Keywords: CdTe grains, nanowires, passivation, footprint energy production: solar energy

CdTe thin films . . . . . . . .

Acknowledgments-Remerciements

Si vous m'aviez demandé il y a 4 ans de cela si je connaissais la ville de Marcoussis, ma réponse aurait surement été « Oui c'est là où le XV de France s'entraine ». Il faut croire que l'on y fait aussi de la recherche, car c'est là-bas qu'il y a 4 ans, cette thèse a commencé. Elle a commencé avec toi, Stéphane, et c'est donc à toi que ces premiers remerciements sont adressés, car c'est grâce à toi que ce projet a démarré. Il faut croire que malgré quelques (nombreuses ?) années de recherche derrière toi, tu restes émerveillé par de « simples » nanofils. Ta joie communicative, ton enthousiasme et ton ambition m'auront permis de passer au travers des périodes plus difficiles. J'aimerais particulièrement te remercier pour ta patience face à mes étourderies chroniques ; je te le promets ce n'était pas fait exprès. Pour me guider entre les terrains de rugby et les champs marcoussissiens, j'ai aussi pu compter sur toi, Andrea. Tu auras été d'une aide précieuse, m'appuyant de nombreuse fois avec ingéniosité, créativité et gentillesse. Au bout de la ligne DM11C se trouvait aussi une petite famille. Andrea S., Ahmed aurait pu être les pères et Julie une mère au sourire rassurant pour 4 enfants turbulents : Romaric, Amadéo, Louis et moi. Nous aurons passé autant de temps à rire qu'à râl. . . travailler ! Merci à vous 3, et surtout à tes pâtes Romaric. Malgré, tes choix rugbystique, je te dois tout de même des remerciements: un grand merci à toi, Christophe pour ta disponibilité, ta patience et ta compétence. Qualités que tu partages avec Nathalie qui m'a elle aussi beaucoup épaulé, je t'en remercie grandement. Non, Amaury, je ne serai pas ingrat. C'est grâce à toi que cette histoire à commencer, à Tokyo. Tu étais mon Sensei de l'époque, et j'ai donc une pensé particulièrement pour toi et pour toute la team Curry: Daniel, Zacharie, Maxime et Samy. Vous m'avez tous beaucoup apporté, humainement et scientifiquement. Je souhaiterais aussi remercier les personnes qui m'auront beaucoup aidé au cours de ces longues journées passées dans le noir de "la grotte" : Sophie, Fabrice et de manière plus générale tous les membres du C2N pour leur aide et leur accueil chaleureux. Un grand merci aussi à l'équipe de foot du laboratoire pour ces heures "musclées" passées ensemble et notamment à vous Abdou, François, Francesco, Stéphane et tous les membres de cette belle équipe.

Une thèse ne prend généralement fin que quand 2 personnes passent de longues heures à relire un long manuscrit, c'est pourquoi je voulais vous remercier chaleureusement. Nicolas Bareau, Laurent Lombez, j'espère ne pas avoir rendu votre tâche trop pénible. Un grand merci aussi à Julien Barjon, de s'être donné la peine de venir au laboratoire malgré la situation sanitaire, ainsi qu'à Anna Fontcuberta i Morral et Ludovic Escoubas ; pour les échanges riches et stimulants que nous avons eu lors de ma soutenance, malgré la dématérialisation de l'évènement.

Motivation, démotivation et badminton, ces mots ne riment peut être pas entre eux mais auront été un lien fort entre nous : Bérangère, Capucine, mes acolytes de bureau, je vous souhaite bon courage pour finir vos projets et vous remercie chaleureusement. Salim, merci pour ta bonne humeur et ton humour, je te souhaite aussi plein de réussite, tout comme à toi Linh. Michel, tu viens de commencer ta thèse entre de bonnes mains : bonne chance ! Ces 4 ans m'auront permis de rejoindre le cercle des docteurs de la famille Bidaud. 3(4 ?) Personnes, comme un roc immuable sans lequel ce projet n'aurait peut-être pas vu le jour, et pour cela je vous serai éternellement reconnaissant. Une thèse c'est avant tout un périple et le moins que l'on puisse dire c'est que je fus bien accompagné pendant celui-ci. Je pense évidemment à mes amis. Autour de cartes, de bières, de jeux de société ou de ballons vous avez su rendre ma vie heureuse, et pour cela je tenais tout particulièrement à vous remercier. J'ai aussi une pensée pour mes camarades de rugby avec lesquels je me suis battu sur d'autres terrains que ceux de Palaiseau et de Marcoussis. Et puis, Joseph merci pour tes coups de pouce, sans toi je n'aurais pas terminé « mon DM de Math ». J'aimerais finir par celle qui fait partie de ma vie, qui m'aura soutenu dans beaucoup de moments délicats et qui aura toujours su me rendre le sourire, oui, je parle de toi Elsa.

General introduction

Global warming has been shown to be a threat for the planet, animals and people. In 2015, 195 countries agreed on making efforts to mitigate the increase of the temperature down to 1.5°C by 2100. But in 2017, still a large part (≈ 80 % [START_REF]renewables 2019[END_REF]) of the energy consumption is non electric and/or originates from fossil sources of energy. Primary energy consumption increased in 2018 by 2.8 %, in terms of million tonnes oil equivalent. In order to meet the objectives settled by the COP21, the switch towards renewable energy sources is necessary. Among them, solar photovoltaics (PV) generates 2.4 % of the global electricity production and benefits from a quasi unlimited source of energy. PV energy production becomes more and more competitive as its levelized cost of energy production (LCOE) decreases fastly, making it the fastest growing energy production in terms of capacity of production. Competitiveness of a PV technology are driven by their power conversion efficiency increase and cost decrease.

More than 95 % of the PV solar cells and modules are based on silicon (Si). Indeed, the power conversion efficiency of Si modules has reached 20 % and their production cost is continuously decreasing. Alternative PV technologies are investigated in order to overcome the theoretical limit of ≈ 30% for Si single junction solar cells.

A first solution which is explored is the development of thinner solar cells. Regardless of the materials consumption, thin-films technologies with the advantage of low production costs and a high manufacturability. Thin-film solar cell technologies are based on direct bandgap materials (CdTe, CIGS, perovskite or a-Si) with high absorptivity. Because of a low growth temperature, these materials are grown in polycrystalline phase. These technologies are more and more competitive with Si based solar cells but are still far from their theoretical efficiencies. Indeed, the two main thin-film PV materials, CdTe and CIGS, record solar cells have reached respectively [START_REF] Brandt | Determining interface properties limiting open-circuit voltage in heterojunction solar cells[END_REF].1% and 23.4% with a theoretical limit above 30 %. The role of grain boundaries, doping levels and minorty carriers lifetimes in the actual performances is still under debate in the community and needs to be elucidate to make thin-films great again.

The combination of high-efficiency solar cells in tandem configuration on a Si platform has recently attracted great attention using thin film III-V materials. Indeed, coupling a Si solar cell with another material allow to increase the theoretical limit of the tandem cell up to 38 % under one sun [64]. III-V materials are good candidate in order to be combined with Si. First, III-V materials have always been at the forefront of high efficiency PV as GaAs single junction has reached 29.1 % efficiency. Another advantage of III-V materials is that they allow for a large bandgap tunability by controlling the alloy. But the direct growth of a high quality III-V materials on a Si substrate is made complex because of the difference in lattice constant. Direct growth of III-V on Si leads to a highly defective top layer. Indeed, the growth of GaAsP on Si only reached 25.9 % efficiency. The growth of III-V in the form of nanowires (NW) lets the strain, induced by the lattice mismatched, to be relaxed [133]. Direct growth in nanosctructures would allow for the crystalline quality to be greatly improved. Besides, the periodic arrangement of NWs in an array should allow for efficient light collection management [START_REF] Anuj | Toward Optimized Light Utilization in Nanowire Arrays Using Scalable Nanosphere Lithography and Selected Area Growth[END_REF]. Proof of concepts have been developed and reached efficiencies up to 17.8 % of PCE.

For both polycrystalline thin-films and NW-based solar cells, efficiencies are still hindered by the materials quality. As these materials are based on nanoscale patterns, the use of a nanoscale spatial resolution characterization technique is fundamental. By coupling the high spatial resolution offered by an electron beam source and the large scope of information which can be unveiled through the luminescence spectrum, cathodoluminescence offers a great opportunity to probe optical and transport properties in nanomaterials. In this thesis, we aimed at developing novel methods in order to assess physical quantities in nanomaterials such as CdTe polycrystalline thin films and GaAs nanowires.

• In Chapter 1, I introduce the context of the thesis by a brief description of the energy production and the photovoltaic landscape. Then the basics the photovoltaic effect and the different phenomena limiting the efficiencies are presented. Subsequently, further explanations on the polycrystalline thin-films and nanowires based solar cells technologies and issues are given. I finish this Chapter with a brief overview of some characterization techniques which are commonly used in the PV community in order to emphasize the aim of my thesis.

• In Chapter 2, The C2N CL setup used during my thesis is described. First, I focus on the C2N setup and main features, such as the spatial, temporal and spectral resolution. Then we describe the process for the data analysis of CL mapping and give insights concerning the spectral corrections. Perspectives, motivations and preliminary results for an absolute calibration of the photon flux is developed.

• In Chapter 3, I present the characterization of CdTe polycrystalline thin-films. First, a state of the art of the CdTe polycrystalline thin films and the open questions are introduced. Then I present a study of impact of the CdCl 2 annealing temperature on the formation and the passivation of defects using room-and low-temperature CL mapping on the exact same microscopic areas. A correlation between statistical and local analyses at RT and the defects distribution at LT has been found to explain the observed annealing optimal temperature observed. Then CdSeTe/CdTe structures are studied, the positive impact of the Se on the radiative efficiency is shown with a series of samples with various Se concentrations.

• Chapter 1.3.3 is dedicated to the study of GaAs NWs. In the first part I detail the growth process and the processes which limit the V oc in these structures. I investigate the presence GaAs crystalline phases and their optical properties in single NWs using CL polarimetry at low temperature. Besides, lifetimes measurements at different locations in a single NW have been performed for carrier dynamics analysis. Then a contactless method for the doping assessment in single nanowires is presented and applied to the evaluation of Te incorporation in GaAs NW. I will end by presenting results on the surface passivation. Different high-bandgap passivation shell will be compared. Time-resolved experiments are combined with a simplified model to bound the surface recombination velocity. Besides, we will present the impact of a lattice mismatch passivation shell on the optical properties of a GaAs NW.

1.1 Outlook of the energy and the photovoltaic markets

The energy consumption and its consequences

Since 1950 the total population have been multiplied by a factor four, causing an increase of the impact of the human activities (energy consumption, manufacturing output etc..) which can be directly linked to the emission of CO 2 , known to increase greenhouse effect. Indeed, since the pre-industrial period, the human activities have caused a climate global warming of 1 °C [START_REF] Masson-Delmotte | Summary for Policymakers[END_REF] (compared to . But global warming is expected to continue over the years [START_REF] Masson-Delmotte | Summary for Policymakers[END_REF], reaching 2 °C in 2060 if no reduction of the CO 2 emission is made (see Figure 1.1). [START_REF] Masson-Delmotte | Summary for Policymakers[END_REF] Global warming represents a danger for many reasons. First it is very likely to impact ecosystems because of an increase of the mean temperature of oceans and their acidity causing dramatic damages on corals, mangroves or arctic region. Depending on the faculty of human to modify their habitat, the global warming may also have dramatic impact for the coastal population or even human health. So, limiting the global warming is mandatory in order to limit the impact of the human kind on earth. To this aim, the Paris agreement is a commitment between 195 countries (over 197 in the United Nations) which agreed on to limiting the warming to 1.5 °C by 2100.

Four model pathways to limit the increase of the temperature to 1.5 °C estimate the need to decrease global anthropological emission of CO 2 down to 45 % by 2030 and net zero by 2050 [START_REF] Masson-Delmotte | Summary for Policymakers[END_REF]. 25 % of the production of greenhouse gas has been estimated to be due to the electricity production [START_REF]Changements climatiques 2014 -rapport de synthèse[END_REF]. But no reduction in the final demand is expected. The switch towards renewable electricity production is mandatory as much as the decrease of the proportion of the primary energy consumption (coal, oil, gas...). [START_REF] Masson-Delmotte | Summary for Policymakers[END_REF] To meet the Paris agreement objectives to limit the global warming to 1.5°C by 2100 efforts should be made to reduce the CO 2 emission. Pathways to mitigate global warming rely on the suppression of the primary energy consumption and the switch towards renewable energy production.

Low CO 2 footprint energy production: solar energy

In 2018, the global energy demand has increased by 2.3 % mainly due to an economic growth and an increasing demand for the heating and cooling systems. This has led to an increase in the global energy related CO 2 emissions by 1.7 %. In 2017, renewable energies accounted for 18.1 % (traditionnal biomass and modern renewables) of the final energy consumption (see Figure 1.3). The share due to the modern renewables to the total final energy consumption has grown by 4.4 % compared to 2016 [START_REF]renewables 2019[END_REF]. But the global electrification of heating and transport may provide an opportunity for renewable energies to further expand. Indeed, in the power sector renewables are increasingly preferred for new electricity generation. [START_REF]renewables 2019[END_REF] In 2018, 181 GW of renewable power production has been installed, setting a new record.

Since 2015, most of the additions of the renewable power capacity is represented by solar photovoltaic (PV), which reached over nearly additional 100 GW for two consecutive years compared to about 50 GW for wind power energy production. [START_REF]renewables 2019[END_REF] The increasing energy production through solar PV may be explained by many factors. The decrease of the module price, making the solar energy more and more competitive with primary energy sources (see Figure 1.5), and the energy payback time of a PV system which is about 2.5 years in northern Europe and 1.5 years in the southern Europe [START_REF]Photovoltaic report[END_REF]. According to several scenarios, by the year 2050, PV electricity should cover more than a 20 % share of the global primary energy demands thanks to the scale of the solar irradiance, its predictability and its ubiquitous nature [START_REF] Haegel | Terawatt-scale photovoltaics: Transform global energy[END_REF]. Still, because of its inability to produce energy over the night or over certain days (cloudy days), solar PV based energy source may impact greatly our energy consumption and use. [START_REF]Photovoltaic report[END_REF] The energy demand and production gradually increase over the years. A global electrification of our way of energy consumption occurs. Renewables will play an important role in order both to meet the demand but also environmental considerations. Because of its cost competitivity and scalability solar photovoltaic (PV) is expected to play a significant role.

The photovoltaic market

PV represents a very promising source of clean energy. From 2017 to 2018, the PV energy production has been increased by 29 % reaching a total of 435 TWh of energy production in the World [START_REF] Bp | BP Statistical Review of World Energy[END_REF]. Different technologies exist for the fabrication of PV modules. The main technology is based on Silicon. It represents 95 % of the total production in 2017. Other materials can also be used instead of Si in order to conceive PV modules, such as CdTe, CuInGaS (CIGS), GaAs for instance. Figure 1.6 shows the market share of the energy production between the different technologies. Within thin-films, CdTe (50 %) and CIGS (41 %) are the main technologies in terms of annual global thin-films PV module production [START_REF]Photovoltaic report[END_REF]. [START_REF]Photovoltaic report[END_REF] Power conversion efficiency of solar cells depends on the type of technology. The record solar cell efficiency achieved a 47.1 % power conversion efficiency under concentration illumination [67]. Mono-Si and multi-Si have reached 26.7 % and 22.3 % efficiencies respectively. Depending on the application, a technology is chosen over an other. For instance, for military or space application high efficiency solar would be more relevant despite their high cost. Thin-film PV is also especially suited to meet particular product needs, such as flexibility or nuild-in integration. Still, Si is the main technology thanks to its cost and relatively high power conversion efficiency.

