cu with (a) a back contact and (b) with a back buffer layer. These figures are based on[START_REF] Corwine | Copper inclusion and migration from the back contact in CdTe solar cells[END_REF]];[START_REF] Li | Controlled activation of ZnTe:Cu contacted CdTe solar cells using rapid thermal processing[END_REF]

Figure 5 . 5 :

 55 Figure 5.5: Schematic of a standard CdSeTe/CdTe sample in superstrate configuration. (a) The MZO layer is first deposited by sputtering on a soda-lime glass coated with FTO. On the right, an SEM image of the FTO surface taken at the rear side of the sample. It is followed by the evaporation of CdSe and CdTe, then a CdCl 2 annealing forms a graded CdSe x Te 1-x bilayer. The SEM on the right highlights the polycrystalline structure of the bilayer with multiple grain sizes. The CuCl 2 treatment performed by the NREL is done before the metallic contact deposition done at C2N, the nomenclature is displayed next to it. A picture of the CdSeTe/CdTe sample with 7 cells on it is shown. The front (FTO, transparent part of the sample) and back contact (golden squares) can be taken at the same side of the sample.

Figure 5 . 6 :

 56 Figure 5.6: Current-Voltage (JV) curves of CdSeTe solar cells under one-sun illumination performed at the NREL. (a) Several standard CdSeTe/CdTe samples are compared with a same CdCl 2 annealing temperature equal to 435℃ (and CuCl 2 equals 200℃). (b) Several 800 nm-thick CdSeTe/CdTe cells at multiple CdCl 2 annealing temperatures are compared (and a CuCl 2 temperature of 200℃).

Figure 5 . 7 :

 57 Figure 5.7: Current-Voltage (JV) curves of CdSeTe solar cells under one-sun illumination (solid lines) and in the dark (dashed lines) showing the comparison between Au [A201014A1_3] and Ag [A201014A1_12] back contacts.

Figure 5 . 8 :

 58 Figure 5.8: EQE of CdSeTe/CdTe solar cells for (a) Au contact [A201014A1_A4] and (b) Ag contact [A201014A1_14] measured at 0 V, and under different bias voltages. Dotted lines represent the ratio of EQE(-/+0.4V)/EQE(0V) for Au contacts and EQE(-/+0.2V)/EQE(0V) for Ag contacts.

Figure 5 . 9 :

 59 Figure 5.9: (a) JV curves under one-sun illumination (solid lines) and in the dark (dashed lines) of CdSeTe solar cells showing the impact of the light soaking on the devices' performances for Au back contact solar cells. (b) Statistics of the performances of the solar cells: J sc , V oc , FF, and efficiency. Several light soaking times were studied. The box plots show the average value (red star) and the box represents the first (25%) and the third (75%) quartile. The error bar gives the maximum and the minimum value. The measurements were carried out on the cells [A102014A1_A1-7]

Figure 5

 5 Figure 5.9(a) shows light and dark IV curves for several light soaking times for CdSeTe/CdTe solar cells with a gold back contact. The main observation is that longer illumination seems to attenuate the rollover. The results conducted on seven cells are summarized in the box plot Figure5.9(b). From 0 min to 10 min exposure, the J sc and FF values slightly increase and seem to stabilize after longer illumination. On the other hand, the V oc decreases until an exposure

Figure 5 . 10 :

 510 Figure 5.10: Current-Voltage (JV) curves of CdSeTe solar cells with (a) Au back contact and (b) Ag back contact. The impact of an Al 2 O 3 passivation layer between the active layer and the back contact is shown. Statistics of the solar cells performances: (c) J sc , (d) V oc , (e) Fill Factor (FF) and (f) Efficiency. The box plots show the average value (red star) and the box represents the first (25%) and the third (75%) quartile. The error bar gives the maximum and the minimum value. Cells come from the sample [A201014A1], Au (A1-7), Ag (A8-14), 1.5 nm Al 2 O 3 +Au (A15-21), 1 nm Al 2 O 3 +Au (A26-28) and 1 nm Al 2 O 3 +Ag (A29-35).

Figure 5 .

 5 Figure 5.11: EQE of typical CdSeTe solar cells comparing Au and Ag back contacts with or without Al 2 O 3 passivation layer. In the order of the legend, the cells are [A201014A1_A12-A33-A3-A27].

Figure 5 . 12 :

 512 Figure 5.12: Comparison of JV sun measurements (yellow dots) and pseudo-JV deduced from J sc -V oc measurements (black dots) for (a) Au back contacts [A201014A1_A4] and (b) 1 nm Al 2 O 3 +Au back contact [A201014A1_A26].

Figure 5 .

 5 Figure 5.13: (a) Sketches of the Cu 2 O deposition steps. (b) SEM images of thin film CdSeTe/CdTe with Cu 2 O deposited on it. It reveals the deposited Cu 2 O crystals. The inset displays a zoomed SEM image, tilted at 45°. The deposited thickness is estimated at ∼ 130 nm. (c) Images of the CdSeTe/CdTe sample after the Cu 2 O deposition, followed by a rank of seven metallic back contacts, with (left) or without (right) Al 2 O 3 .

Figure 5 .

 5 Figure 5.14: Current-Voltage curves of CdSeTe/CdTe solar cells under one-sun illumination (solid lines) and in the dark (dashed lines) and External Quantum Efficiency, comparing the performance of cells with Cu 2 O contact with or without a passivation layer with (a,c) Au and (b,d) Ag back contacts, respectively. The sample under study is [A201014A2].

Figure 5 . 15 :

 515 Figure 5.15: J sc of the CdSeTe/CdTe cells with a p-TCO Cu 2 O layer for several back contacts. J sc values are shown depending on their measurement method: using EQE (plain bars) or IV curves (hatched bars). Results for Au and Ag back contacts with or without Al 2 O 3 layer are shown. These are plotted as a function of the cells position on the sample [A201014A2]

Figure 5 . 16 :Figure 5 .

 5165 Figure 5.16: JV curves of CdTe cells with Cu 2 O back contact (+Au contact). Squares: without any treatment, triangles: after 5-min annealing at 150 ℃ and circles: after a thermal treatment of 25 min at 150 ℃ and a additional 5 min at 200 ℃. Figure reprinted from[START_REF] Türck | Copper (I) Oxide (Cu2O) based back contact for p-i-n CdTe solar cells[END_REF]].

Figure 5 . 18 :

 518 Figure 5.18: Evolution of J sc , FF, V oc and R s (serie resistance in logscale) for the solar cells with Cu 2 O oxide at the rear side as the function of their position on the sample. Devices before (filled triangles) and after (hollow triangles) annealing and with Au (orange) or Ag (grey) metallic back contact are compared. Cells with 1 nm of Al 2 O 3 layer are highlighted with a light green background. J sc are extracted from EQE measurements.

Figure 5 .

 5 Figure 5.19: SEM image tilted at 45°of CdSeTe/CdTe rear side with (a) as deposited Cu 2 O (130 nm) and (b) ∼90 nm of Cu 2 O after 6 min of dry etching. The inset shows SEM tilted at 90°. On the left of each SEM images, a photo of the rear surface of a CdSeTe/CdTe sample with (a) as deposited Cu 2 O and with (b) ∼ 50 nm of Cu 2 O, after 20 min etching.

Figure 5 .Figure 5 .

 55 Figure 5.20: (a) JV sun plot showing the performances of the best CdSeTe devices with a ∼50 nm Cu 2 O+Au back contact with [A201014A2-3_A2] or without [A201014A2-3_A5] an Al 2 O 3 passivation layer in between. Before (solid lines) and after (dashed lines) a 5 min annealing at 300℃ in air [A201014A2-3]. (b) The EQE are plotted for several cells after the annealing step only.

Figure 5 .

 5 Figure 5.21: (a) JV curves under different illumination levels, conducted with neutral density to attenuate the illumination. (b) JV curve under one-sun illumination (red curve) compared to a pseudo-JV curve (hollow black dots). (c) JV curves measured under specific wavelength. These measurements were conducted for a cell with Cu 2 O + Au back contact after a 150 ℃ annealing [A201014A2-0_A5].

2 Figure 5 .

 25 Figure 5.22: EQE measurements under two different monochromator beam sizes for CdSeTe/CdTe devices with Cu 2 O + Au back contacts, (a) with [A201014A2p_A2] or (b) without a passivation layer [A201014A2p_A5]. (c) EQE under two different monochromator beam sizes for CdSeTe/CdTe cell with Au contact only [A201014A1_A6] for comparison. (d)EQEs measured (blue dashed curves) with a reduced beam size localized at distinct location on one cell compared to the EQE with a larger beam size (solid black line) for CdSeTe/CdTe devices with Cu 2 O + Au back contacts [A201014A2p_A4].

Figures

  Figures 5.22(a) and (b) show that higher EQE are measured with a large beam that covers the entire cell than with a small focused beam, and regardless of the presence of the passivation. Annealed samples exhibit similar outcomes. For comparison, EQE measured with only a Au back contact shows almost no dependency on the beam size or displays a slightly improved curve for a smaller beam (Figure5.22(c)). Finally, inFigure 5.22(d), we plot EQEs at several positions on the cells (blue dashed curves) and compare it to the EQE obtained with a large beam (solid black line). The significant dependence on the beam position suggests inhomogeneously distributed defects.

Figure 5 .

 5 Figure 5.23: (a) Current-Voltage (JV) curves of 800 nm-thick CdSeTe solar cells under one-sun illumination (solid lines) and in the dark (dashed lines) for Au and Ag metallic back contacts. (b) EQE of thin CdSeTe solar cells for Au and Ag back contacts with or without an Al 2 O 3 passivation layer. The sample under study is [A201016A1], the cells are [A4,A10] for Au and Ag contacts respectively and [A23,A16] for passivated Au and Ag contacts respectively.

Figure 5 . 24 :

 524 Figure 5.24: Current-Voltage (JV) curves of 800 nm-thick CdSeTe solar cells with (a) Au back contacts and (b) Ag back contacts. The impact of an Al 2 O 3 passivation layer between the active layer and the back contact is also investigated. Statistics of the solar cell performances: (c) J sc (extracted from EQE except for cells with Ag contact), (d) V oc , (e) Fill Factor (FF) and (f) efficiency. The box plots show the average value (red star) and the box represents the first (25%) and the third (75%) quartile. The error bar gives the maximum and the minimum value. The sample under study is [A201016A1]

Figure 5 . 25 :

 525 Figure 5.25: Comparison of opto-electronic characteristics of CdSeTe solar cells. (a) JV curve under one-sun illumination, (b) EQE and (c) optical transmission for two different thicknesses of the Cd-SeTe/CdTe solar cell's active layer (without metal contact): 3450 nm [A201014A1_A3] and 800 nm [A201016A1_A4]

Figure 5 . 26 :

 526 Figure 5.26: Photoluminescence response comparing standard [A201014A1] and thin [A201016A1] Cd-SeTe/CdTe layers. A spectrum of the sample part with 1 nm of Al 2 O 3 at the back is also shown. PL was measured with 532 nm excitation from the back side of the sample.

Figure 5 . 27 :

 527 Figure 5.27: Sketch of the sample layers used for the simulations with TMM.

Figure 5 . 29 :

 529 Figure 5.29: Transmission measurements of (a) standard and (b) thin CdSeTe/CdTe compared to the simulated transmission of an optimized complete solar cell and to a CdTe cell only and CdSeTe cell only for (a) and (b), respectively. EQE compared to the optimized simulation of the CdSeTe/CdTe structure for (c) standard and (d) thin sample. The calculated absorption in the CdTe layer only is plotted for both devices in (c) and (d).

1

 1 Effect of thickness reduction in planar CdTe solar cells . . . . . . . . . . 141 6.2.2 Nanostructured back mirror . . . . . . . . . . . . . . . . . . . . . . . . 144 6.2.3 Alternative architectures for reducing parasitic losses . . . . . . . . . . . 147 6.3 Analysis of the multi-resonant light-trapping in CdSeTe solar cells 150 6.4 Experimental perspectives for ultrathin CdTe cells . . . . . . . . . . 153 6.4.1 Description of fabrication steps . . . . . . . . . . . . . . . . . . . . . . . 153 6.4.2 Cost assessment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 155 6.5 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157

Figure 6 .

 6 Figure6.1 reports non exhaustive results of the state-of-the-art of thin and ultrathin CdTe solar cells. Remarkably, all the experimental results lie below the single-pass absorption, as observed for CIGS solar cells[START_REF] Massiot | Progress and prospects for ultrathin solar cells[END_REF]]. This can be partly justified by the predominance of the CdS layer in theses results (few results with MZO) that generated parasitic absorption at short wavelength.