Conclusion of the section

We have seen that there is a growing demand for energy. Renewable energy production increases in order to meet both the demand and the goal of lowering the global warming. Among renewable sources, photovoltaics represents an interesting source of energy. PV is composed of very different technologies with various costs, efficiencies and aspects. Such technologies vary from their cost and efficiencies and so their scope of applicability.

Physics of solar cells 1.2.1 Solar spectrum and photovoltaic effect

Solar irradiance on earth

The power surface density emitted by the sun, at the distance of the Earth, is 1.3kW.m -2 at the Earth's surface before entering into the atmosphere. Assuming that 70% of the incoming light impinges the ground, the power surface density is about 1kWh.m -2 . For instance, the energy per day that can be absorbed from solar irradiation is between 2.9 kWhr.m -2 and 5.8 kWh.m -2 (at Stuttgart) [152]. This amount of energy is largely sufficient to produce enough energy to meet the needs for energy for the entire world. [104] The solar spectrum can also been first approximated by a blackbody radiation at T=5800K. The standardized solar spectrum at the Earth's surface (at a latitude of 41.8 °), is plotted in Figure 1.7. The presence of molecules (O 2 , CO 2 , O 3 or H 2 O) in the atmosphere causes absorption and scattering, and induces absorption and dips in the solar spectrum.

Basics of the photovoltaic effect

A typical scheme of a solar cell is drawn in Figure 1.8 a). The light emitted by the sun is absorbed in the active region (red zone). The absorbed light generates carriers which can then flow to the metallic contact (black bands) and be collected by an external electric circuit.

Most of the solar cells are composed of semiconductor (SC) materials. A SC materials is defined as a materials with a small difference in energy between the last occupied energy band, named valence band (VB), and the first unoccupied band energy, the conduction band (CB). The difference in energy between the top of the VB and the bottom of the CB is the bandgap (E g ). In Figure 1.8 b) is plotted a band diagram of a simplified solar cell. In the case of a SC material, the photovoltaic effect is composed of the different following steps :

1. Light absorption: If the energy of the incoming photons is lower than the bandgap (hν<E g ), the light is transmitted. If the incoming photons have an energy higher than the bandgap (hν>E g ), an electron of the VB is excited into the CB creating an electronhole pair. This step corresponds to step A in Figure 1.8. Carriers generated within the SC quickly loose their kinetic energy and reach the bottom of the CB and the top of the VB via a process named thermalization (step B).

2. Carriers transport: Once generated, carriers can move freely in the SC in a random direction if no external electric field is applied. The collection of light-generated carriers by the internal electric field induced by the p-n junction causes a movement of electrons to the n-type side and holes to the p-type side of the junction.

3. Carriers collection: This step is described as the step C in Figure 1.8, where carriers flow to the metal contacts described as grey rectangles. The current is collected by metallic contacts and can power an external load.

Solar cell efficiency definition

A p-n junction presents a current-voltage dependence (J dark ) which can be characterized by a diode-like behavior (Equation 1.1), if no resistance effects are taken into account:

J dark (V ) = J 0 × exp( qV nkT ) -1 (1.1)
Where q is the elementary charge, k is the Boltzmann constant and V is the external voltage applied to the p-n junction. The factor n is the ideality factor of the diode. The current dependence with the voltage is characterized by the ratio qV /kT where kT /q corresponds to a thermal voltage. The pre-factor J 0 is the saturation current which can be related to the flux of minority carriers generated with the thermal excitation (kT ) under equilibrium.

When the p-n junction is under illumination, the photovoltaic effect (described in Subsubsection 1.2.1.2) generates a photogenerated current J ph . The photogenerated current is added to the J dark and is assumed independent from V. The equivalent circuit of an ideal solar cell is presented in Figure 1.9 a), where the p-n junction is represented by a diode and the photogenerated current by a current generator. Characteristic J-V curves for a diode in dark (blue curve) or under illumination conditions (red curve) are plotted in Figure 1.9 b). The open circuit voltage (V oc ) corresponds to the voltage when the cell (under illumination) is not connected to an external circuit (J(V oc )=0 A). The short circuit current (J sc ) corresponds to the current when no external voltage is applied (J(0)=J sc ).

J(V ) = J dark (V ) -J ph (1.2)
Figure 1.9: a) One-diode equivalent electric circuit for a p-n junction under illumination (Equation 1.2).b) Characteristic J-V curves for ideal solar cell under illumination (red) and in the dark (blue), corresponding respectively to Equation 1.1 and Equation 1.2.

The maximum power P = J × V is obtained at the maximum power (MP) point, plotted in Figure 1.9. Efficiency η of a solar cell can then be defined as the ratio between the incident solar irradiation per surface unit (Φ sun =1 kW.m -2 ) and the MP. η is related to V oc and J sc (Equation 1.3) by defining the Fill Factor (FF). The FF is defined as the ratio between the area of the rectangle represented by V oc J sc and V mp J mp .

η = V oc × J sc × F F Φ sun (1.3)
The V oc is directly linked to the separation of quasi-Fermi levels within the SC under illumination, and so depends on the quantity of photogenerated carriers. On the other hand, J sc is directly linked to light absorption efficiency and to the quantity of carriers collected. These two parameters represent the two main figure of merits for a PV materials.

Limitations of solar cells effiencies

The Shockley-Queisser limit

Shockley and Queisser (S.-Q.) have developed a model in order to estimate the upper limit, the detailed balance limit, of the solar cell efficiency as a function of the bandgap based on thermodynamical considerations [START_REF] Green | Energy conversion approaches and materials for high-efficiency photovoltaics[END_REF][START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF][START_REF] Shockley | Detailed Balance Limit of Efficiency of p-n Junction Solar Cells[END_REF].

In the S.-Q. limit, the different stages needed for the photovoltaic effect is described by five assumptions. The first one is that each incident photon with an energy higher than E g (hν > E g ) is assumed to generate an electron-hole pair. The SC is considered as transparent for photons with energies hν < E g . This photogeneration contributes to the J sc . Then the electron-pair is assumed to lose the excess energy above E g . Once electron-hole pairs have been generated, they may either recombine radiatively or be collected at the selective contacts. Only a part of the photons can be emitted to the ambient as some photons are reabsorbed. Emission of photons from the SC can be modeled by the generalized Planck law (further explanations are given in Subsection 2.4.1). Collection of photogenerated carriers implies two different contacts. The last assumption is that these two contacts are ideal and exchange only one type of carriers.

Figure 1.10: Shockley-Queisser limit under AM1.5G (blue curve) and under full concentration (dashed green curve). Best solar cell efficiencies (multi color diamonds) are compared with the Shockley-Queisser limit. Figure reprinted from [START_REF] Green | Energy conversion approaches and materials for high-efficiency photovoltaics[END_REF].

On the basis of the assumptions described above, efficiencies can be determined as a function of the bandgap. The S.-Q. limit (blue curve) is plotted as a function of the bandgap in Figure 1. [START_REF] Barnard | 3D Lifetime Tomography Reveals How CdCl 2 Improves Recombination Throughout CdTe Solar Cells[END_REF]. By comparison with best solar cells efficiency, plotted in Figure 1.10, we can realize that there is still room for improvements for all PV technologies.

Sources of loss in a solar cell

The theoretical model explained above, the S.-Q. limit, leads to J sc SQ and V oc SQ which only depends on E g . In the following we will discuss the different mechanisms which limit the "real" solar cells efficiencies. J-V curves where the losses mechanisms are successively taken into account are plotted in Figure 1.11.

In fact, no absorber with a step function absorption exists. The finite thickness of the cell can be given as one of the explanation. Besides, a fraction of light can be absorbed in layers which are not the absorber and does not lead to the generation of electron-hole pairs. These two effects decrease the Jsc value from J sc SQ to J sc rad . The current-voltage characteristic defined for the pn-junction in the previous section is based on the assumption that only radiative recombination takes place. This ideal situation enabled us to determine an upper limit for the open-circuit voltage and the efficiency of a solar cell [178]. However, in real solar cells recombination via bulk or surface defects predominates. Photogenerated carriers can indeed be lost due to non-radiative recombinations reducing the cell voltage to the actual V oc . Resistive losses must also be taken into account. They hardly affect the V oc or the J sc values but mainly the FF. 4) radiative losses with non ideal absorption [START_REF] Alberi | Measuring long-range carrier diffusion across multiple grains in polycrystalline semiconductors by photoluminescence imaging[END_REF] non radiative losses are taken into account. For each J-V curve, a rectangle represents the maximum power achievable. The final white square correspond to the model of an actual solar cell taking into account radiative and non-radiatives losses, non ideal absorption, thermalization processes and resistive losses. Figure reprinted from: [START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF] In order to take into account all the effects depicted previously, a more complex equivalent electric circuit can be determined (Figure 1.12). To include non-radiative recombination, a second diode with an ideality factor n 2 is added to the model. The two-diode model considers both radiative and non-radiative recombination and reproduces the J-V curves of real solar cells relatively well [178]. Besides, in order to be more realistic two resistances corresponding respectively to the shunt resistance (R p ) and the series resistance (R s ) are added to model the non ideal behavior of the contacts. The Shockley-Queisser limit gives a theoretical limit for the solar cell efficiency. But, most of the record solar cell are still far from their radiative limit. This effect can be explained by different loss processes (resistive losses, parasitic and non-ideal absorption, non radiative recombination).

J(V ) = J 01 exp( q(V -J × R s ) n 1 kT ) -1 + J 02 exp( q(V -J × R s ) n 2 kT ) -1 + V -J × R s R p -J ph (1.4)

Recombination processes in a semiconductor

In the following we will describe processes in which electrons and holes, which have been generated through the absorption process, are annihilated in reactions named recombination. These reactions can be described by radiative and non radiative recombination that can occur within a SC.

A radiative recombination is the recombination in which an electron from the conduction band directly recombine with a hole in the valence band and releases a photon. The recombination between an electron in a state E i with a hole in a state E f results of the emission of a photon with an energy hν = E i -E f [179]. A set of transitions have been presented in Figure 1.13. A band-to-band recombination occurs when an electron of the CB recombines with a hole of the VB (Reaction (1)). The resulting photon energy is hν≈ E g and is therefore weakly absorbed. This type of recombination mostly occurs in direct bandgap SC (GaAs, CdTe, CIGS, Perovskite) and is very limited in indirect bandgap (Si) as the momentum needs to be preserved.

e + h → γ (1.5)
The total radiative recombination rate (R rad ) is directly related to the product of the quantity of electrons (n) and holes (p). For a given transition the rate of radiative recombination R rad can be described by Equation 1.6:

R rad = B × n × p (1.6) R net,rad = B × (np -n 2 i ) (1.7)
where B is a coefficient characteristic of a materials which can be calculated from absorption coefficient. In equilibrium with the ambient radiation at 300K, both the generation and the recombination are equal (detailed balance principle). It follows that the net radiative recombination rate (R net,rad ) is balanced with an opposite generation rate G 0 =Bn 2 i (Equation 1.7).

If we consider an intrinsic SC (no intentional doping), with an injection level (∆n,∆p) and that ∆n=∆p, we can emphasize that the band-to-band recombination rate increases with ∆n 2 :

R rad,BB = B BB (n 0 + ∆n)(p 0 + ∆p) (1.8)
Transitions ( 3) and (4) in Figure 1.13 result from the transition of an electon/hole in a localized state of a defect (donors/acceptors) within the bandgap. Donors/acceptors can be formed by many ways: substitution by an element with a different electronic structure (more or less electrons than the element they replace), vacancies or interstitials. The predominant presence of one type of impurities in a SC (donors or acceptors) can induce an increase of the electron/hole concentration which will define as n-type (majority carriers: electrons) or p-type (majority carriers: holes).

If we consider a p-type doped materials with an impurity concentration N a , such as p 0 N a , with a low injection level (N a ∆n), the defect-related recombination rate can be written as in Equation 1.9, taking into account that N a n 0 ≈ n 2 i . For the defects related transitions, the radiative recombination rates depend linearly on the injection level: R rad,def ects ≈ B def ects N a ∆n (1.9)

The transition that starts and finishes in defect states (transition (5)) can be described through the Donor-Acceptor recombination model (DAP). This type of recombination occurs when an overlap between an acceptor and a donor defect state exists. The recombination rate and energy is directly related to the density of electrons/holes occupying the acceptors/donors states. The photon energy resulting from the DAP transition is given by Equation 1.10:

ω DAP = E g -E a -E d + e 2 4π 0 SC R DA (1.10)
where E d and E a are the binding energies of holes (E a ) and electrons (E d ), respectively. These energies are determined by considering that the electron/hole is bound by coulombic interaction to a positively charged/negatively center (donnor/acceptor). The last term corresponds to the coulombic interaction between the two centers separated by a distance R DA , screened by the matrix with a permitivity 0 SC . Two cases can be distinguished, widely separated (R DA →∞) and associated DAPs. In the case of widely separated DAP, the coulombic term vanishes and the photon energy is only given by the difference between the bandgap and the binding energy. When the average separation between the donnor and acceptor is close enough, the two centers interact. This repulsion result in a blue-shift of the energy of the emitted photon.

During a recombination the released energy needs to be captured by other particles which, in the case of non-radiative recombination, are electrons or holes (Auger recombination) or phonons (non-radiative recombination). During Auger recombination, the energy set free when the recombination happens is transferred to an electron or a hole as kinetic energy. This kinetic energy is then lost through collisions with phonons (thermalization). This type of recombination may dominate in indirect bandgap materials, such as Si.

Recombinations through defects, also called Shockley-Read-hall recombination, occur when an electron is trapped by an energy state in the bandgap (E imp ) and a hole moves up to the same energy level and recombines. Such energy states are introduced though defects in the lattice. The rate at which electrons (or holes) are captured is described by Equation 1.11 (or Equation 1.12) [178]: R e,imp = σ e ν e n e n h,imp (1.11) (1.12) where σ e (σ h ) is the capture cross section for the electron, ν e (ν h ) the thermal velocity of electrons (or holes) and n e is the electron density and n h,imp is the hole density of impurities occupied by a hole. Then the density n h,imp and n e,imp is given by the concentration of impurities n imp and by the equilibrium Fermi distribution [START_REF] Green | Solar cells: operating principles, technology and system applications[END_REF]. By introducing all these equations in the steady state continuity equations for the different carriers concentration n e n h and n imp (occupied impurities) the non radiative recombination rate R non,rad (Equation 1.13) is given by:

R h,imp = σ h ν h n h n e,imp
R non,rad = np -n 2 i τ h0 (n + n 1 ) + τ e0 (p + p 1 ) (1.13)
τ e0 =1/(n imp ν e σ e ) and similarly τ h0 are the mean lifetimes of an electrons and a hole when all impurities are occupied by the other type of carriers. This defines the smallest lifetime for an electron and a hole. Using Equation 1.13, for equal τ e0 and τ h0 , the maximum value for R non,rad is reached when the energy level E i lies in the middle of the bandgap. In other words, the defects is more damaging if the level is in the middle of the bandgap.