Figure 6 . 1 :

 61 Figure 6.1: State-of-the-art of ultrathin CdTe solar cells. Short-circuit current of ultrathin CdTe solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models described in Chapter 2. a.[START_REF] Gupta | High efficiency ultra-thin sputtered CdTe solar cells[END_REF]], b. and h.Bothwell et al. [Bothwell et al. 2020], c. and[START_REF] Paudel | Ultrathin CdTe solar cells with MoO3-x /Au back contacts[END_REF]], e.[START_REF] Salavei | Influence of CdTe thickness on structural and electrical properties of CdTe/CdS solar cells[END_REF], f. Amin et al.[Amin et al. 2001], g. and i. Plotnikov et al. [Plotnikov et al. 2011], j.[START_REF] Paudel | Ultrathin CdS/CdTe solar cells by sputtering[END_REF], k.[START_REF] Luk | Surface plasmon polariton enhanced ultrathin nano-structured CdTe solar cell[END_REF]].

Figure 6 . 3 :

 63 Figure 6.3: (a) Calculated absorption in the CdTe layer, plotted as a function of the CdTe thickness. (b) CdTe solar structure used in simulations in superstrate configuration (with a gold back contact in (a)).

Figure 6 . 4 :

 64 Figure 6.4: Absorption spectra of a complete CdTe device calculated for (a,b) a 3450 nm-thick and (c,d) a 200 nm-thick absorber layer. The back contact is either (a,c) in Te/Ni or (b,d) in Au. Absorption in each layer of the stack is shown (colored areas) and J th determined using the CdTe layer absorption is given for each configuration.

Figure 6 . 5 :

 65 Figure 6.5: (a) Simulated absorption spectra with a 200 nm-thick CdTe cells comparing the effect of Au and Ag metallic contact. (b) J th calculated for each layer and for several absorber thicknesses.

  Figure A.11 in the previous chapter.

Figure 6 . 6 :

 66 Figure 6.6: Sketch and corresponding optical simulation results of an ultrathin CdTe cells with a flat Au mirror and a nanostructured silver back mirror. The parameters under study (in red) are: h corresponding to the layer height on top of the nanostructures, h n and w the height and the width of the nanostructures, p the period of the nanopillars.

AgFigure 6 . 7 :

 67 Figure 6.7: Schematics of a 200 nm-thick CdTe solar cells with three different back side architectures. The corresponding optical simulation for each structure are plotted. On the left, the graph displays a comparison of calculated absorption in the 200nm-thick CdTe absorber with and without light trapping structures: (yellow) solar cell on a flat Au mirror, (purple) on an optimized nanostructured back mirror whit different layers. The plots on the right show the absorption spectra in each layer of the solar cells with optimized geometry. The dimension are given inTable 6.2

Figure 6

 6 Figure 6.8: (a) Calculated transmission/absorption of SiO 2 +SnO 2 :F. The spectra were corrected by a 0.96 factor, corresponding to light transmission through bare glass. (b) Original extinction coefficient of FTO, and reduced by a factor 5.

Figure 6 . 9 :

 69 Figure 6.9: (a) Calculated absorption spectra of the structure C (dashed lines), compared to the optimized same structure but with tampered extinction coefficient k FTO/5 (solid lines). J th is shown for each structure. (b) With an updated structure C, comparison of the absorption spectra in the active layer: 200 nm-thick CdTe in dashed line, a 200 nm-thick CdSeTe in solid line

Figure 6 . 11 :

 611 Figure 6.11: Dispersion diagrams showing the total absorption as a function of the CdTe thickness and the grating period. (a) 1-R as a function of the CdTe thickness. The horizontal black dashed line indicate 200 nm and the vertical one marks the CdTe bandgap (period 400 nm, width 200 nm). (b) 1-R as a function of the grating period for a 200 nm CdTe thickness. The horizontal black dashed line indicates the period 400 nm that gives rise the optimum broadband absorption. The width-to-period ratio is kept constant (50%). Dotted lines are calculated dispersion curves of guided-mode resonances.

Figure 6 . 13 :

 613 Figure 6.13: Costs in $/W grouped by major fabrication step for CdTe module production in Southeast Asia. The modules have a 18% efficiency. Figure reprinted from[START_REF] Smith | Photovoltaic ( PV ) Module Technologies : 2020 Benchmark Costs and Technology Evolution Framework Results[END_REF]

Figure 6 . 15 :

 615 Figure 6.15: State-of-the-art of ultrathin CdTe solar cells. Short-circuit current of ultrathin CdTe solar cells as a function of the absorber thickness. Experimental and numerical results are indicated with filled and open circle respectively. Values are compared to the reference absorption models described in Chapter 2. a.[START_REF] Gupta | High efficiency ultra-thin sputtered CdTe solar cells[END_REF]], b. and h. Bothwell et al. [Bothwell et al. 2020], c. and d. Paudel et al. [Paudel et al. 2014], e.[START_REF] Salavei | Influence of CdTe thickness on structural and electrical properties of CdTe/CdS solar cells[END_REF], f. Amin et al.[START_REF] Amin | Highly efficient 1 µm thick CdTe solar cells with textured TCOs[END_REF], g. and i.[START_REF] Plotnikov | Thin-film CdTe cells : Reducing the CdTe[END_REF], j.[START_REF] Paudel | Ultrathin CdS/CdTe solar cells by sputtering[END_REF], k.[START_REF] Luk | Surface plasmon polariton enhanced ultrathin nano-structured CdTe solar cell[END_REF]. The blue and the purple stars are the J sc obtained with a CdTe absorber and a CdSeTe absorber, respectively.

  Figure A.2 shows images of the different steps of this thick bilayer process.

Figure A. 2 :A. 1 . 3

 213 Figure A.2: (a,b) Microscopic images of the front contact lithography step on GaAs substrate using the bilayer process with nLOR30B and AZ5214. (c,d) Tilted SEM images of the bilayer, emphasizing the overhang of the AZ5214 easing the lift-off process. (e) Tilted SEM image of the bilayer and 1.5 µm of gold deposited on it. (f) Sketch of the lithography mask for the front electrode of a 0.5×0.5 cm 2 cell: the width of the fingers is 5 µm and that of the busbars is 30 µm and the area of the two pads is 250×750 µm 2 .