Lifetime and diffusion length

When carriers are generated through light absorption or other mechanism and that the number of minority carriers is increased above that at equilibrium, the excess minority carrier concentration will decay through recombination processes. The continuity equation for the additionnal electron concentration can be written as in Equation 1.14 [START_REF] Würfel | Physics of solar cells: from basic principles to advanced concepts. 2nd[END_REF]:

∂(∆n) ∂t = ∆G e -∆R e (1. 14
)
Where ∆n is the excess minority carrier concentration, ∆G e is the generation due to excitation and ∆R e is the additional recombination rate for excess minority carriers. Then when the excitation is turned off (∆G e = 0) Equation 1.14 can be written as using Equation 1.6:

∂(∆n) ∂t = -∆R e = -Bp 0 ∆n (1.15)
If we assume a p-type doping, Equation 1.15 can be solved in order to determine the evolution of the minority carrier density after the excitation has been turned off:

∆n(t) = ∆n(0) exp -t τ (1.16)
The characteristic time is defined as τ =1/Bp 0 . This lifetime has been defined taking into account all the recombination processes. This mean lifetime describes the time until a carrier recombined. Using this definition, the recombination rate for one type of recombination can be written as:

R = ∆n τ (1.17)
Different recombination processes can occur in parallel: radiative and non-radiative recombination. These processes can be described by R rad = ∆n/τ r (radiative recombination rate) and R nonrad = ∆n/τ nr (radiative recombination rate). Based on electrical analogy, recombination processes can be seen as current, so as for electric circuit, The recombination current are summed leading to Equation 1.18:

∆n τ = ∆n τ nr + ∆n τ r (1.18) 1 τ = 1 τ nr + 1 τ r (1.19)
Then an average diffusion length, which is the average length a carriers moves between generation and recombination, can be related to the carrier lifetimes following the Equation 1.20, where D is the diffusion coefficient [START_REF] Mathieu | Physique des semiconducteurs et des composants électroniques: cours et exercices corrigés[END_REF]: Because of dangling bounds or presence of oxide, the surface of a crystal may contain a higher density of defects states than the bulk materials. Thus, effective properties of a material can be modified by the surface properties. Indeed, the measured lifetime of a materials can be modified by the surface recombinations (τ surf ace ). The effective lifetime of a materials is thus defined as in Equation 1.21:

L = √ Dτ (1.20)

Surface recombinations

1 τ = 1 τ bulk + 1 τ surf ace = 1 τ nr + 1 τ r + 1 τ surface (1.21)
The recombination rate for the surface states can be characterized by a surface recombination velocity S surf ace . Then the surface recombination rate is given by Equation 1.22, where ∆n(x = x s ) is the excess minority carrier at the surface:

R surf ace = S surf ace × ∆n(x = x s ) (1.22)
In this subsection we explored the different recombinations processes and factors that influence the recombination rates. Recombination processes are one of the pathway to explain the difference between the Shockley-Queisser limit and the existing experimental efficiencies.

PV technologies for low cost production of energy

In this section I will focus on dominant PV technologies which are both mono-crystalline or multicrystalline Si, thin-films technologies and I will also discuss of Tandem solar cells as a perspective for the future. I will first describe briefly the dominant Si technology in subsection 1.3.1. Then I will describe thin-film technology based on different materials such as the CdTe, CIGS or the emerging perovskite. Then I will describe tandem III-V solar cells and the opportunity represented by nanowire-based solar cells.

Silicon solar cell technology

Si PV modules are the most commercialized technology. Silicon has a nearly optimum bandgap for sunlight absorption and silicon solar cells reach a high photoconversion efficiency due to the good material quality and widespread technological know-how [START_REF] Yoshikawa | Silicon heterojunction solar cell with interdigitated back contacts for a photoconversion efficiency over 26%[END_REF]. Si technology mostly benefits from a very abundant raw material, a low cost of cells and modules production. Latest improvements in the efficiency have been made thanks to two technological developments: interdigitated back contact (IBC) [START_REF] Smith | Silicon solar cells with total area efficiency above 25 %[END_REF] and the use of an heterojunction for contact passivation (HJ) [START_REF] Masuko | Achievement of More Than 25% Conversion Efficiency With Crystalline Silicon Heterojunction Solar Cell[END_REF]. Yoshikawa et al. [START_REF] Yoshikawa | Silicon heterojunction solar cell with interdigitated back contacts for a photoconversion efficiency over 26%[END_REF] has shown that coupling these two technologies allows to reach a 26.3 % efficiency in laboratory on a 180 cm 2 (Figure 1.15). The maximum efficiency considering the radiative limit for Si is determined to be at 3 %. However, taking into account Auger recombinations, Richter et al. have demonstrated a new theoretical limit for Si cells at 29.4 % [START_REF] Richter | Reassessment of the Limiting Efficiency for Crystalline Silicon Solar Cells[END_REF].

Many reasons, such as Si abundance, technology maturity or its stability, have pushed Si modules to be the dominant PV technology. However, despite its nearly optimal bandgap for sunlight absorption, Si PCE suffers from its indirect bandgap and Auger recombinations, limiting its efficiency below 30 %.

Polycrystalline thin-film technologies

Polycrystalline thin-films (pc-TF) technologies represent an attractive alternative. It is especially well suited for built-in PV or other integrated PV as it can be grown on a large surface of glass or a flexible substrate. In this subsection we will discuss on the polycrystalline thin-films technology as an opportunity to produce competitive modules. We will describe the different materials, their deposition techniques, the stacks used for thin-films and finish by discussing solar cells by discussing their limitations. Polycrystalline thin-films technology, as Cu(In,Ga)Se (CIGS), CdTe or perovskite represent 4 % of the global PV shipment in 2017. The high potential for the polycrystalline thin-film PV materials relies on low active material consumption and costs [START_REF] Powalla | Thin-Film Solar Cells Based on the Polycrystalline Compound Semiconductors CIS and CdTe[END_REF]. Unlike thick Si wafers (> 150µm), thin-film PV cells are based on the growth of thin absorbers of a few microns. Furthermore, active materials are deposited using low pressures and temperatures making it compatible with large scale industrial process. Unlike Si, they are based on direct bandgap SCs which enable complete absorption in a few microns absorber thickness. Reprinted from [151] Figure 1.16 presents usual polycrystalline solar cell architecture. The p-type absorber is combined with a high bandgap n-type window layer (CdS, MZO or TCO) in order to form a p-n heterojunction. pc-TF are grown both in the superstrate (CdTe, perovskite) or substrate (CIGS) configuration (see Figure 1.16) :

• CdTe: TCO is sputtered at room temperature on SnO 2 :F -coated glass to form the underlying substrate [7]. Then CdTe is grown with typical growth temperatures under 600 °C using a closed-space sublimation technique (CSS). Classic CdCl 2 post-deposition annealing is then performed to improve materials quality. Metal contact layers (Cu,Au) are then deposited to form the back contact. By adding Se in the absorber structure, record cells efficiency have just reached 22.1 % [67]. Main advantages of CdTe are the cheap deposition technique and the high stability. About 20 GW of this solar cell technology has been installed, with a module PCE averaging at 17-18 %. But the main drawbacks are the scarcity of Te, and still low efficiency compared to Si.

• CIGS : CIGS can be grown using co-evaporation or sequential process (where precursors are first deposited and then annealed). p-type CIGS is grown directly on the back contact which is classically Mo, but other back contact are currently explored [START_REF] Gouillart | Reflective Back Contacts for Ultrathin Cu(In,Ga)Se 2 -Based Solar Cells[END_REF]. Then a CdS buffer is deposited to create the p-n heterojunction. Alkali treatment (K,Na) are used to improve the materials properties like to passivation and the doping [2]. CIGS is a quaternary alloy which allows managing the bandgap (bandgap grading) in the structure. The highest efficiency achieved in laboratory has recently reached 23.35 %.

• Perovskite: Pervoskite is the name of a crystalline structure with a composition ABX 3 . The deposition technique for the perovskite is performed under low temperature and ambient pressure, making it very cheap. An other advantage is the tunabality of the bandgap by only changing the composition, making it compatible with tandem with Si or even tandem perovskite-perovskite. Nearly all of the cells are prepared in a superstrate configuration. As shown in Figure 1.17, except for pervoskite, pc-TF modules still lag behind Si in terms of power conversion efficiencies. Despite theoretical (S.-Q. limit) pc-TF PCE efficiencies over 30%, highest cells efficiency are still below 25%. These limited performances can be explained by the materials quality. Indeed, efficiencies are dragged by low minority carriers lifetimes and low diffusion lengths because of high bulk defect densities, compared to Si or III-V materials. Besides, cells suffer from high interfaces (back and front) recombination velocities and low doping levels. Because of their low deposition pressures and temperatures, these materials are grown under polycrystalline structures. The role played by the grain boundaries [2,[START_REF] Visoly-Fisher | How Polycrystalline Devices Can Outperform Single-Crystal Ones: Thin Film CdTe/CdS Solar Cells[END_REF] is still under debate and may hinder the performances [6,50,95] of polycrystalline thin-films. Polycrystalline thin-films present an interesting competitiveness with Si thanks to a good combination of low production cost and good modules efficiencies. Still Si remains the dominant PV technology because polycrystalline efficiencies are still lower than Si. These low efficiencies are mainly explained by poor V oc values. Chapter 3 explores thin-films characterization of CdTe.

Tandem III-V on Si: nanowires-based solar cells

The principle of a tandem III-V/Si solar cell

As we have seen previously in Subsection 1.1.3, Si (mono or poly) is the main PV technology due to its low cost and relatively high module efficiency. In this subsection we will describe the opportunity to reach very high efficiency solar cells by coupling two absorber layers. We will first describe the tandem III-V/Si solar cells principles and their theoretical optimization. Then we will explain why and how nanowire-based solar cells offer the opportunity to grow directly high quality III-V materials on Si.

Tandem cell is based on the use of a Si bottom cell (E g = 1.1eV) with a top cell with a higher bandgap. The maximum efficiency of a tandem cell as a function of the top and the bottom cell bandgap is plotted in Figure 1.18. The maximum efficiency which can be reached for a tandem cell with Si, with a 1.78eV top cell, is 38.1 % efficiency [100][64]. The white circle presents the maximum efficiency which can be achieved with a bottom cell Si with a E g =1.15eV. This figure has been reprinted from [64] III-V materials present the best single junction efficiencies: GaAs 29.1 %. As shown in Figure 1.19, III-V compounds allow to grow a large variety of bandgaps. Indeed, III-V materials allow to make some quaternary and ternary alloys to set the bandgap to the optimal value of ∼1.7 eV. This bandgap can be obtained with InGaP with a certain composition in P. As presented in Figure 1.19, III-V and Si have a high lattice mismatch.

No III-V semiconductors, with E g =1.7eV, is lattice-matched with Si. For instance, GaAs presents a 4 % lattice mismatch with Si. Strain induced by this lattice mismatch will generate dislocations. Other dislocations can be induced through the high difference in thermal expansion coefficient and the high growth temperature for III-V. Aside the lattice mismatch, the growth of polar crystals on non polar crystals leads to anti-phase domains. Besides, stress can arise from the difference in thermal expansion coefficient between the materials, leading to dislocation generation during cool-down from the growth temperature [START_REF] Bett | III-V compounds for solar cell applications[END_REF] [104]. These Si/III-V cells based on direct growth only produces limited efficiencies (25.9 % for a GaAsP/Si [67]). The low efficiencies are mainly due to the high density of defects induced by the lattice mismatch.

In order to overcome this issue, mechanical bounding and wafer bounding solutions are under investigation. They allow to fabricate both cells separately and then to bound them together, reducing density of defects. Mechanical stacking has reached 34.1 % for a GaAs/Si cell [67]. An other solution which has been explored is the use of nanowires. The tolerance to lattice mismatch for nanowires is plotted in Figure 1.19. Elastic relaxation of the strain allows epitaxial growth of nanoscale materials. Besides, difference in thermal expansion lattice-mismatched structures at the nanoscale are likely to respond elastically to dissimilar expansions under thermal stress without breaking the structure. III-V/Si Tandem solar cells have been shown to be able to reach up to 38 % efficiency theoretically. Si is a good candidate for the bottom cell as it is a cheap and efficient technology. Because of the highly tunable bandgap and the high material quality, III-V materials are good candidates for the top cell. But because of a high lattice mismatch, it is difficult to grow directly high quality III-V on Si. To overcome this difficulty, solutions are explored such as mechanical or wafer bounding. In order to grow a dislocation-free planar layer, thicker than 10 nm, a < 1 % lattice mismatch is required. But, for nanostructures with a radius of 100 nm, it is possible to grow a structure with almost 3 % lattice mismatch as observed in Figure 1.19. Nanowire and nanopillars are nanostructures with a radius of about 100nm. This small footprint can accommodate the strain without the creation of dislocations. Besides, nanowires present also other advantages compared to planar structures [START_REF] Borgstrom | Towards Nanowire Tandem Junction Solar Cells on Silicon[END_REF]. They are well suited for PV as there is an enhancement of the broadband absorption and anti-reflection. The use nanoscale materials is interesting as it induces light scattering (because of their size) and so light collection enhancement [111].

III-V nanowires direct growth on Si

Proof concept of GaAs nanowires grown directly on a Si substrate have been developed reaching 3.3 % [19] and 4.1 % [41] power conversion efficiency. But efficiencies are still hindered by low V oc . These low V oc can be explained by a high defects density introduced during the nanowire growth or low doping levels due to the particular growth technique. Besides due to the high surface to volume ratio, carriers lifetimes in the nanowires are very dependent of surface recombinations.

Tandem solar cells present an opportunity to fabricate high efficiency solar cells. We have seen that the use of nanowires offers the opportunity to grow directly high quality III-V materials on Si. Proof of concepts have been fabricated reaching up to 3.3% [19] and 4.1 % [41] power conversion efficiencies. But, cells efficiencies are still very dependent of the improvement of the absorber quality (lifetimes, doping and surface passivation).

1.4. Materials characterization and aim of the thesis

Materials characterization and aim of the thesis 1.4.1 Overview the main characterization techniques of PV materials

We have seen previously that solar cells based on polycrystalline thin-films and nanowire suffer from poor V oc values which can be explained by many factors such as low doping levels, short lifetimes and high defects densities. In this Subsection we will briefly describe some characterization tools and techniques, their working principles and how they can be used to determine these quantities. We will discuss on their relevance for the characterization of the doping level and optical or transport properties. Besides, their spatial resolution will be discussed for the applicability on nanostructures. Examples on nanowires or thin-films solar cells will be described.

Hall effect measurement

Doping is a fundamental property of SC and its control is crucial for PV applications but also other applications. A standard way to measure doping levels in thin SC layers is the Hall effect measurement (Figure 1.21 a)). The charge carrier type (electrons or holes), carrier concentration and mobility are determined from the measurement of the Hall voltage, which results from the externally applied magnetic field to an electrical current. It requires four contacts (Figure 1.21b)) and the layer thickness is known in advance.

Hall-effect measurements have been applied to measure the electron concentration in InAs nanowires using an optimized electron-beam lithography process for the contacts deposition [START_REF] Essig | Progress Towards a 30% Efficient GaInP/Si Tandem Solar Cell[END_REF]. The spatial variations of the carrier concentration in InP was also analyzed by placing multiple contacts [167]. [32] and c) [START_REF] Ch | Hall effect measurements on InAs nanowires[END_REF] This characterization technique is perfectly suited for thin-films doping level characterization as it is a very straightforward and precise method. But, contacting nanostructures such as nanowires is a complex and time consuming task this is why it is difficult to see such characterization methods as a viable solution. Indeed, contactless methods are more suited for the characterization of nanomaterials. In this thesis, we aim at developing contactless method of determining the doping levels, among other materials properties, in nanostructures SC.

Resonant Raman spectroscopy

Resonant Raman spectroscopy has been used by Hakkarainen et al. [73] as a contactless method of determining the doping level in GaAs NWs. This characterization method is commonly used to elucidate the phonon-electron interactions for both bulk and nanoscale materials. Lightmatter interactions between an incident photon and a SC gives rise to different mechanisms: elastic scattering (Rayleigh) or inelastic scattering (Stokes or anti-Stokes) [START_REF] Smith | Modern Raman Spectroscopy: A Practical Approach[END_REF]. Most of the interactions are due to elastic scattering interactions where the matter is excite or desexcite through a virtual state corresponding to an other arrangement of the electron states, without loss or gain of energy. When the incoming photon has an energy near an electronic state, the scattering process can be enhanced inducing an increase of the signal 1.22. This technique has also been used by Ketterer et al. [91] in order to determine the different phases in a crystalline single GaAs nanowire. Besides, they use the link between the incoming light polarization and the probability of light scattering to analyze the existence of zinc-blend (ZB) and wurtzite (WZ) phases within a GaAs NW. Resonant Raman spectroscopy can be used to probe both the doping level, crystalline phase but its spatial resolution is limited to the micrometic scale because of its light based source of excitation (see discussion below). In order to estimate the homogeneity of these quantities in nanomaterials, a nanoscale spatial resolution characterization will be used in this work.

Hyperspectral luminescence mapping

Luminescence of SCs contains a wealth of information on material properties, like absorption spectra, quasi-Fermi level splitting, or energy levels within the band gap as will be further described in Subsection 2.4.1. This is why luminescence coming out of the SC can be used as a diagnostic tool for the quality assessment of a SC. Tools for quality assessment such as electroluminescence (EL) and photoluminescence (PL) are widely used for solar cell characterization.

Electroluminescence:

The most popular imaging method to investigate the influence and properties of electrical losses within a large-area PV module is the electroluminescence (EL). EL is the phenomenon in which a p-n junction emits light in response to an electric current that is passed through it by forward biasing the p-n junction, when the carriers injected recombine in a radiative process. The radiative efficiency, at a given voltage, and the analysis of the emitted luminescence spectrum gives information regarding the absorption, resistance losses, or defect levels.

By imaging the luminescence, we can analyze the spatial decay of EL signal from the electric contact to determine the sheet resistance of the window layer [START_REF] Paire | Measuring sheet resistance of CIGS solar cell's window layer by spatially resolved electroluminescence imaging[END_REF]. Further analyses of the EL maps, recorded with an absolute calibrate setup, lead to the quantitative determination of the spatial variations of properties such as the quasi-Fermi level splitting [44]. EL imaging is commonly used as it allows direct imaging of defects in cells, which appear as darker areas because less radiative emission is expected near defects, over wide surface areas. EL imaging has a spatial resolution which depends on the CCD camera and the collection optics system. The spatial resolution is limited by Rayleigh criterion. Besides, this characterization technique is limited by the fact that it requires contacts to apply a voltage bias.

Photoluminescence: Photoluminescence (PL) is a contactless method where a laser is used as source of excitation. The incoming photons are chosen to have an energy higher than the SC bandgap in order to generate minority carriers through light absorption in the SC. The signal induced by the radiative recombination of the electron-hole pairs is called photoluminescence. The working system for a classic PV setup is presented in Figure 1.23. A laser is used as a light source. Then the light can be focused on an excitation point (local injection) or set in order to enlighten the whole cell (broad field). A beam splitter reflects the PL signal toward a CCD camera in order to record the luminescence map coming out of the sample. An hyperspectral imager can be added to the system in order to record hyperspectral maps. Besides, the use of a laser allows setting rapidly and efficiently the power of the excitation source.

Chang et al. [30] have analyzed the evolution of the optical properties of a single GaAs NW using PL (see Figure 1.24) by moving the PL spot along the NW axis. They estimated the evolution of the PL peak position and linewidth as a function of the position of the PL spot along the NW, and related the trends with surface passivation of the NWs. Pulsed excitation for lifetimes measurements can also be performed by using a pulsed LASER. Indeed, as described in 1.2.4, the carriers density decreases when turning off the light source in order to reach the equilibrium. As the luminescence is directly proportional to the product of the carriers densities, measurement of the PL signal decrease gives a way to measure the carriers [30] lifetime (Equation 1.16).

PL is used industrially in order to assess for the materials quality as a fast and non destructive characterization technique. PL mapping coupled with advanced modeling allows assessing for the transport properties of PV materials. Still, because of the use of the light as an excitation source, the spatial resolution is limited due to light diffraction. Confocal microscopy allows for spatial resolution to go down to sub-micrometric scale (≈500 nm). But in this thesis we aim at characterizing nanomaterials whose size is of about 500 nm or below. This is why photoluminescence is not suited for determining homogeneity of properties within a nanomaterials. However we have seen that luminescence-based techniques are well suited for the characterization of PV materials.

Transmission electron microscopy

Electron microscopy has already been shown for a long time to be able to reach very high spatial resolution. The de Broglie wavelength of electrons is below 0.1 Å, for electron energy above 20 keV [START_REF] García De Abajo | Optical excitations in electron microscopy[END_REF]. The excitation source is no longer limited by the diffraction, as in the case of the photon source. Electron microscopes thus allows for the analysis of materials with both localized (STEM) and extended (TEM) source of excitation. In transmission electron microscopy (TEM) and scanning electron microscopy (STEM), an energetic electron beam (50-400 keV) is used on thin samples (<100 nm). The sample thickness makes it almost transparent to the electron beam. Deviations of the electron beam are induced by the interactions with the electronic cloud of the sample. The contrast of the STEM and TEM images depends on the atomic number Z of the atoms. Besides, elastic interactions between the electron beam and the sample give rise to diffraction patterns, for crystalline materials, which can be analyzed to determine the crystalline phases.

TEM is used to investigate the crystal structure of individual NWs at the atomic scale.The high potential for very local analysis is highlighted in Figure 1.25 where TEM images and micrographs on a single GaAs NW show the Wurtize and Zinc-Blende phases formed during the growth [32]. Besides, more insight on their chemical composition is gained when using X-rays detectors to perform electron diffraction X-ray spectroscopy measurements (EDX). de Lépineau et al. [START_REF] De | Evidence and control of unintentional As-rich shells in GaAs 1-x P x nanowires[END_REF] have evidenced the growth of an unintentional As-rich shells during the growth of GaAsP NWs. They accurately measured the core and shell alloy compositions and they estimated the shell thickness, using high angle annular dark field (HAADF) and EDX profiles.

TEM experiments have shown to be a very efficient way to characterization crystalline phases, composition fluctuations or dislocations within a SC at the atom scale. But, no direct information on the transport (mobility, lifetimes,...) or the optical (E g ,...) SC properties can be extracted from TEM measurements. Besides, they can only be performed on very thin samples and so necessitate complex sample preparation in the case of thick samples (>200 nm).

In this Subsection, we presented some characterization techniques which care about similar issues as we are interested in: doping level estimation, optical and transport properties. We described how they can characterize these quantities and their working principles. But we also raised the questions of their applicability for nanostructures in the case of the Hall effect measurements and their spatial resolution in the case of the contactless characterization techniques (Raman spectroscopy, PL and EL). Besides, we discuss of the relevance of luminescence-based techniques for the characterization of PV materials. Electron microscopy presents a very high spatial resolution thanks to an electron-based source of excitation but it does not allow for the analysis of the optical and the transport properties.

Aim of the thesis: Characterization of nanomaterials by cathodoluminescence

Previously we highlighted that thin-films and the nanowire-based solar cells are still hindered by high defects density, poor doping levels and short minority carriers lifetimes. In order to further improve these technologies, an estimation of these quantities is of great importance. But, because of their very small size, conventional characterization techniques cannot be employed. Indeed, it is necessary to use a characterization tool that has a resolution comparable to their size.

Electron microscopy has been shown to be able to reach atomic scale spatial resolution. High energy electrons also induce electron-hole pairs within the SC which can recombine radiatively to generate a luminescence signal: it is the cathodoluminescence (CL) [START_REF] Ozawa | Cathodoluminescence Color of Y2O2S: Eu Phosphors[END_REF] . CL combines the very high spatial resolution of the electron microscopy with the luminescence which has already been shown to be a very relevant way to characterize PV materials. The mechanisms leading to the photon emission can be compared with other luminescence-based techniques such as PL, even if differences associated with the details of the excitation process may arise. Besides, in contrast to TEM experiments the sample thickness does not have to be of hundreds of nanometers as the electron beam does not have to go through the film thickness. CL can thus be performed on solar cells structures.

A scheme of the working principle of the CL is presented in Figure 1.26. The electron beam, described as a red cone, is focused on the SC. First, secondary electrons will be generated through the first interactions between the electron beam and the material, forming a SEM signal (Figure 1.26, mechanism 1)). Electron-hole pairs can also be generated through inelastic interactions between the SC and the incident high energy electrons. Just to provide an order of magnitude, ≈ 10 3 electron-hole pairs are estimated to be generated by each incident electron with energy of 10 keV. Then, carriers can either recombine radiatively or be collected resulting in: CL (Figure 1.26, mechanism 3)) or electron beam induced current(Figure 1.26, mechanism 4)), respectively. Besides, a mapping of the sample can then be performed by exciting and collecting a CL spectrum at each point of excitation.

In this thesis we are interested in the investigation of nanomaterials for PV application using CL. We based our analyses on the unique specifications of the CL tool located at the C2N. We developed methods to explain the mechanisms which hinder the efficiency within CdTe poly- In this Chapter I will describe the C2N cathodoluminescence setup, its main features and its particularities. I will first focus on the description of the tool and the technological aspects, going from the tip of the electron gun to the photon collection and the signal processing, in Section 2.1. Besides, electron-matter interactions will be discussed for insights on the spatial resolution and the injection level.

In Section 2.4, theoretical background of the luminescence of semiconductors will be explained. Motivations for the absolute calibration will be clarified in Subsection 2.4.1. Absolute calibration of the CL is needed in order to reach the full set of information accessible with a fit of the luminescence spectrum. In Subsection 2.4.2, the protocol for the development of the absolute calibration of the C2N CL setup will be explained. Then preliminary results will be given.

The CL setup: main features

An Attolight Allalin 4027 Chronos tool equipment has been used for all the experiments. This tool [3] possesses very specific capacities as quantitative comparison between samples and the abilities to perform time-resolved experiments. In this Section, the C2N setup will be described in details starting from the electron source to the photon detector, starting with a brief overview of the full setup in Subsection 2.1.1. Electron source settings will be introduced. Then I will focus on the description of the light collection optics and the dispersion and detection system. CL mapping principles and the data processing will be described. The change from continuous to pulsed conditions will be explained in Subsection 2.2. Comparison with other time-resolved systems will be introduced. I will end by studying the electron-matter interactions for CL spatial resolution and injection levels assessment.

Overview of the cathodoluminescence setup

The CL setup is composed of four distinct ensembles. In Figure 2.1, they have been distinguished using different colors: the emission system (grey), the light collection optics (yellow), the dispersion and detection system (red) and the time-resolved system (green). Cathodoluminescence setup scheme where four ensembles have been distinguished: the emission system (grey), the light collection optics (yellow), the detection and dispersion system (red) and the time-resolved system (green). The different system have not been drawn at scale.

The emission system (Subsubsection 2.1.2) is composed of an electron source which provides an electron beam with a current (I) at a given electron energy (acceleration voltage) and a magnetic lenses to focus the beam on the sample. Secondary electrons are detected to form SEM images.

The light emitted by the sample is collected via a collection mirror above the sample, and sent out of the CL chamber by a reflective objective. The light which comes out of the chamber is collimated.

The light is then sent to the different spectrometers by reflective mirrors. Achromatic objectives at the entrance of the spectrometers focus the light on the entrance slits of the spectrometers. The spectrometers possess different exits with different detectors. At the exits of the "Spectrometer 1" a CCD camera and a streak camera have been placed. For the "Spectrometer 2", exits are occupied by an InGaAs camera and a time-correlated single photon counting system (TCSPC). Time-resolved cathodoluminescence is achieved with a pulsed LASER excitation source system and is associated with the streak camera or the TCSPC detectors. The electron beam current and the electron energy are set by the emission system. The system is very similar to a classic electron microscope which is used for SEM measurements. Electrons are extracted from a tip (ZrO coated Tungsten) which has been thermally excited. Electrons are then gathered and accelerated by electromagnetic lenses. The acceleration voltage is set at the very top of the column. In our system, it can be set between 1 kV and 10 kV. The incident current is set using two parameters: the gun lens (GL) and the aperture. The gun lens is an electromagnetic lens at the top of the column which controls the current density. The diameter of the aperture (from 25µm to 100µm), positioned after the gun lens, is set in order to change the current. The variation of the incident current along the day is estimated to be less than 5%. The electron microscope alignment is then made manually. A set of four different parameters needs to be optimized to obtain a perfect focus of the beam: 1) The vertical position of the sample 2) an objective lens 3) a deflector and 4) a stigmator. Once the focusing is completed, the beam current can be measured using a Faraday cup, in which the beam is focused, and an electron beam induced current (EBIC) system which measures the current collected by the cup.

The electron source

In Figure 2.3, the incident current as a function of the GL and the aperture current is plotted. As we can observe, by setting the GL and the aperture, we are able to change the injection current by orders of magnitude. Consideration on the incident current has been made prior to any CL experiments. Impact of the electron beam energy on the measurements is discussed in Subsection 2.3.

Light collection system

The mirror is integrated in the column (see Figure 2.2) so its position in the column is fixed. Collection optics optimization is simply made by setting the sample in the focal plane of the collection mirror using the nanopositioning system (see Subsection 2.1.4). It is done by optimizing the luminescence signal. This configuration of the collection system allows for quantitative comparison between samples as no optical alignment is needed when modifying the sample, except for the setting of the sample height.

The collection optics is composed of a Cassegrain mirror disposed above the sample for photons collection (Figure 2.20. It possesses a constant and high collection efficiency over a large field of more than 150 µm in diameter. Collection optics has a numerical aperture (NA) of 0.72 corresponding to a half angle of 46°with a central opening of 9°for the beam. Once collected, the light is collimated and sent out of the CL chamber by a reflective objective. The exit of the chamber is a transparent window made of fused silica.

Nanopositioning system and control of the temperature

The stage is composed of a motorized nanopositioning system (see Figure 2.2) and a cryostat. The cryostat monitors the temperature of the stage using a He flux. Temperature is monitored with a 0.1 K precision on an interval of temperatures ranging between 10 K and 150 K. Below 10 K, the monitoring of the temperature is less stable because of a very high He flux which can vary abruptly, causing shift in the positioning. Above 150 K, the control of the temperature is more challenging as a very weak He flux and a high heating are required. Considering a stable temperature, a drift of the stage has been measured to be less than 250 nm/h at low temperature (10 K).

For some experiments it is necessary to accumulate during a long time in order to acquire a sufficient signal/noise ratio. At low temperature it takes 24min for the beam to be drifted from the center of a 200 nm diameter nanowire to the very edge of this nanowire. It allows for a 20 min acquisition before the beam position needs to be reset. Besides, for CL map acquisition, the very stable stage allows us not to have to correct the final image because of drift artifacts.