Figure A. 3 :

 3 Figure A.3: Typical JV curve under 1 sun illumination (orange dots) and fit wth a two diode model (red curve).Sample [C2853b-12]

Figure A. 4 :

 4 Figure A.4: (a) Absorption coefficient of CdTe calculated with refractive indices provided by the NREL[START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF] and found in the literature[START_REF] Adachi | Optical properties of CdTe: Experiment and modeling[END_REF];[START_REF] Treharne | Optical design and fabrication of fully sputtered CdTe/CdS solar cells[END_REF]]. (b) Absorption coefficient of CdSe x Te 1-x calculated with refractive indices provided by the NREL depending on the Se concentration[START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]. The graph on the right is a zoomed plot corresponding to the red rectangle, the absorption coefficient α is plotted between 1 eV and 2.5 eV.

Figure A. 5 :

 5 Figure A.5: Absorption coefficient of CdSex Te 1-x calculated with refractive index provided by the NREL[START_REF] Junda | Optical Properties of and Alloys and Their Application for CdTe Photovoltaics[END_REF]] (solid lines) compared to absorption coefficient calculated with equations 4.2 and 4.3 (dashed line), the inset is in logscale for (a) x=0.1, (b) x=0.2, (c) x=0.3, and (d) x=0.4.

Figure

  Figure A.7: SEM image with Se distribution superimposed on it, for four samples with nominal Se concentration ranging from x=0.1 to x=0.4. The associated elemental quantification is plotted next to it. This calculation is conducted for both CdTe and CdSeTe regions.

Figure 4 .

 4 Figure 4.17 reports the diffusion coefficient found for CdSe x Te 1-x samples (x=0.1, 0.2 and 0.3). We show in Figure A.10 along which lines the diffusion coefficients were evaluated. Table A.2 summarizes the values estimated from the erfc fit.

Figure A. 10 :

 10 Figure A.10: Maps of the Se concentration x Se extracted from the CL maps of the peak energy) and equations(4.6,4.7) for nominal compositions of the CdSeTe region of x=0.1, x=0.2, and x=0.3. Dashed vertical line for each map are the linescans along which we evaluate the decrease in selenium concentration and the associated diffusion coefficient.

Figure

  FigureA.12(a) recalls the structure of an ultrathin CdTe cell with a nanostructured back mirror. It is used in Chapter 6 as a reference for the implementation of a light-trapping strategy in this material.Figure A.12(b) shows maps of J th in the active layer as a function of the Ag nanopillars width and the grating period for several refractive indices (from 1.2 to 2 by step of 0.1) taken by the dielectric layer (optically transparent). The other parameters, h (the thickness of Ag above the nanostructure) and h n , the height of the nanostructures are kept constant, equal to 200 nm and 130 nm respectively.

  Figure A.12: (a) Scheme of a 200 nm-thick device with a dielectric/Ag nanostructured mirror at the back side. (b) Maps of theoretical short-circuit current J th , plotted as a function of the nanopillars width and the period of the gratings, for several refractive indices of the dielectric. The calculations was conducted on the structure sketches in (a)

2 (

 2 Figure B.1(b)).

2 (

 2 Figure B.1(c)).

Figure B. 1 :

 1 Figure B.1: (a) Caractéristique courant-tension (JV) sous illumination, montrant les performances de la cellule solaire utrafine record en GaAs de première génération (19.9% d'efficacité), développée par l'équipe. (b) Spectre d'absorption dans la couche de GaAs de la seconde génération de cellules ultrafines en GaAs avec une architecture arrière faite en a-Si:H/SiO 2 et un miroir plan en argent (courbe jaune continue) comparé au spectre d'absorption de la structure initiale (coube jaune en pointillée). L'absorption totale de la nouvelle structure est représentée en noir. Les spectres d'absorption dans le réseau et dans le miroir d'argent sont tracés en violet et en gris respectivement. (c) Spectres d'absorption dans l'emitteur GaAs, dans la base et dans le BSF(couche de passivation arrière) pour l'empilement suivant : n-GaAs/p-Al 0.6 Ga 0.4 As/BSF Al 0.6 Ga 0.4 As avec un réseau de nanostructure en Al 0.6 Ga 0.4 As. L'épaisseur respective de chaque couche est de 200/50/50 nm.

Figure B. 2 :

 2 Figure B.2: (a) Caractéristique courant-tension (JV) d'une cellule solaire ultrafine en GaAs avec pour structure arrière le réseau nanostructuré en a-Si:H/SiO 2 / miroir en argent plan sous illumination à un soleil (courbe rouge) et dans le noire (courbe noire). (b) Réponse spectrale (EQE) correspondantee, la densité de courant de court-circuit à un soleil est donné.

Figure B. 3 :

 3 Figure B.3: (a) Evolution du gap du CdSe x Te 1-x en fonction de la composition en Se. (b) Intensité de CL à température ambiante en fonction de la position du pic CL des quatre échantillons de CdSe x Te 1-x dont la teneur en Se varie entre 0,1 et 0,4. L'intensité CL moyenne est calculée pour chaque gamme d'énergie et chaque échantillon. (c) Carte panchromatique à température ambiante du CdSe 0.2 Te 0.8 où le profil des lignes analysées est indiqué par des lignes pointillées colorées. (d) Carte de concentration en Se correspondante soulignant une teneur en Se plus élevée au joint de grain.(e) Profil de concentration en Se à travers l'un des joints de grains (ligne pointillée blanche en (b)). (f-h) Données et ajustements des profils de concentration de Se en profondeur. Les couleurs de l'ajustement correspondent aux lignes indiquées en (a). (i) coefficients de diffusion de trois échantillons de CdSex Te 1-x (où x=0,1, 0,2 et 0,3) recuits à 420°C pendant le traitement post-déposition de CdCl 2 . Pour les échantillons de CdSe 0.2 Te 0.8 interieur de grain et joint de grain sont distingués.

  et minces (800 nm) ont été étudiées expérimentalement (Figure B.4(a-e)).

Figure B. 4 :

 4 Figure B.4: Caractéristiques courant-tension (JV) de cellules solaires standards CdSeTe/CdTe avec (a) un contact arrière en or, (b) un contact arrière en argent. L'impact d'une couche de passivation d'alumine entre la couche active et le contact arrière est illustré pour les différents types de contacts. (c) Un contact arrière en Cu 2 O+or, l'effet d'un récuit à 150°C est étudié. Comparaison des caractéristiques optoélectroniques des cellules solaires CdSeTe/CdTe standards et fines (800 nm) (d) caractéristiques courant-tension sous une illumination d'un soleil avec un contact en or, (e) Réponse spectrale correspondante (EQE). (f) EQE expérimentale d'une cellule solaire de CdSeTe/CdTe d'épaisseur 800 nm comparée à l'absorption simulée dans une cellule CdSe 0.424 Te 0.576 /CdTe d'épaisseur 500 nm/300 nm.