Dispersion and detection system

At the exit of the CL chamber, the light is collimated and sent to one of the two different spectrometers, depending on the material we study. The light is dispersed by the grating of the spectrometer and detected by the chosen detector. The choice of the grating and the detector depends on two considerations: the spectral resolution and the detection/collection system response.

The grating diffracts light with a spectral dispersion which depends on the groove density. The two grooves densities in the spectral range of interest (600 nm-900 nm) were 150 gr.mm -1 and 600 gr.mm -1 . The two gratings are blazed at 500 nm, meaning that their maximum external quantum efficiency (EQE) is reached at λ=500 nm as observed in Figure 2.4 (red line). The spectral dispersion were then 0.53 nm.pixel -1 and 0.12nm.pixel -1 , for 150 gr.mm -1 and 600 gr.mm -1 respectively. A perfectly punctual source at the entrance of the spectrometer leads to a 4-pixel image on the CCD (CCD Andor Newton DU920P-BEX2-DD), when no dispersion from the grating was considered. It corresponds to a spectral broadening of the incoming signal by the dispersion and detection system of about 2 nm and 0.5 nm for the two grating respectively. This broadening can be used to define a spectral resolution.

CL(λ) = CL recorded (λ) ÷ EQE system (λ) (2.1)
The detectors and the spectrometers have diffraction efficiencies which depend on the wavelength. Indeed, the theoretical spectral response (extracted from datasheets) for the CCD detector and the grating "150/500", with a groove density of 150 gr.mm -1 and blazed at 500 nm, have been determined from the products datasheet and plotted in Figure 2.4. The system response which corresponds to the product of the two responses has been plotted (green line). From the spectral response we can notice that this configuration (grating/CCD) is not suited for measurements of CL signals above 1100nm and under 300 nm. In order to take into account the non perfect spectral response, the recorded CL spectrum is divided by the theoretical spectral In this Subsection I described the C2N Cathodoluminescence tool. A key point of the system is its special collection optics which allow for high collection efficiency over a large surface (>150µm in diameter). Besides, quantitative comparison can be performed between samples as there is no need for optical alignment when moving the sample holder. An other specificity of the tool is the very stable (drift<250nm/h) which allows for very long acquisition times at low temperature.

CL mapping and hyperspectral cube

A CL mapping is performed by scanning a selected area at the surface sample. In Figure 2.5 a) the O point corresponds to the column axis (when no deflection of the beam has been made). The electron beam is deflected from the axis to excite at the M point. At each excitation point a spectrum is recorded. An hyperspectral cube is then obtained by collecting CL spectra for each excitation point.

In Figure 2.5, the scheme of the different steps for the image formation on the CCD from the selected area is presented. First, an image of the scanning area is rotated by 90°at the entrance of the spectrometer. In addition, the scanning area has been magnified by a factor 13 because of the collection system magnification. Then image on the CCD is formed as when no dispersion of the light is performed. CL spectra are obtained by summing vertically over the 256 pixels of the CCD camera. When considering a point away from the column axis (M or M') a spatial shift on the CCD is observed. This spatial shift can be explained by geometrical considerations as drawn in Figure 2.5 a), b) and c). This artifact can be predicted theoretically considering the size of the map, the grating dispersion (0.53nm/pixel) and the CCD pixel width (26µm/pixel) (Equation 2.2). The size of the map is calculated from size=131100/Magnification, where the Magnification is a setting parameter in order to adjust the scanning area. The position of the pixel in the hyperspectral map is defined with its horizontal position a and its vertical position b. The spectral shift F spatial shift

d) and e) represents CL spectra at different location O (black) M (blue) and M'(green) of the hyperspectral map when d) no correction is applied and e) if the spatial shift is perfectly corrected.

due to the deviation of the beam from the horizontal axis can be calculated at any point of the hyperspectral map by Equation 2.2. In this Equation, the CCD camera (1024×256 pixel array with pixel size of 26 µm) and the grating "150/500" (with a light dispersion of 0.53 nm/pixel) were considered. The maximal shift is expected at the edges of the CL maps(M and M' in Figure2.5). We estimate theoretically a shift of 4 nm for a 31µm wide CL map using Equation 2.2:

F spatial shift (a, b) = 1 2 - b pix y ) × 131100 M agnif ication ) 13 × 0.53(nm/pixel) 26(µm/pixel) (2.2)
We measured a CL map at RT on a GaAs thin film. The peak energy position at each pixel has been estimated and plotted in Figure 2.6. The peak energy map in Figure 2.6 a) corresponds to the map before the spatial correction was applied. A clear blue-shift from the top to the bottom of the image is observed. In Figure 2.6 b), the correction of the spatial shift has been applied on the same map as the one presented in Figure 2.6. The shift of the peak is no longer observed and only induces a spectral shift of a few meV.

The different steps in order to determine the emitted signal per unit of energy from the raw data have been exposed in Figure 2.7. First, the background level is assessed. The CL signal which is obtained after the background removal is normalized by the exposure time of the CCD camera (Figure 2.7 b)). Spectral corrections are then applied to the CL spectrum (Equation 2.1). When comparing the uncorrected and the corrected CL spectrum we can observe a modification of both long and short wavelength slopes.

As we used a grating which dispersed the light linearly with the wavelength, the CL spectrum is plotted in cts.s -1 .nm -1 . This luminescence signal corresponds to a spectral density of photon flux per unit of wavelength. In this thesis all luminescence spectra were converted into a spectral density of photon flux per unit of energy (cts.s -1 .eV -1 ) as presented in Figure 2.7 d) using Equation 2.3: Carrier lifetime in a material is an important quantity to assess the transport properties. To do so, a pulsed source must be used. In this Section I will describe how to produce pulses of electrons in order to perform time-resolved cathodoluminescence (TRCL). The first time-resolved system which has been developed used a beam blanker. This technique is very convenient as it allows switching directly from one (continuous) mode to an other (pulsed). But the blanker configuration induces a loss of spatial resolution and is limited by the performances of the blanker in terms of temporal resolution [START_REF] Meuret | Complementary cathodoluminescence lifetime imaging configurations in a scanning electron microscope[END_REF].

∂I ∂λ = ∂I ∂E × h × c λ 2 (2.3)

Time-resolved cathodoluminescence

Figure 2.9: SEM images of the same GaAs NW under CW injection (10kV) a) and Pulsed excitation b). The contrast and the brightness of each image has been readjusted.

In our setup, pulsed electrons beam are generated with the use of LASER pulses (λ=355nm) with a power up to ≈100mW. First results using TRCL with the LASER configuration have been obtained in 2005 [START_REF] Merano | Probing carrier dynamics in nanostructures by picosecond cathodoluminescence[END_REF]. Prior to pulsed conditions, the column settings need to be made under continuous source condition for simplicity. Once the column is set, the thermal excitation of the tip is turned off. Then the LASER is focused on the tip using motorized optical lenses. LASER pulses act as an excitation energy source to generate pulses of electrons which are then accelerated and sent to the sample. A pulse is sent to the SEM tip each 12ns with a 5ps width. In Figure 2.8, the blanker and the LASER configurations are compared.

As we change the injection regime from continuous to pulsed, the electron flux is modified. The SEM and the CL signals are directly correlated with the current of electrons. When switching from continuous to pulsed excitation conditions the signal is by about of 3 orders of magnitude. In Figure 2.9 we compared SEM images of the same nanowire under continuous and pulsed injection. As we can see, the spatial resolution is the same as a comparable contour can be observed but the signal is much noisier in the case of pulsed conditions. The time-resolved lumninescence signal is sent to a spectrometer exit with a streak camera. The spectrometer splits the luminescence signal according to the photon wavelength and outputs a dispersed line signal along the horizontal axis. A streak camera is composed of a photocathode which collects this signal and converts into photo-electrons. Then these electrons are swept along the vertical axis with a bias which has been synchronized with the laser frequency and deflects the photo-electrons as a function of time. The first incoming electrons are directed to the top of a phosphor screen, whereas the last ones are deflected to the bottom. A CCD sensor images the phosphor screen and transmits the acquired spectral-temporal 2D image to the computer [START_REF] Adrien Bercegol | Charge transport in photovoltaics via multidimensional luminescence imaging[END_REF]. Finally, The streak camera measures 2D images with a temporal (vertical) and a spectral (horizontal) dimension (see Figure 2.10).

Figure 2.11: TRCL decay measured on a short lifetime sample (A2033) with a streak camera (red curve). The decay has been fit (blue curve) using the convolution of a Gaussian function (FWHM = 12 ps) with an exponential decay (τ =8ps) in order to estimate the time response of the system. The time scale has been shifted by 400 ps compared to

The temporal resolution of our setup has been determined using a sample for which the lifetime has been measured to be about ≈10 ps. A decay for the reference sample is plotted in Figure 2.11. The system response is assumed to be a 1-Gaussian function, characterized by a standard deviation σ. So we fit the decay with the convolution of the system response and an exponential decay. The FWHM of the Gaussian has been measured to be F W HM = 2 2 ln(2)σ = 12ps and the decay time of the 1-exponential function of 8 ps. The FWHM of the Gaussian function is assimilated to our temporal resolution.

The combination of the high spatial resolution and the pulsed excitation conditions allows probing carrier dynamics in nanostructures. In this thesis, it has been used for lifetimes measurements within micrometric grains (see Subsection 3.4) or at different locations of single NWs (see Subsection 4.15).

Electron-matter interactions

The study of the electron-matter interactions is primordial in order to assess the interaction volume of electrons, which impacts the spatial resolution and the carriers injection level. In Subsection 2.3.1 I will focus on the estimation of the interaction volume using a Monte Carlo simulation. Then simple and complex model will be used in order to assess the injection level based on these simulations.

Interaction volume: CASINO simulations

The interaction volume is determined prior to each measurement using CASINO simulations. The CASINO software is used in order to model the trajectory of an electron, based on Monte-Carlo simulations [START_REF] Drouin | CASINO V2.42-A Fast and Easy-to-use Modeling Tool for Scanning Electron Microscopy and Microanalysis Users[END_REF]. CASINO assumes a Gaussian-shaped incident electron beam, in our case we considered a 10 nm beam radius (see Subsection 2.1.2). The electron-matter interactions are described by the trajectory of electrons within the sample where only elastic collisions are assumed. The distance between each collision is estimated from the weight fraction and the atomic weight of each chemical element and the material density ρ (g.cm -3 ). Energy is transmitted to the matter through continuous energy loss until the electron energy becomes lower than 50 eV.

Trajectories of incident electrons within the material are described for a 6kV electron beam energy in a GaAs substrate in Figure 2.12 a). Contour plot of the dissipated energy within the material is presented in Figure 2.12 b). A pear shape of the interaction volume is observed. Besides we can notice that 90 % of the incident energy is dissipated within a distance of about 150 nm. Thereafter we will assimilate the characteristic distance R e to the distance from which 90% of the incident energy is dissipated.

Figure 2.12: Results of the CASINO simulation for a 6kV acceleration voltage and a 10 nm probe size in a GaAs sample. a) Trajectories of electrons within the material. Red lines describe back-scattered electrons. b) Distribution of the dissipated energy summed over the y-direction (perpendicular to the plane). The contour plot indicates the percentage of total dissipated energy within the corresponding contour.

Several factors may limit the CL spatial resolution. The first one is the probe size, which is in our case of 10 nm. However, we just observed that the interaction volume is larger than the gun probe size (R e > 10nm) thereby reducing the spatial resolution. Besides, once electronhole pairs are generated within the interaction volume a concentration gradient is created so diffusion can take place. Since the light is collected from the whole surface, eventual radiative recombination away from the injection spot induced by the diffusion of carriers may occur. For mostly all PV materials the diffusion length is in the order of L dif f ≈ 1µm or longer so the diffusion phenomenon limits the CL spatial resolution. Still, the diffusion length can be affected by potential barriers (heterostructures) and/or by the sample temperature. Indeed, a decrease of the temperature has been shown to reduce the diffusion length, such a phenomenon is also observed in this Thesis. Reducing the acceleration voltage tends to improve the spatial resolution by decreasing R e . But, for multi-layer structures (such as passivated NWs, see Subsection 4.4) it is important to estimate if the carriers are generated in the active layer. CASINO simulations can also be run for multi-layers with the aim of estimating the penetration depth of the beam. In Figure 2.13 two CASINO simulations are compared for the same sample but with different acceleration voltages. For 3kV (Figure 2.13 a)), the penetration depth is not sufficient to generate enough carriers into the active layer (GaAs). When the acceleration voltage is increased up to 6kV, electrons (blue trajectories) reach the active layer but the width of the pear-shape interaction volume is much larger (130nm for 6kV compared to 50nm for 3kV).

Estimation of the injection level

The density of carriers injected within the material may impact the luminescence spectrum and/or the transport properties. Indeed, defects saturation, screening of local potential or band filling can be observed with high injection levels. Carriers generation with an electron beam has been studied for a long time for EBIC or CL studies [START_REF] Davidson | Advances in the electrical assessment of semiconductors using the scanning electron microscope[END_REF][START_REF] Donolato | Influence of the Generation Distribution on the Calculated EBIC Contrast of Line Defects[END_REF][START_REF] Shimizu | Experimental and theoretical study of energy dissipation profiles of keV electrons in polymethylmethacrylate[END_REF]179]. Electron-hole pairs generation rate with CL (G) is described by Equation 2.4:

G = (1 -γ) × i b × E b q × E i = (1 -γ) × E b × i b × Q q × E g (2.4)
Equation 2.4 relates the electron-hole pairs generation rate (G), where E b and i b describe the beam energy (eV) and the beam current (A), respectively. γ is the percentage of back-scattered electrons which can be estimated form CASINO simulations. E i is the ionization potential which is related to the bandgap energy and describes the potential for a high energy electron to ionize electrons of a SC for a given bandgap. The realtion between E i and E g can be approximated by

E i = Eg Q
, where Q is the quantum efficiency for the electron-hole pair generation which depends on the electron beam energy, for typical semiconductors (CdTe, GaAs...) Q ≈ 0.3. As we can notice in Equation 2.4, the generation rate can be set with the acceleration voltage and the current. Generation rates for GaAs and CdTe have been summarized in Table 2.1 considering standard operating conditions: 6kV and I = 2nA. Besides, the relevant parameters for the calculation of the generation rate have been given in Table 2.1.

CdTe

GaAs

ρ (g.cm -3 ) 5.85 5.33 γ 0.33 0.3 E g (eV )
1.51 1.41 G (pairs.s -1 ) 1.0×10 13 1.1×10

Carrier distribution can be described with different degrees of precision (see Figure 2.14).