  Figure B.5(a)). J'ai brièvement évalué les mécanismes de résonances optiques en faisant varier les paramètres de la structure du réseau et l'épaisseur de l'absorbeur (Figure B.5(b)).

Figure B. 5 :

 5 Figure B.5: (a) Spectres d'absorption d'une cellule ultrafine avec un miroir nanostructuré à l'arrière, calculés dans la couche active de 200 nm pour une cellule de CdTe (trait continu) et une cellule de CdSeTe (trait pointillé). (b)Diagramme de dispersion montrant l'absorption totale en fonction de la période du réseau. 1-R en fonction de la période p du réseau pour une épaisseur de CdTe de 200 nm. La ligne horizontale en pointillés noirs indique la période de 400 nm qui donne lieu à la meilleure absorption sur une large bande spectrale. Le rapport largeur/période est maintenu constant (50%). Les lignes pointillées représentent les courbes de dispersion des résonances.

  2 of the available roof area in France [Le solaire photovoltaïque en France : réalité, potentiel et défis, version du 13/06/2022 2023]. -The lifetime of a solar panel is defined by the IEA as the time after which its efficiency has been reduced by 20%. Although statistical studies are scarce, it has been shown that module lifetime can exceed 35 years in temperate climates [Le solaire photovoltaïque en France : réalité, potentiel et défis, version du 13/06/2022 2023]. -In 2022 a worker group at the Fraunhofer ISE, fabricated a 19.7% efficient module using only wafers made of recycled silicon. With an efficiency slightly below the premium silicon module at 22.2%, it remains largely competitive [Fraunhofer Institute for Solar Energy 2022].

  [START_REF] Atwater | Plasmonics for improved photovoltaic devices[END_REF]

	Single-pass absorption
	Short wavelengths	ARC
	Short decay of light intensity	Absorber
	Transmission of unabsorbed photons	ARC
	Figure 2.2: Single-pass absorption scheme (reprinted from
	[Massiot et al. 2020])

Table 2 .1: Notable

 2 results summary related in the literature, on ultrathin GaAs solar cells performances. The double bar separates the experimental results from the numerical ones.

		et al. 2022]	[Micha	2022]	Abad et al.	[Camarillo		et al. 2021]	[Buencuerpo	2020]	Abad et al.	[Camarillo			et al. 2022]	[Buencuerpo				2015]	[Lee et al.		2022]	[Sayre et al.				al. 2020a]	[Eerden et	2011]	[Liu et al.	et al. 2008]	[Nakayama	Ref
		+DLARC	AlInP/GaAs /AlGaAs 100				InGaP/GaAs /AlInP +ARC 80		+ARC	AlInP/GaAs /AlGaAs 100			InGaP/GaAs /AlInP +ARC 80	+DLARC	AlGaAs/ AlGaAs contact	GaAs / p-InGaP/ patterned	AlInP+GaInP window /n-260			contact + DLARC	AlGaAs BSF / p-type GaAs	AlGaAs /n-GaAs /p-GaAs/ 200	tact	GaP BSF / p-type GaAs con-	AlInP/n-GaAs /p-GaAs / In-80			contact	GaP BSF / p-type AlGaAs	AlInP/n-GaAs /p-GaAs / In-300	tact	+ window layer/p-GaAs con-	GaAs emitter/AlGaAs BSF	GaAs sub/n-GaAs base/p-200	tact	+ window layer/p-GaAs con-	GaAs emitter/AlGaAs BSF	GaAs sub/n-GaAs base/p-200	[nm]	Stack Thick.
			27.2				25.3			25.3			16.5				26.4					22			11.3					24.8	9.42	11.9	[mA/cm 2 ]	J sc
	side	tured grating at the rear	1.152 86.4 AlGaAs/MgF 2 nanostruc-	ture at the rear side	random scattering struc-	AlGaAs/SiO 2 quasi-	mirror with SiO 2 spacer	in AlGaAs and flat back	Quasi-random structure	back mirror	Back photonic crystal, flat		mirror with SU-8 spacer	at the rear side, flat back	1.119 75.7 22.35 Quasi-random structure				layer and back mirror	0.942 78 16.2 Nanostructured TiO 2 top	nanophotonic grating	lar cells with hexagonal	1.012 79.1 9.04 Rear side transferred so-			lic mirror	lar cells with rough metal-	1.027 83.9 21.4 Rear side, transferred so-	0.662 68.8 4.3 Top side	by surface plasmons	0.760 65 5.9 Top side, strong scattering	[V] [%] [%]	V oc FF Eff Nanostructuration
			No contact layer				No contact layer	Algorithm. No contact layer	based on the Iterative Fourier Transform	Photonic crystals in the reciprocal space			No contact layer		transferred through wet etching	quasi-random photonic crystals. Pattern	Polymer blend for maskless fabrication of	printing method	diffuse) back surface reflector. Transfer	Lithography and dry etching + (specular-	on the window layer by Nanoimprint	Hexagonal nanopatterns of TiO 2 deposited		ment Talbot lithography + silver mirror	Ag/SiN grating patterned using displace-	layer	tact, though AlGaAs pillars in the silver	tact deposition. Use of localized back on-	etching (AlGaAs layer) and Ag back con-	Textured back contact by wet chemical	N 2 atmosphere	nealing e-beam evaporated Ag films in a	Ag nanoparticles were fabricated by an-	position through anodic aluminum oxide	Ag nanoparticles fabricated by masked de-	Fabrication/Simulation highlights

A 19.9% efficient ultrathin solar cell

Table 2 .3: J th

 2

	dJ grating
	[mA/cm 2 ]

Al composition n-GaAs emitter [mA/cm 2 ] dJ p-Al 0.4 Ga 0.6 As layer [mA/cm 2 ] and parasitic absorption in the p-Al 0.4 Ga 0.6 As layer and in the grating as a function of the Al content in the nanostructured grating.

Table 2 .

 2

	2 ]

4: Different rear side configurations for ultrathin GaAs cells, with theoretical J th in the main absorber, the base, the BSF and in the nanostructured layer. The total absorption is also specified.

Table 3 . 1 :

 31 1.1. Data of the semiconductor layers considered for the front and back contact losses in ultrathin GaAs solar cells. Mobility data were extracted from

	Material	Thickness Doping	Mobility	Resistivity Sheet resistance
		[nm]	[cm -3 ]	[cm 2 /(Vs)]	[ cm]	[ /x 2 ]
	n-AlInP window	25	2 × 10 18	100	0.031	12500
	n-GaAs emitter	100	1 × 10 18	2700	0.0023	230
	p-GaAs base	100	1 × 10 18	170	0.037	3670
	p-AlGaAs BSF	100	2 × 10 18	60	0.052	5200

Table 3 . 2

 32

	Solution	GaAs	AlGaAs			InGaAs	AlInP
	H 3 PO 4 :H 2 O 2 :H 2 O	∼ 100 nm/min	> 100 nm/min		etch	stop
	(3:1:40)
	H 3 PO 4 :H 2 O 2 :H 2 O	∼ 5 µm/min	> 5 µm/min			etch
	(3:4:1)
	citric acid : H 2 O	∼ 300 nm/min	Depends on the composition.	∼ 260 nm/min
	(1:5)		Selectivity higher than 100
			for GaAs over Al 0.3 Ga 0.7 As.
			1.5 nm/min for Al>80%.
	HCl	stop	etch	for	Al>60%,	∼	stop	∼ 6 µm/min
			100 nm/min for Al>80%
	NH 4 OH:H 2 O 2 :H 2 O	∼ 8-15 µm/min,	∼ 30 nm/min
	(1:5)	depends on the
		solution temper-
		ature
	NH 4 OH:H 2 O 2 :H 2 O	∼ 1 µm/min
	(2:1:10)

: Overview for the wet chemical etching of III-V semiconductors. Stop means that the solution will not etch the material (or very slowly). The etching rate is specified for some solutions.

Table 3 . 3 :

 33 .3, along with the theoretical current in the GaAs layer. Summary of the parameters used for the calculation of the absorption in a 200 nm-thick GaAs solar cell. The first row shows the parameters used for the initial structure, the second row shows the updated parameters to match the experimental results. The last column shows the theoretical current J th in the GaAs. h N is the grating height, p the period, w the a-Si:H nanopillars width, n the dielectric nanogrid refractive index and h spacer the dielectric thickness below the grating, or spacer thickness.

		DLARC	h N	p	w	n (diel)	Ti layer	h spacer	J th
			(nm)	(nm)	(nm)		before	(nm)	in GaAs
							Ag		(mA/cm 2 )
	Original	MgF 2 /Ta 2 O 5	100 750 340 1.2		180	28.2
	structure	(90/60 nm)
	Updated		110 750 350 Refractive	1 nm	280	18.25
	structure					index in
						Fig. 3.4(a)

Table 4 .

 4

		511	1.501	10.5	1.38	100
	x se = 0.1	1.465	1.454	11.5	1.31	80
	x se = 0.2	1.451	1.437	11.5	1.285	75
	x se = 0.3	1.446	1.432	14.0	1.27	72
	x se = 0.4	1.411	1.396	12.2	1.223	70

1: Peak energy values extracted from experimental CL data at room temperature. Fit parameters determined using the generalized Planck law and an absorption model at 296 K.

Table 4 .

 4 ]. Experimental energy values are reported in Table 4.2.

		Low temperature (15K)
	Sample	Peak (eV) Defect (eV)
	x se = 0	1.590	1.419
	x se = 0.1	1.543	1.332
	x se = 0.2	1.527	1.294
	x se = 0.3	1.525	1.277
	x se = 0.4	1.489	1.21

2: Peak energy and defect energy values extracted from experimental CL data at low temperature.

  ].

		1.8
			E G (Optical index at 296K)
	CdSeTe bandgap (eV)	1.4 1.5 1.6 1.7	E peak (Varshni at 0K) E peak (296K) Bowing Fit	b=0.71 b=0.73 b=0.67
		0.0 1.3	0.2	0.4	0.6	0.8	1.0
				X Se

Table 5 . 1 :

 51 Summary of notable results found in the literature on the passivation of CdTe solar cells by Al

	3 .
	2 O

Table 5 . 2 :

 52 Description of CdSeTe solar cells layer: stack and annealing temperature steps carried out by the NREL. The differences in the process steps depending on CdSeTe semiconductor layer thickness are indicated.

  .3).

	Contact description (Best cells)	J sc (EQE) (mA/cm 2 ) V oc (mV) FF (%) Eff (%)
	Au contact (C2N -CdSeTe/CdTe)	28.05	749	60	12.48
	Au contact (NREL-CdSeTe/CdTe)	31.6	770	70.4	17.1
	Ag contact (C2N -CdSeTe/CdTe)	27.6	531	35	5.14
	Ag contact (CSU -CdTe)	27.0	505	62.3	8.5
	Ag contact (CSU -CdS-CdTe)	∼ 22	∼ 820	∼ 58	∼ 10.5

Table 5 . 3

 53

: Summary of the highest efficiencies reached in standard processing at the C2N and the NREL (the corresponding IV is not shown) with a gold back contact. The performances obtained with a silver back contact at the C2N are compared to literature results with samples fabricated at Colorado State University. Silver back contact was evaporated at the C2N while it was sputtered

Table 5 .4: JVsun

 5

		1.64 × 10 -12	1	6.03 × 10 -6	1.94	0
	Au contact only, JVsun	1.64 × 10 -12	1	6.03 × 10 -6	1.89	7.27
	Al 2 O 3 + Au contact, pseudo-JV	1.71 × 10 -12	1	3.69 × 10 -6	2	0
	Al 2 O 3 + Au contact, JVsun	1.71 × 10 -12	1	3.69 × 10 -6	1.8	7.26

and pseudo JV parameters for CdSeTe cell solar cells for Au contact with and without a passivation layer in between, fitted with a 2-diodes model (J 01 & J 02 : saturation currents, n 1 & n 2 : ideality factors, and R s : series resistance).

Table 5 . 5 :

 55 Summary of the performances of CdSeTe/CdTe solar cells with Cu 2 O at the back side and Au contacts, before and after cells modification. Average values and standard deviation are calculated from 7 devices except for the last row, where only 5 are considered since 2 were shunted. No distinction as to whether or not they have a passivation layer has been made. The J sc extracted from IV curves are given as an indication only.