First, injection can be described as a punctual source generated at a depth d (Figure 2.14). Despite its convenience, the punctual generation is only theoretical. If we want to determine the carrier density we need to represent more accurately the generation distribution. Figure 2.14 b) presents a model where the generation is assumed to be homogeneous within a truncated sphere. But CASINO simulations (see below) and experimental results [START_REF] Shimizu | Experimental and theoretical study of energy dissipation profiles of keV electrons in polymethylmethacrylate[END_REF] suggest that the excitation volume is pear-shaped (at least in GaAs). Indeed, the carrier generation rate is proportional to the energy loss and Figure 2.12 suggests a pear shape of the energy loss. A Gaussian model of the generation distribution (g) has been shown to be a good approximation of the pear-shaped volume of generation (Figure 2.14 c)). In this model the shape of the energy loss of the incident electrons is assumed to be similar to a Gaussian function (Equation 2.5), where a is related to the penetration depth, r is the distance to the excitation point, and θ is the angle with respect to the electron beam. Davidson and Dimitriadis [START_REF] Davidson | Advances in the electrical assessment of semiconductors using the scanning electron microscope[END_REF] proposed to inject directly the Gaussian form of the carrier generation distribution into the diffusion equation. It is difficult to define a precise injection level, but it would be interesting to define an approximated injection level to know whether we are under a high or low injection regime. In order to simplify the model, two limit cases have been considered, the case of infinite surface recombination velocity (v s =∞) and the case of negligible surface recombination velocity (v s =0). Upper (Φ high ) and lower (Φ low ) bounds for the carrier density per second can be determined from these two cases:

g(r, θ) = 4 × G π 3/2 a 3 exp - r a 2 cos(θ) (2.5)
Φ low = G 16 × L 2 × a (2.6) Φ high = G 5 × L 2 × a (2.7)
In Equation 2.6 and Equation 2.7, L is the diffusion length and a is related to the factor R e as a = R e /2. CASINO simulations have been conducted on CdTe and GaAs in order to estimate R e . In our case, we take a diffusion length L = 1µm. Calculated carrier density per second generated by CL in standard conditions (E b =6 kV and I=2 nA) have been resumed in Table 2.2. Then they have been been compared with working conditions under 1-sun excitation (Φ sun ). Considering J sc ≈30 mA.cm -2 and a cell thickness of t = 2µm we can estimate the carriers density per second with Equation 2.8:

Φ 1-sun = J sc q × t (2.8) a Φ low Φ high J sc Φ 1-sun nm (cm -3 .s -1 ) (cm -3 .s -1 ) (mA.cm -2 ) (cm -3 .s -1 )
CdTe 95 6.6×10 24 2.1×10 From the comparison between CL and the estimated 1-sun conditions, we can conclude on the very high injection levels induced by the CL injection. Subsequently, concentrations of carriers injected have been estimated from ∆n = Φ high × τ . Considering τ = 1 ns orders of magnitude of ∆n≈ 10 16 cm -3 have been found. Therefore, the CL injection level should be lower than the effective density of states for the conduction band (N c (GaAs) = 4.2 × 10 17 cm -3 [34] and N c (CdT e) = 9.1 × 10 17 cm -3 [START_REF] Marple | Effective Electron Mass in CdTe[END_REF]) avoiding for band filling. High injection levels may also impact the diffusion because of ambipolar diffusion effects [START_REF] Cadiz | Surface recombination in doped semiconductors: Effect of light excitation power and of surface passivation[END_REF][START_REF] Paget | Imaging ambipolar diffusion of photocarriers in GaAs thin films[END_REF].

Electron-matter interactions have been studied in order to estimate both the injection levels and the interaction volume. From CASINO simulations we observed that the interaction volume is larger than the probe size thus limiting the spatial resolution. Besides, discussion on the role played by the diffusion on the spatial resolution has been discussed. Calculations of the injection levels have also been performed. They have been compared with 1-sun excitation conditions (Equation 2.8). High injection levels have been determined for standard conditions. Still, no artefacts on the luminescence spectra are expected as the injection levels are still below the conduction band density of states.

CL emission and calibration

The luminescence process is the emission of light from a material caused by stimulated or spontaneous emission. The luminescence signal contains crucial information such as its absorptivity, defects energy levels or the fermi level splitting. Besides, from detailed balance considerations it is possible to show that a good radiative emitter is also a good solar cell. This is why, for PV application, characterization techniques based on luminescence provide valuable inforamtion on the material optical and transport properties.

In this Section, we will deal with the theoretical background of the SC luminescence and its application for PV material. First the generalized Planck law and the absorption model will be discussed in order to explain the full luminescence spectrum. Thereafter, a protocol for the calibration of the C2N CL tool and preliminary results will be presented in Subsection 2.4.2.

Luminescence of semiconductors

This part of the thesis has been written based on the reasoning developed by Würfel in Physics of solar cells [178] and summarized by Chen [32]. It aims at describing the luminescence emission by a SC and its application for PV material.

The generalized Planck law

The rate of spontaneous emission of photons, at the energy ω, emitted by a semiconductor (SC) is given in Equation 2.9. The spontaneous emission rate (r sp ( ω)) depends on the refractive index n and the absorption coefficient α at the energy ω. Under excitation, excess carriers are generated and are characterized by the quasi-Fermi levels of electrons E f c and holes

E f v . r sp ( ω) = ( ω) 2 π 3 (c 0 /n) 2 α( ω) exp( ω-(E f c -E f v ) k B T
) -1

(2.9) Equation 2.9 describes the emission rate in the volume of the SC. But what is truly measured is the photon current emitted through a surface. The integration of the difference between the photon emission and absorption rates over the thickness leads to Equation 2.10 which gives the general expression of the light emitted by a semiconductor Φ lum (generalized Planck law):

Φ lum ( ω) = A( ω) × ( ω) 2 4π 3 c 2 0 × 1 exp( ω-(E f c -E f v ) k B T ) -1 (2.10)
Here A( ω) is the absorptivity of the semiconductor layer of a thickness d, and R the reflectivity of an incident light from outside:

A( ω) = (1 -R( ω))(1 -exp(-α( ω)d)) (2.11)
We can then simplify Equation 2.10 by assuming that the emitted photon energy and the splitting of quasi-Fermi levels is large compared to the thermal energy k B T, where k B is the Boltzmann constant and T the temperature: hω -(E fc -E fv ) >> k B T . This approximation leads to a Boltzmann distribution, Equation 2.10 is simplified to Equation 2.12. Where Φ bb is the emission of the black body radiation at a temperature T.

Φ lum ( ω) = A( ω) × Φ bb (T, ω) × exp( (E f c -E f v ) k B T ) (2.12)
We can notice that the Fermi level splitting only impacts the intensity of the luminescence. But absorption and temperature directly impact the shape of the spectrum.

Absorption model

Equation 2.12 can be described by the product of three factors (the black body radiation, absorption and the quasi-Fermi level splitting). In the following we will analyze the model use to describe the absorption in a SC. Different mechanisms can take place within the absorption process. The band-to-band absorption corresponds to the transition of an electron from the valence band to an unoccupied state of the conduction band by absorption of an energy ω. Only photons with energies greater than ω G (E G = ω G ) can be absorbed through this mechanism. Sub-bandgap emission is also observed leading us to assume for a sub-bandgap absorption. Exponential decays of the absorption are mainly observed at low energies, Urbach first modeled the sub-bandgap absorption using a exponential tail with a parameter γ called the Urbach tail [START_REF] Pankove | Absorption Edge of Impure Gallium Arsenide[END_REF][START_REF] Urbach | The Long-Wavelength Edge of Photographic Sensitivity and of the Electronic Absorption of Solids[END_REF]. Katahara and Hillhouse [START_REF] Katahara | Quasi-Fermi level splitting and sub-bandgap absorptivity from semiconductor photoluminescence[END_REF] described absorption in a semiconductor with the convolution of a sub-bandgap Urbach's tail absorption and an above bandgap absorption described by a parabolic absorption model:

α( ω) = 1 2γ ω -∞ α ideal (E)exp( |E -E g | γ )dE (2.13)
Absorption above the bandgap α ideal is proportional to the DOS and can be defined as in Equation 2.14. The term (f v -f c ) takes into account the occupancy of initial and final states. A fit based on the Planck law (Equation 2.12) has been used to analyze the luminescence spectrum for a slightly n-doped GaAs sample (Figure 2.15). As we can see, the fit (red line) is accurate on at least two orders of magnitude. Discrepancy for the high energy slope is due to an accurate absorption model for high-energy photons. For the same reason, sub-bandgap absorption leads to a small discrepancy between the luminescence signal and the fit.

α ideal ( ω) = α 0 ω -E g (f v -f c ) (2.14)
Luminescence spectra have been shown to be directly linked to material properties such as the absorptivity, as described by Equation 2.12. But the fit applied in Figure 2.15, only takes into account for the normalized luminescence spectrum. Additional information regarding the radiative efficiency or the quasi-Fermi levels can be determined from a fully calibrated luminescence spectrum.

Calibrated CL Setup: protocol and preliminary results

Calibration of the setup is motivated by the fact that it allows to directly probe the emission rate induced by a CL excitation. Determination of the radiative efficiency by coupling these measurements with an estimation of the injection level is thus possible (Subsection 2.3.2). Besides, probing the carrier dynamics may enable us to decompose radiative (τ r ) and non-radiative (τ nr ) lifetimes.

The aim is to measure the number of photons emitted by the surface S during of a diode placed in the CL chamber. We designed and fabricated a reference diode for which we aim at determining the photons flux per unit of energy. In order to do so, electroluminescence (EL) spectra have been recorded with a calibrated PL system. Indeed, Delamarre et al. [44,[START_REF] Delamarre | Contactless mapping of saturation currents of solar cells by photoluminescence[END_REF] have developed a calibrated PL imager at the IPVF/IRDEP lab, which will be used as a reference. Then the same EL spectra have been recorded with the CL setup and compared with the calibrated spectra. An electrical bias of 1.1V has then been applied to the diode and a panchromatic electroluminescence image (Figure 2.16 b)) has been recorded. Bright areas correspond to areas from where the light has been emitted. We can clearly distinguish the central part of the diode but we can also observe that luminescence also comes out from the edges.

Calibration of the reference diode irradiance

b) Scheme of the patterns used for the fabrication process of the calibration diode. Red inset presents the full patterns where the diodes (as presented in the yellow inset) with different sizes appear. Back contact are also evidenced (green inset). c) Photo of the CL holder where the pattern and the grooves position have been informed by arrows.

Sample description

Irradiance calibration protocol

In the following paragraph I will explain the protocol we used to calibrated the irradiance of the GaAs diode. The irradiance calibration has been performed using a calibrated PL/EL imager. Raw electroluminescence (EL) signal (Φ(λ, r)) is first normalized with the exposure time ∆t of the CCD after background removal. The HI spectrometer will integrate photons over a wavelength bandwidth of ∆λ, in order not to depend from the data generated by the grating, the collected signal must be normalized by the wavelength interval ∆λ. The signal which is recorded at each pixel of the CCD camera corresponds to a certain surface S of the emitter. This surface can be determined using the pixel size and the magnification of the microscope objective. The signal is thus normalized by the surface in order to obtain a signal in counts.s

-1 .nm -1 .cm -2 . Π diode (λ, x, y) = Φ(λ, x, y) ∆t × S × ∆λ (2.15)
Subsequently, the dependence of the collection efficiency with the position and the wavelength (syst. resp.(λ, r)) must be taken into account. This quantity has been assessed by measuring an hyperspectral image of a reference halogen lamp (I ref λ) coupled with an integrating sphere for which the intensity is assumed to be homogeneous. Thus, the spectral and spatial correction are carried out at the same time based on Equation 2.16:

syst. resp.(λ, x, y) = I recorded (λ, x, y) I ref (λ) (2.16) Φ corr (λ, x, y) = Φ(λ, x, y) syst. resp.(λ, x, y) (2.17)
We state that the angular dependence of the emission of the sample follows a Lambertian law: I(θ) = I 0 × cos(θ). Because of this emission, a part of the photons emitted by the sample is not collected by the system. Hence to obtain the total number of photons emitted by the surface, we must estimate the fraction of photons collected by the sample. In Equation 2.18 we calculated the number of photon truly emitted. Then we calculate the proportion of a lambertian emission signal collected by the microscope objective (Equation 2.19) assuming a numerical aperture NA. A collection limit angle α can be defined as N A = sin(α).

I collected tot = α 0 2π 0 I 0 cos(θ)sin(θ)dθdφ = 1 2 πI 0 sin(α) 2 = I 0 πN A 2 (2.19)
The number of photons truly emitted (Equation 2.20 ) by the sample in cts.s -1 .nm -1 .m -2 is therefore determined by considering all the correcting factor discussed above:

Π emitted (λ, x, y) = Φ corr (λ, x, y) ∆t × S × ∆λ × N A 2 (2.20)
Then absolute calibration is performed by converting an intensity (counts) to a number of photons. To do so, we estimated the Ratio (cts.photons -1 ) to convert a count in terms of photons (Equation 2.21). The ratio value for this collection system has been determined with the use of a LASER. The photon flux emitted by a LASER N photons has been measured using a Si diode and compared to the correponding number of counts measured by the CCD N LASER :

Ratio = N LASER N photons (2.21)
Once absolute calibration has been performed, each spectrum which is recorded at each pixel of the EL hyperspectral cube is given in terms of a photon flux density per unit of energy (photons.eV -1 .s -1 .m -2 ). The photons fluxes emitted by the edges and the center of the diode of the EL map (Figure 2.17 a)) have then been determined by integrating over the energy and the surface of the features. They have been summarized in Table 2

Calibration of the CL setup

The photon flux under a 1.1V bias has been determined in Subsection 2.4.2.1. The aim is to convert the number of counts into a number of photons. The CL collection optics are designed to collect the light emitted by the material after being excited by the electron spot in the field of view (> 150µm in diameter). In Figure 2.17, we observed that a non-negligible part of the light is emitted by the edges of the diode.

The diode which was used in Subsection 2.4.2.1 has been placed in the CL chamber with electrical contact in order to biased the diode under 1.1V. Then EL spectra were measured with different entrance slit apertures of the spectrometer. On the first row of Figure 2.18, images recorded on the CCD camera for different apertures are presented. We can clearly recognize the diode, with intensity from the center of the diode but also from the edges. We can observe that as we reduce the entrance slit aperture, the proportion of the intensity coming out of the edges is reduced. EL spectra have been determined by averaging vertically intensities on the CCD images. A main peak at 870 nm is observed but also shoulders at 830 nm and 900 nm. The two shoulders are due to the fact that we don't have a punctual source. Total intensities are then determined by summing intensities recorded for each wavelength. We already showed that the spectral response of our dispersion and detection system depend of the wavelength (see Figure 2.4). We assumed that all photons are emitted at 870 nm, which corresponds to the maximum of the EL spectrum for the lowest entrance slit aperture. At this wavelength, the system response is syst.resp.(870nm) = 0.31.

An other factor which needs to be taken into account is the angle dependence of the sample emission. Indeed, the collection efficiency is limited to a given angle. As before, we assume a lambertian emission (Figure 2.20). The theoretical collection efficiency, is determined by the integral of the emission collected between α 1 and α 2 . α 1 corresponds to the maximum collection angle and α 2 to the minimal angle due to the aperture in the collection optics for the electron beam. A collection efficiency of f coll. eff. = 0.4 of the Cassegrain mirror is thus calculated. In Figure 2.19 is plotted the evolution of the EL intensity as a function of the area proportion of the diode which is intercepted with the entrance slit. When the entrance slit is fully opened (slit= 7000µm in Figure 2.18 a)) luminescence coming out from all the pattern is intercepted (edges and center) corresponding to areaproportion = 100%. On the other hand, when the slit is slightly closed the contribution from the edges to the total luminescence starts decreasing. An affine function has been used in order to fit the evolution of the intensity with the area proportion. From the fit we can estimate the luminescence intensities, in cts.s -1 , coming out of the center of the diode (I diode ) and the global intensity (edges and center: I diode+edges ). By combining these results with the photon fluxes (photons.s -1 ) estimated previously and summarized in Table 2.3, we can thus determine two values for the Ratio (Ratio diode and R diode+LASER ) for the conversion of counts into a number of photons.

An other method has been used in order to attest for the Ratio value. A LASER (λ = 532nm) has been introduced directly into the spectrometer and spectra have been recorded for different LASER power. For each LASER power, a calibrated Si diode has been used in order to measure the incident power. A linear function has then been used in order to fit the evolution of the total intensity, in cts.s -1 , and the LASER incident photon flux. The calibration factor (Ratio LASER ) can be determined directly with the slope of the curve. Very similar results are obtained for the Ratio values between the two methods (LASER and the calibration diode methods).

All results for the calibration factor determined from the different methods are summarized in Table 2.4. The order of magnitude of the ratio we obtained can be justified using data sheet of the CCD. The CCD is described by a 256×1024 pixels (vertical×horizontal) array. The theoretical conversion factor of counts into photons is 10. Theoretical collection efficiency of the Cassegrain mirror and quantum efficiencies of CCD and the grating have already been discussed to be 0.41 and 0.31 respectively at 870nm. Besides, 5 deflector mirrors can be numbered in the light path. Assuming a 0.9 reflectivity for each mirror, a f mirror =0.59 needs to be taken into account. When multiplying these factors, a theoretical Ratio R theoretical = 3.6×10 4 photons.cts -1 is obtained. A deviation of a factor 3.75 between the theoretical and experimental values which can be explained by non-ideal transmission of the achromatic objective at the entrance of the spectrometer or of the window at the exit of the CL chamber. Besides, variations between actual performances of the CCD and the theoretical performances extracted from datasheet can be discussed.

In this Subsection we aim at determining the calibration ratio which corresponds to the conversion factor between an incident photon and the resulting number of counts. Different methods have been used to this end. First, a calibration GaAs diode has been designed and fabricated. Then, a reference calibrated hyperspectral imager (IPVF) has been used to assess for the photon flux emitted by the diode. By comparing the resulting photon flux with the electroluminescence signal collected in the CL for the same diode, we were able to determine the conversion ratio of ≈ 1.5 × 10 5 photons.cts -1 . An additional method based on a LASER has also been deployed in order to assess the order of magnitude of the calibration ratio. Similar calibration ratio have been found with the two methods. Besides, these values have been compared to a theoretical considerations to check for the consistency of the calibration ratios.

Conclusion of the section

Luminescence spectra can be described by the generalized Planck law 2.10. The fit of the luminescence spectrum allows obtaining a set of relevant parameters for PV applications. In order to go a step further in the materials analysis we explored the absolute calibration of our CL tool. To this end, a GaAs calibration diode has been elaborated and its photon flux under a 1.1 V bias has been measured using a reference calibrated hyperspectral imaging system. By comparing this result with the EL signal recorded with the CL setup, the conversion ratio (cts.photons -1 ) between the number of counts and the number photon emitted by the diode has been determined. A conversion ratio of ≈ 1.5 × 10 5 photons.cts -1 has been found. This achievement is a preliminary result and is still under investigation for accuracy. It has not been used in the following thesis.

Summary

In this chapter I described the cathodoluminescence setup installed in the C2N laboratory, its main features, and the basic principles of the technique.

• We first presented the parameters which control the acceleration voltage and the electron current. We showed, in Subsection 2.1.2, that we are able to control the electron beam current over several orders of magnitude (Figure 2.3). As a result of the specific collection optics which has been described in Subsection 2.1.3, we can perform quantitative comparison of the light emitted between the samples. The dispersion system is composed of two spectrometers with four detection systems. The spectral resolution of the system used in our experiments (grating: 150/500 and CCD) is estimated to be of 2 nm. We ended by describing in Subsection 2.1.6 the hyperspectral mapping of the sample, the data processing and the spectral correction.

• The experimental configuration used for time-resolved CL experiments has then been described (see Subsection 2.2). A streak camera is used a detection system for time-resolved experiments, allowing to be both spectrally and time-resolved. The temporal resolution of our setup with the streak camera has been estimated (Figure 2.11).

• In Section 2.3, the electron-matter interactions were described. The distribution of dissipated energy is imaged using Monte-Carlo simulations. Injection levels and its dependence on the different setting parameters (E, I b ) was described in Subsection 2.3.2. Injection levels were estimated for standard settings parameters (6kV, 2nA) and compared to the 1-sun injection level.

Chapter 3. Polycrystalline CdTe thin-films

We described in Section 1.3.2 polycrystalline thin films solar cells as promising materials. As of now, these technologies suffer from poor V oc which hinders their efficiencies.

In the following chapter we will study polycrystalline CdTe thin films. The main limitations are ascribed to the very nature of polycrystalline thin films, and has multiple origins: the limited grain size, the detrimental role of grain boundaries, bulk defects, and the low carrier concentration. These limitations will be described in Section 3.1. We end by highlighting why we choose to use high resolution CL mapping in order to analyze polycrystalline CdTe layers. During the past decades, the development of post-deposition treatments has enabled tremendous improvements of the material quality and doping levels. In the case of CdTe layers, CdCl 2 annealing has historically been used in order to improve greatly efficiencies. The annealing process has been optimized but no detailed explanation of the role played by the temperature has been explicated. This is why we conducted a study, in Section 3.2, on the exact role of CdCl 2 treatments on CdTe, and espacially the role played by the temperature. Recent improvements have been made by adding Se in the CdTe absorber. The add of Se in the absorber first aimed at improving the J sc by lowering the bandgap, but surprisingly it had not reduced the V oc value. We analyzed the influence of Se quantity on the radiative efficiencies, the grain sizes and the bandgap of the layers. Besides, preliminary results on the lifetimes assessment in single grains will be presented in Section 3.4. An original protocol, based on lifetimes and grain size assessment, for the quantitative assessment of grain boundary recombination velocity and bulk lifetimes will be explained. Feasibility will be proven with low temperature TRCL measurements.

Background and state of the art of CdTe

In this chapter we will focus on CdTe polycrystalline thin films. Based on a bibliographic study, this Section studies the questions at stake for future material improvements. CdTe efficiencies and cells parameters will be compared to other mono or polycrystalline thin films technologies in Subsection 3.1.1. Reasons for the observed high V oc deficit are identified and explained in Subsection 3.1.2. A focus on the complex role played by the grain boundaries or other interfaces on the efficiencies will be made in Subsection 3.1.3. First results on the analysis of CdTe layers using CL we will shown in Subsection 3.1.4 as a justification of our use of the high resolution CL mapping for CdTe optical or transport properties assessment.

State of the art for CdTe polycrystalline thin films

CdTe thin films technologies performances have been greatly improved since 2010 as shown in Figure 3.1. First improvements have been made through the J sc optimization by minimizing the parasitic absorption within window layers. In addition of this achievements, annealing of CdTe thin films under CdCl 2 atmosphere has been found to be an efficient way to boost the performances of CdTe solar cells [START_REF] Hajimammadov | Effect of "CdCl 2 Treatment" on Properties of CdTe-Based Solar Cells Prepared by Physical Vapor Deposition and Close-Spaced Sublimation Methods[END_REF][132] [7]. It is often referred as an activation treatment even if the specific mechanism involved is not yet fully understood. Last achievements have been realized by adding Se in the absorber allowing to improve the J sc by decreasing the bandgap. But it has also been shown to improve the radiative efficiencies within the absorber [START_REF] Kuciauskas | Radiative Efficiency and Charge-Carrier Lifetimes and Diffusion Length in Polycrystalline CdSeTe Heterostructures[END_REF][54] [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]. Despite efficiencies improvements for CdTe polycrystalline thin films, they are still lagging behind Si crystalline cell efficiencies. One explanation for this is the V oc . Indeed, for 20 years the V oc have been limited under the 900meV value for polycrystalline CdTe [START_REF] Gloeckler | CdTe Solar Cells at the Threshold to 20% Efficiency[END_REF]. When calculating the V oc deficit (Figure 3.2), defined as V oc,def =E g -V oc , for several materials we can observe that the CdTe is clearly far from its theoretical limits (red line in Figure 3.2) compared to the other materials. For instance, the V oc,def for the GaAs is 280meV and for CdTe it is 630meV. If we succeed in reaching the same V oc,def for the CdTe, considering the same Jsc and FF than the one of the record cell [67], its efficiency will be improved from 21.0% up to 29.5%. The record efficiency for GaAs thin film is of 29.1%. This simple calcul shows that th CdTe mainly differs from high efficiency cells by a high V oc,deficit , Figure 3.2: V oc deficit as a function of the bandgap for the record solar cells reported in [67]. The continous red line represents the calculated Voc deficit when taking into account only radiative losses [START_REF] Brandt | Temperature-and intensity-dependent photovoltaic measurements to identify dominant recombination pathways[END_REF][175] [START_REF] Brandt | Determining interface properties limiting open-circuit voltage in heterojunction solar cells[END_REF] V oc for CdTe polycrystalline cells have been limited to <900meV. As presented in Figure 3.2, CdTe suffers from a very high V oc deficit compared to GaAs, Si or CIGS. The road towards 60 3.1. Background and state of the art of CdTe higher efficiencies is now driven by the increase in V oc [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1 V barrier[END_REF]. This V oc deficit can be explained by many factors. For clarity we divided these factors in three families separated in space (see Figure 3.3):

V oc deficit in polycrystalline CdTe

• Grain interiors: The grain interiors describe the bulk of the absorber. The CdTe bulk properties are characterized by the doping level and the minority carrier lifetime. The most common approach to dope CdTe semiconductors is to substitute Cd with type I. This leads to states at approximately 50-200 meV above the valence band [START_REF] Kuciauskas | The impact of Cu on recombination in high voltage CdTe solar cells[END_REF]. The Cu dopant is introduced in the film via a post deposition treatment. Despite a 10 -18 cm -3 dopant concentration in the film, classic hole concentrations are about 10 13 cm -3 -10 15 cm -3 [50], this can be explained via compensation with Cl defects [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1 V barrier[END_REF][START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]. Researches have been conducted to improve the hole density in CdTe by switching from Cu to group V dopants (As,P,N,Sb) substitution a Te atom. But group V incorporation needs to be achieved during the growth (low diffusion of the group V atoms within the layer [START_REF] Kuciauskas | The impact of Cu on recombination in high voltage CdTe solar cells[END_REF]). Using vapor-transport-deposition (VTD), instead of the classic CSS process, allowed producing highly-doped CdTe layers (10 17 cm -3 ) with group V dopants and a 20.8% efficiency [START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]. Minority carriers in CdTe has been limited to tens of nanoseconds. Indeed a high defective bulk material is obtained due to the incorporation of a high defects density for high doping or inherent defects. But, re-crystallization process has shown to be able to increase grain crystallinity [START_REF] Mendoza-Alvarez | Luminescence and particle size in microcrystalline CdTe thin films[END_REF] and promising group V dopants [START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF], with higher activation rates, already shown to succeed in increasing minority carrier lifetimes by orders of magnitudes [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1 V barrier[END_REF][START_REF] Metzger | CdCl2 treatment, S diffusion, and recombination in polycrystalline CdTe[END_REF]. Besides, the add of Se within the absorber has already shown to be very promising for bulk defects passivation [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF].

• Grain boundaries: Grain boundaries (GBs) are highly defective areas that can enhance non-radiative recombination and reduce the power-conversion efficiency of polycrystalline solar cells. Indeed, the frontier between two mis-oriented grain results in defective areas [START_REF] John | The electrical properties of polycrystalline silicon films[END_REF]. The mis-orientation between grains will result in dangling bounds, traps or other types of defects which will result in the creation of deep and shallow defects [START_REF] Moseley | Recombination by grain-boundary type in CdTe[END_REF]. As the grain size is of the same order of magnitude than the diffusion length, the control of grain boundaries recombination velocity is of great importance. Kelvin probe force microscopy [START_REF] Jiang | Simultaneous Examination of Grain-Boundary Potential, Recombination, and Photocurrent in CdTe Solar Cells Using Diverse Nanometer-Scale Imaging[END_REF], electron-beam-induced current [START_REF] Galloway | Characterisation of thin film CdS/CdTe solar cells using electron and optical beam induced current[END_REF][START_REF] Poplawsky | Direct Imaging of Cl-and Cu-Induced Short-Circuit Efficiency Changes in CdTe Solar Cells[END_REF][START_REF] Poplawsky | Nanoscale doping profiles within CdTe grain boundaries and at the CdS/CdTe interface revealed by atom probe tomography and STEM EBIC[END_REF], and other microscopic measurements [START_REF] Cohen | Unusually low surface recombination and long bulk lifetime in n -CdTe single crystals[END_REF] have shown that GBs can develop electrostatic potentials of ∼50-100 mV that attract and channel minority carriers to the pn junction. According to numerical simulations, these potentials can give small increases in current collection, but they also increase recombination in forward bias and reduce the open-circuit voltage (V oc ). The impact on efficiency is predicted to be negative for GB recombination velocity and potential values that are consistent with measurements. The grain size is a critical parameter determining the magnitude of GB performance impacts. For example, the V oc is predicted to vary as the natural logarithm of the grain size (assuming columnar grains that are not fully depleted) [START_REF] Jensen | Long carrier lifetimes in large-grain polycrystalline CdTe without CdCl2[END_REF][START_REF] Matthew | Characterization of Sputtered CdTe Thin Films with Electron Backscatter Diffraction and Correlation with Device Performance[END_REF]. The potential to improve V oc through increasing grain size was demonstrated by monocrystalline CdTe solar cells [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1 V barrier[END_REF][START_REF] Burst | Interface Characterization of Single-Crystal CdTe Solar Cells With VOC > 950 mV[END_REF][START_REF] Cohen | Unusually low surface recombination and long bulk lifetime in n -CdTe single crystals[END_REF]. These cells achieved V oc >1 V, which is substantially above the highest reported value for polycrystalline cells [95,[START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]. Unfortunately, the epitaxial and bulk-crystal growth methods used for these monocrystalline cells are not compatible with rapid and large-scale manufacturing.

• Interfaces: The most common front contact which has been used with CdTe is the CdS, which is known to alloy with CdTe during processing. Recent total replacement of CdS with MZO window layer has enabled to greatly reduce parasitic absorption (with a bandgap of 3.75 eV) [START_REF] Kephart | Sputter-Deposited Oxides for Interface Passivation of CdTe Photovoltaics[END_REF].

But, Burst et al. [START_REF] Burst | Interface Characterization of Single-Crystal CdTe Solar Cells With VOC > 950 mV[END_REF] have shown with modeling that even if the minority carrier lifetime is successfully increased from 10ns to 100ns, the V oc can be limited to <900mV for CdTe. This is explained by a very defective buffer/absorber interface with a high recombination velocity S>10 5 cm.s -1 .

Recent improvements in efficiencies have been performed by adding a grading of Se in the absorber, with high Se concentration near the front contact [START_REF] Amit | Polycrystalline CdSeTe/CdTe Absorber Cells With 28 mA/cm 2 Short-Circuit Current[END_REF]. Adding Se in the CdTe absorber has been shown to modify the bandgap [START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF]. One reason which has been proposed for the higher carrier lifetimes is the reduction of interfacial recombination because Se improves band alignments between the absorber and buffer layers. An other reason is the bandgap grading which is now possible with Se. Indeed, the wider-bandgap non-alloyed CdTe at the back of the device acts as an electron reflector and reduces backsurface recombination.

The V oc is the main limiting factor for the efficiency improvements. We showed that, compared to GaAs, CdTe suffers for a very high V oc deficit. In Subsection 3.1.1 we showed that by reaching the same level of the V oc,def than the GaAs technology, the CdTe polycrystalline thin-films efficiency may be increased from 21.0% up to 29.5%. In Subsection3.1.2, the reasons for the poor V oc have been described by dividing the causes between three families: the grain interiors, the grain boundaries and the interfaces.