Table 5 . 6 :

 56 Summary of the highest efficiencies reached in standard processing at C2N with a gold and a silver back contact for a 800 nm-thick CdSeTe/CdTe sample. It is compared to a CdTe cell of the same thickness fabricated at CSU,

Table 6 . 1 :

 61 Optimized parameters deduced from the simulation, with a nanostructured rear side mirror, with Ag nanopillars embedded in a dielectric layer for a 200 nm-thick CdTe cell.

Table 6

 6

	.2

Table 6 . 2 :

 62 Optimized parameters used in the simulations for the three structures detailed in Figure6.7. J gain is the theoretical short-circuit gain between the optimized structure and a CdTe cell with a Au contact. dJ is the parasitic absorption losses occurring in the mirror and in the nanostructured grating.

  .10, the absorption spectrum in a 200 nmthick CdTe layer with an Au contact only, and with a nanostructured back side (structure C) is plotted. Several resonant wavelengths are labeled.

		1.0	A B	C D E	F G
			460	510	614	710	805
		0.8			654
	Absorption	0.4 0.6	FP FP 9 7 394 nm 464 nm 519 nm FP 6
		0.2	Structure C optimized, J th = 26.41 mA/cm 2
			Total Absorption, Structure C optimized
		300 400 500 600 700 800 900 1000 0.0 Au contact , J th = 21.23 mA/cm 2
				Wavelength (nm)

845 Figure 6.10: Calculated absorption spectra of the structure C in a 200 nm-thick CdTe layer and total absorption, compared to the absorption spectrum in a 200 nm-thick CdTe with a planar Au contact alone. J th is shown for each structure. Several resonant wavelengths are indicated and are labeled by capital letters A to G.

  Al 2 O 3 is selectively deposited by ALD. It has the same thickness as the localized gold contacts. SAM and SiO 2 are removed by reactive ion etching. A silver layer is then evaporated at the surface.While Al 2 O 3 deposition selectivity has been observed with PMMA, it is for much larger pattern size and pitch[START_REF] Färm | Selective-area atomic layer deposition using poly(methyl methacrylate) films as mask layers[END_REF]]. In the following, we present the main fabrication steps for ultrathin CdTe devices with a nanostructured rear side. Localized Au (or alternative material such as Te) back contacts are deposited at the rear side, followed by an Al 2 O 3 selective deposition by ALD. Two fabrication steps are examined. On the one hand, the method developed in the first model is used. On the other hand, we consider larger localized back contacts (in the micrometer range, the required period depends on the minority carrier diffusion length) processed with optical lithography. The alumina layer is deposited by selective ALD using PMMA (the alumina is deposited everywhere except on the PMMA layer). For the next step, depending on the model, either a a-Si:H is deposited by PECVD, or SiO 2 is spin coated on the Al 2 O 3 layer.

	Steps 1 & 2: Localized back
					contact
					CdTe
					Al O deposition 2 3
					2 3 Al O
				CdTe
				Model 3	Model 2
	Step 3: CdTe	aSi:H aSi:H deposition CdTe	SiO 2 spin coating SiO 2
					pitch
					PMMA	d
		CdTe
	Step 2: 30 nm of gold is evaporated. A large array of localized Au back contacts at the back of the CdTe is assembled monolayer (SAM) deposited by vapour phase. obtained. 5 nm of SiO 2 is sputtered, followed by a self-	SAM SiO 2 CdTe Au	SAM SiO 2 Au
	PMMA is removed.
			SAM	SAM
			SiO 2	SiO 2
	Step 3: CdTe	Au	2 3 Al O	Au
	Step 4: CdTe	Au	2 3 Al O	Au

Ag

For this model, the steps 1 & 2 are still under development. The SAM is expected to prevent the deposition of Al 2 O 3 , however no such results have been observed so far with control samples.

  ]. CdTe production cost in $/Wp for several module efficiencies (16% is the most relevant in our case) and tellurium price (from 170 to 700$/kg, note that it around 400$/kg today). Figurereprintedfrom[START_REF] Candelise | Implications for CdTe and CIGS technologies production costs of indium and tellurium scarcity[END_REF], published in 2012 which explains the cost differences with the one calculated by the NREL in 2021.

	$1.10
	$1.00
						$170
	$0.90					$200
	$0.80					$300
	$0.75					$400
	$0.70					$500
						$600
	$0.60					$700
	$0.50
	11.6%	12%	13%	14%	15%	16%
			Efficiency
	Figure 6.14:

Table A . 1 :

 A1 Peak energy values extracted from experimental CL data at room temperature. Fit parameters determined using the generalized Planck law and an absorption model at 296 K.

			1.454	1.437	1.432	1.396
	γ (meV)	10.5	11.5	11.5	14.0	12.2

† https://www.firstsolar.com/

is not yet investigated.

* https://sam.nrel.gov/

* 2-diode model fit program https://nanohub.org/resources/14300

Remerciements

//www.nrel.gov/pv/cell-efficiency.html Indices extracted from https://refractiveindex.info * PVEducation website available here

Appendix

A Additional simulation and experimental results

Chapter content

A In this Appendix, we provide additional calculation and experimental results, to illustrate and complete the approaches developed in the thesis. Each section is related to a chapter of the thesis.

A.1 Chapter 3

A.1.1 Modeling of the front contact losses

To calculate the resistive losses and optimize the front grid design, we develop a simple theoretical model based on simulations developed in the literature [START_REF] Honsberg | Photovoltaics Education Website Cited on[END_REF]; [START_REF] Liu | Optimization of grid design for solar cells[END_REF]; [START_REF] Mäckel | Analytical models for the series resistance of selective emitters in silicon solar cells including the effect of busbars[END_REF].

We demonstrate in the following the delicate balance between lateral resistivity losses and shadowing losses. The first one requires a small spacing (S) between the fingers while the other benefits from a larger spacing. Based on the grid design in Figure A.1, the relative loss related to the grid shadowing which groups the fingers and the two busbars is expressed by:

where w f and w b are the width of the fingers and the busbar respectively, S the finger spacing and L the cell length. Front contacts shading will directly impact the J sc as a higher ratio of the surface is covered by the metal. Note that only one busbar placed at the edge of the cell may be sufficient, but in case of broken fingers it still allow the current recovering. The relative loss due to resisting effect in the emitting layer is:

where R sq is the sheet resistance of the semiconducting layer, S the finger spacing and J mp , V mp the current density and the voltage at the maximum power point. We take J mp =30 mA/cm 2 and V mp =0.9 V. R sq is computed using the relation: R sq =ρ/t. ρ is the semiconductor resistivity, and t is its thickness. The resistivity can be calculated from the doping concentration (n,p) and the doping dependent mobility (µ n ,µ p ):

This equation can be simplified according to the doping type of the sample. Since the window layer is only 25 nm thick it will not contribute to the lateral conduction. Therefore, the emitter is the most important layer for electron lateral conduction to the top contacts (from the layer to the fingers). From equation A.3, we deduce that R sq would benefit from a high doping in the layer, especially since the emitter thickness is fixed at 100 nm. However, the doping of the active region must be minimized to limit the influence of Shockley-Read-Hall recombination. Therefore, the finger spacing is of utmost importance for the minimization of losses.

We add two more losses contribution linked to the metallic contacts. The first one is the relative losses due to the transport of carriers from the fingers to the busbar:

A.2.2 Se gradient characterization

In section 4.5, we focus on the chemical composition of CdSe x Te 1-x cells. We want to determine precisely the Se content in each sample. With CL characterization we managed to retrieve estimation values. To check the reliability of these results SEM-EDS was performed on each CdSe x Te 1-x sample and STEM EDS was used to characterize the CdSe 0.2 Te 0.8 device.

A.2.2.1 SEM-EDS

To collect chemical data on the samples, energy dispersive spectroscopy (EDS) measurements are operated with an acceleration voltage of 6 keV in a Zeiss Merlin VP SEM. Figure A.7 shows the Se distribution superimposed on the SEM image for each sample with nominal Se content ranging from 10% to 40%, and the associated elemental quantification. Figure A.9 presents a higher magnification HAADF image of a grain boundary within the interdiffusion area and associated elemental maps. This region is chosen to detect/highlight potential elemental segregation and associated composition changes at the GB. Indeed, in this zone, presence of higher Se content is shown with CL map acquisition. No Se segregation is visible, but similar results (absence of visible Se segregation at GB by STEM-EDS) were reported in the literature [START_REF] Fiducia | Selenium passivates grain boundaries in alloyed CdTe solar cells[END_REF]]. STEM-EDS observation of Se segregation at GB was sometimes reported in the literature [START_REF] Guo | Effect of selenium and chlorine co-passivation in polycrystalline CdSeTe devices[END_REF]]; [Munshi et al. 2018b]. However, when this phenomenon is reported, the Se concentration is at least twice as high the one from CdSe 0.2 Te 0.8 sample.

A.3 Chapter 5 A.3.1 Optical constants used for the electromagnetic simulations

Appendix

B Résumé en français

Dans son sixième rapport d'évaluation publié en 2022, le Groupe d'Experts Intergouvernemental sur l'Evolution du Climat (GIEC) a confirmé l'urgence de réduire drastiquement nos émissions de CO 2 pour lutter contre le réchauffement climatique. L'Agence Internationale de l'Energie (AIE) et le GIEC ont rappelé l'importance de la transformation de notre système énergétique (les émissions liées à l'énergie représentent trois quarts des émissions mondiales de CO 2 ) et la nécessité de remplacer les energies fossiles par des énergies à faible teneur en carbone, voire sans carbone. En 2020, les énergies renouvelables représentaient 12,6% de la consommation totale d'énergie. L'énergie solaire, l'une des ressources les plus abondantes au monde, apparaît comme une alternative essentielle. Au cours des dix dernières années, l'énergie solaire est devenue l'option la moins coûteuse pour la production d'électricité dans la plupart des pays du monde. La technologie PV réalisée avec du silicium représente 95% du marché du PV mondial, grâce à ses modules photovoltaïques peu chers dont l'efficacité dépasse 20%. Cependant, l'efficacité du silicium est limitée à moins de 30% et nécessite toujours l'utilisation d'une grande quantité de matériau avec des épaisseurs de cellules solaires de 150-200 µm et un budget thermique considérable pour la fabrication de wafer de silicium monocristallin de haute pureté. Pour augmenter l'efficacité et réduire les coûts de production l'industrie photovoltaïque est à la recherche d'alternatives. Les technologies dites couches minces (3-4 µm), basées sur des matériaux à gap direct absorbant fortement la lumière comme le CdTe, le CIGS ou encore les III-Vs représentent une solution attractive.

Les cellules solaires en GaAs détiennent le record d'efficacité des cellules simple junction (29.1%). Cependant, le coût des matières premières III-V limite l'industrialisation de cette technologie et la restreint au domaine d'application du spatial. Pour augmenter son rayonnement industriel, l'un des voies d'amélioration est la réduction de l'épaisseur de l'absorbeur. A ce jour, même les cellules solaires les plus modernes sont proches de la limite d'absorption simple passage et un piégeage efficace de la lumière pourrait réduire l'épaisseur des cellules solaires d'au moins un ordre de grandeur tout en conservant des rendements équivalents, avec des avantages considérables en termes de consommation de matériaux et de capacité de fabrication industrielle. Les développements récents dans le domaine de la nanophotonique fournissent de nouveaux outils pour aller dans cette direction. En 2019, une cellule solaire en GaAs d'une efficacité record de 19,9% avec une épaisseur d'absorbeur de seulement 205 nm a été réalisée par notre équipe en utilisant l'absorption multi-résonante à l'aide d'un miroir arrière nanostructuré [START_REF] Chen | A 19.9%-efficient ultrathin solar cell based on a 205-nm-thick GaAs absorber and a silver nanostructured back mirror[END_REF]. Les perspectives à court terme prévoient d'atteindre des rendements supérieurs à 25%.

Il serait particulièrement intéressant de transférer cette preuve de concept aux technologies commerciales comme le CdTe qui représente la majeure partie des 5% restants de la capacité de production de modules photovoltaïques en 2021 (environ 4%). La compétitivité industrielle de ces matériaux polycristallins serait quant à elle stimulée par la réduction de l'épaisseur de l'absorbeur tout en maintenant une efficacité similaire. Cela permettrait de réduire la consommation de matériaux rares tels que le tellure, de raccourcir le temps de dépôt et de faciliter la collecte des porteurs. Cependant les cellules solaires à base de CdTe restent encore bien en dessous de leur limite théorique (32%) avec des cellules solaires record d'une efficacité de 22,1%. Plusieurs défis se posent avec un rendement de conversion d'énergie sévèrement entravé à la fois Oral presentations