Recombinations at interfaces in polycrystalline materials

Non-radiative recombinations occur in defective areas, such as structural defects or interfaces. As a polycyrstalline material, CdTe presents different types of interfaces: Front/back surfaces and grain boundaries. In the following Subsection 3.1.3.1, I explain the impact of the different interfaces on minority carriers lifetimes. Besides, the room for improvement for reduction of their detrimental impact with surface passivation using post deposition annealing (CdCl 2 ) will be developed.

Role of interfaces in polycrystalline materials

Grain boundaries: Recently published studies report significant increases in carrier lifetimes and doping levels for both polycrystalline and monocrystalline CdTe [START_REF] Burst | CdTe solar cells with open-circuit voltage breaking the 1 V barrier[END_REF][START_REF] Mccandless | Overcoming Carrier Concentration Limits in Polycrystalline CdTe Thin Films with In Situ Doping[END_REF][START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]. As bulk properties have been improved, interfaces (grain boundaries, back and front surfaces) start limiting the performances. Surface passivation is now a pathway for lifetime improvement. For a bulk lifetime τ bulk =10ns, the diffusion length can be estimated to L diff =2.8µm (assuming D= 7.84×10 -4 m 2 .s -1 [START_REF] Kanevce | Quantitative Determination of Grain-Boundary Recombination Velocity in CdTe by Cathodoluminescence Measurements and Numerical Simulations[END_REF]). It means that if the grain size is 1µm, grain boundaries recombinations should prevail against grain interiors lifetime. But CdTe suffers from a high grain boundary recombination velocity S GB >10 5 cm.s -1 . The detrimental effect of grain boundaries on the efficiency is summarized in Figure 3.4 where theoretical efficiencies have been calculated as a function of the grain boundary recombination velocity (S GB ), for different grain size (1 µm (green), 3 µm (red), 20 µm (black)), lifetimes (1 ns (circles), 10 ns (squares), and 100 ns(triangles) and interfaces recombination velocities (S=100 cm.s -1 (solid line), S=10 5 cm.s -1 (dashed line)).

As we can observe in Figure 3.4, for grains of 1µm and 3µm, efficiencies are very sensitive to the grain boundaries even for limited lifetimes (τ bulk = 1 ns). Efficiencies are clearly improved when S GB is reduced by an order of magnitude (see Subsubsection 3.1.3.2). During the past decade, grain boundary recombinations have been lowered using post deposition treatments (PDT). For CdTe, chlorine treatments have been explored to passivate the grain boundaries [81, To compare, S int =10 5 cm.s -1 and 20µm grains is shown with dashed lines. This is Figure is reprinted from [START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grainboundary recombination for 25% efficient CdTe solar cells[END_REF] 109]. In particular, CdCl 2 annealing treatments have been proved to reduce S GB [START_REF] Kanevce | Quantitative Determination of Grain-Boundary Recombination Velocity in CdTe by Cathodoluminescence Measurements and Numerical Simulations[END_REF][START_REF] Moseley | Luminescence methodology to determine grain-boundary, graininterior, and surface recombination in thin-film solar cells[END_REF] and is now widely used.

Back and front interfaces: For CdTe solar cells, an unpassivated back surface presents a high recombination velocity S=10 5 cm.s -1 . Assuming this back surface recombination velocity, efficiency is expected to be limited under 22% [START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grainboundary recombination for 25% efficient CdTe solar cells[END_REF]. Theoretical impact of the front surface recombination velocity (S Front ) and the back surface recombination velocity (S b ) are plotted in Figure 3.5. We can observe an increase in efficiency with the bulk lifetime. But for both doping levels, for long lifetimes, efficiencies reached a plateau. This plateau is a consequence of the fact that properties are driven after τ >4×10 2 ns by interfaces properties (S front ,S back ). Realization of a bandgap grading for back surface reflection of minority and absorption opti-mization in CdTe has been done by adding Se in the absorber. Indeed, alloying with Se induces a reduction in the bandgap of CdSeTe, despite a higher bandgap for CdSe (E g (CdSe)=1.70eV), thanks to a bowing effect [START_REF] Lane | A review of the optical band gap of thin film CdSxTe1-x[END_REF][START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF]. This result allows managing the bandgap within the absorber which can induce a back surface field (BSF). Partly thanks to this bandgap grading technology efficiencies have been improved up to 22.1% [START_REF] Zheng | Roles of bandgrading, lifetime, band alignment, and carrier concentration in high-efficiency CdSeTe solar cells[END_REF].

In order to increase the minority carriers lifetime, attention must be paid to the reduction of the detrimental impact of interfaces. Grain size, grain interior lifetime, grain boundary surface recombination velocity or back/front surfaces recombination are relevant properties to characterize material improvements.

Post deposition CdCl 2 annealing

Jensen et al. [START_REF] Jensen | Long carrier lifetimes in large-grain polycrystalline CdTe without CdCl2[END_REF] achieved long lifetimes in CdTe by increasing the grain size, from few microns to hundreds of microns, in absence of chlorine treatment. This result shows the high dependence of lifetimes on grain size and so grain boundaries. It also highlights the necessity to passivate grain boundaries. [START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grainboundary recombination for 25% efficient CdTe solar cells[END_REF] For CdTe various chlorine treatments [START_REF] Major | In-depth analysis of chloride treatments for thin-film CdTe solar cells[END_REF] have been tried but CdCl 2 has emerged as the most efficient post deposition treatment (PDT). The exact role of CdCl 2 is still under debate. CdCl 2 PDT has various effects. The first one is to increase the grain size [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][START_REF] Matthew | Characterization of Sputtered CdTe Thin Films with Electron Backscatter Diffraction and Correlation with Device Performance[END_REF] by recrystallisation. A second effect of CdCl 2 PDT is to passivate grain boundaries [START_REF] Moseley | Recombination by grain-boundary type in CdTe[END_REF]. Cl ions segregate at the grain boundaries [START_REF] Li | Grain-Boundary-Enhanced Carrier Collection in CdTe Solar Cells[END_REF][START_REF] Poplawsky | Nanoscale doping profiles within CdTe grain boundaries and at the CdS/CdTe interface revealed by atom probe tomography and STEM EBIC[END_REF] and substitute to Cd vacancies. This process allows substituting a deep level state with a state close to the valance band [START_REF] Li | Grain-Boundary-Enhanced Carrier Collection in CdTe Solar Cells[END_REF]. It has also been suggested theoretically that Cl and Cu cause a co-passivation of the GB [START_REF] Zhang | Effect of Copassivation of Cl and Cu on CdTe Grain Boundaries[END_REF]. Figure 3.6 a) presents cross-section EBSD images of CdTe thin films where the CdCl 2 annealing temperature has been varied. We can observe that for high CdCl 2 temperature columnar grains can be obtained, removing GB along the z-axis. Using high CdCl 2 temperatures high grain sizes can be reached regardless of the deposition temperature and the window layer thickness and nature [6]. Indeed, window layers can critically affect key material properties such as grain morphology, but aggressive post-deposition enables enhancing the grain size by recrystallizing polycrystalline CdTe films.

Using numerical simulation, the impact of both the grain size and the S GB has been plotted in Figure 3.6. Studies suggest that using CdCl 2 allows reducing S GB down to few 10 5 cm.s -1 [START_REF] Kanevce | Quantitative Determination of Grain-Boundary Recombination Velocity in CdTe by Cathodoluminescence Measurements and Numerical Simulations[END_REF][START_REF] Moseley | Luminescence methodology to determine grain-boundary, graininterior, and surface recombination in thin-film solar cells[END_REF]. By increasing the grain size up to 10 µm V oc is expected to be above 0.98V.

To date, despite the strategy that consists of increasing the annealing CdCl 2 temperatures to get higher grain sizes, V oc peaks did not coincide with the maximum grain size. For instance, an optimal CdCl 2 annealing temperature around 390-430°C is used in state-of-the-art polycrystalline CdTe and CdSeTe solar cells [6,[START_REF] Barnard | 3D Lifetime Tomography Reveals How CdCl 2 Improves Recombination Throughout CdTe Solar Cells[END_REF][START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF][START_REF] Kumarasinghe | Effect of post deposition heat treatment on microstructure parameters, optical constants and composition of thermally evaporated CdTe thin films[END_REF][START_REF] Lee | Effects of heat treatment of vacuum evaporated CdCl2 layer on the properties of CdS/CdTe solar cells[END_REF][START_REF] Nakamura | Influence of CdCl 2 Treatment on Structural and Electrical Properties of Highly Efficient 2-µm-Thick CdS/CdTe Thin Film Solar Cells[END_REF][START_REF] Matthew | Characterization of Sputtered CdTe Thin Films with Electron Backscatter Diffraction and Correlation with Device Performance[END_REF]. These results indicate that there are other factors besides GB recombination, such as grain interior (GI) or interface defects, that limit performances in large-grain films produced by these strategies.

Post-deposition treatments (PDT) are important to improve the quality of polycrystalline photovoltaic thin films such as CdTe. We showed that the CdCl 2 annealing contributes to both the passivation of GBs and the increase of the grain sizes. Besides, increasing the annealing temperature has been found to increase greatly the grain sizes up to hundreds of microns. Despite these results, an optimal temperature has been observed during the process optimization which does not coincide with the highest grain sizes. It highlights the fact that there are other factors which have not been elucidated for the CdCl 2 annealing.

CL on polycrystalline CdTe thin films

State of the art CdTe solar cells present an average grain size of about a few microns [6,7]. For instance for a CdTe grown by close space sublimation grains are about 1 µm width [START_REF] Moutinho | Electron backscatter diffraction of CdTe thin films: Effects of CdCl2 treatment[END_REF]. Figure 3.7 presents CL and electron backscatter diffraction (EBSD) maps performed on the same sample with a top view. On CL maps we can clearly distinguish dark (low intensity) and bright (high intensity) areas. As shown in Figure 3.7, dark areas correlate with grain boundaries [START_REF] Moseley | Recombination by grain-boundary type in CdTe[END_REF][START_REF] Moutinho | Grain boundary character and recombination properties in CdTe thin films[END_REF]. Grain-boundaries correspond to highly defective areas where non-radiative recombinations are enhanced indeed. EBSD and CL maps coincide perfectly except for a slight shift, only due to an error in positioning. CL provides an excitation source, with a nanoscale spatial resolution which efficiently dissociates grain interiors from grain boundaries.

Mendis et al. [START_REF] Mendis | A contactless method for measuring the recombination velocity of an individual grain boundary in thin-film photovoltaics[END_REF] has coupled the high spatial resolution of the CL with a simplification of the Van Roosbroeck diffusion equations [START_REF] Van Roosbroeck | Injected Current Carrier Transport in a Semi-Infinite Semiconductor and the Determination of Lifetimes and Surface Recombination Velocities[END_REF] to develop a contactless method for extracting the bulk minority carrier diffusion length and reduced recombination velocity of an individual grain. This method is now commonly used for the extraction of bulk minority carrier diffusion length and reduced recombination velocity of an individual grain boundary [2,[START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF][START_REF] Mendis | A contactless method for measuring the recombination velocity of an individual grain boundary in thin-film photovoltaics[END_REF][START_REF] Moseley | Explanation of red spectral shifts at CdTe grain boundaries[END_REF]. Moutinho et al. [START_REF] Moutinho | Grain boundary character and recombination properties in CdTe thin films[END_REF] have studied the different types of grain boundaries and their impact on luminescence intensity. They also showed the very good correlation between dark regions on CL panchromatic maps and the grain boundaries determined with EBSD experiments. Moseley et al. [START_REF] Moseley | Luminescence methodology to determine grain-boundary, graininterior, and surface recombination in thin-film solar cells[END_REF] has assessed S GB and τ GI with a fit of the evolution of the grain interior CL intensity as a function of the grain sizes with a model for a large number of grains. The obtained values have been checked by introducing them in a Sentaurus device model and the simulated performances have been compared to the measured one. Besides, Moseley et al. [START_REF] Moseley | Explanation of red spectral shifts at CdTe grain boundaries[END_REF] estimated the defects density within a grain by studying the observed spectral red-shift of the CL signal. An original sample processing for CL in depth analyses has also been developed. It has first been used on CdTe samples for the investigation of the impact of CdCl 2 and Cu PDTs on the recombination at the grain boundaries through the thickness of the film [START_REF] Moseley | Cathodoluminescence spectrum imaging analysis of CdTe thin-film bevels[END_REF]. Then it has been applied on CdSeTe sample for the investigation of the evolution of the radiative efficiencies as a function of the Se content [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF][START_REF] Thomas | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF]. Further insights on this latter point are given in Section 3.3.

Conclusion of the section

We highlighted in Subsection 3.1.1 that state of the art CdTe thin films suffer from a high V oc which lower their efficiencies. Causes for these poor results have been explained in Subsection 3.1.2. Insights on the complex role played by the interfaces on the solar cells efficiencies have been given (Subsection 3.1.3.1). In order to lower the detrimental impact of interfaces, high temperature post-deposition CdCl 2 annealing has been used to both increase the grain sizes and passivate the GBs. Results of the impact of this annealing on the grain sizes have been shown in Subsection 3.1.3.2. But, the exact role of this PDT on the CdTe polycrystalline cells have been shown not to be perfectly understood. High resolution CL mapping has been shown to be able to probe the radiative efficiency within the grains (Subsection 3.1.4), allowing for both statistical and local analyses of the grains. In the following chapter we focus on the characterization of the grain interiors radiative efficiencies, the grain boundaries recombination velocities and the grain sizes based on CL mapping at both low-and room temperature.

Impact of the CdCl 2 annealing temperature on the defects passivation and formation

Variations of typical thin film processing have produced films with relatively large grains. These include increasing the temperature or duration of the post-deposition CdCl 2 annealing step. An increase of the CdCl 2 annealing temperature has been shown to be able to increase grain size near 10µm. While much is known about CdCl 2 's transformative effects -recrystallization, grain growth, and grain boundary passivation [START_REF] Lee | Effects of heat treatment of vacuum evaporated CdCl2 layer on the properties of CdS/CdTe solar cells[END_REF][START_REF] Mendoza-Alvarez | Luminescence and particle size in microcrystalline CdTe thin films[END_REF][START_REF] Metzger | CdCl2 treatment, S diffusion, and recombination in polycrystalline CdTe[END_REF][START_REF] Moutinho | Electron backscatter diffraction of CdTe thin films: Effects of CdCl2 treatment[END_REF][START_REF] Nakamura | Influence of CdCl 2 Treatment on Structural and Electrical Properties of Highly Efficient 2-µm-Thick CdS/CdTe Thin Film Solar Cells[END_REF]-much can still be learned at a more fundamental level.

In this Section 3.2, we used high resolution CL mapping to study CdTe thin films for a range of CdCl 2 annealing temperatures and grain sizes. We collected room-and low-temperature CL maps on the very same microscopic areas of samples annealed under different CdCl 2 temperatures (400°C, 420°C, 440°C, 460°C) and as-deposited. By combining local and statistical analyses, we describe the impact of the annealing temperature on the passivation and formation of defects at the grain boundaries and in the grain interiors. Our results highlight the importance of GI defects for improving the device performances.

Description of the series of CdTe thin films

TCO/MZO/CdTe superstrate thin-films were fabricated at the NREL with a standard process and various CdCl 2 post-deposition treatment temperatures. A 100nm MgZnO (MZO) film has first been deposited via RF planar magnetron sputter on a SnO 2 :F coated soda-lime glass. A 4-5µm thick CdTe was then grown on the MZO by CSS (as described in Figure 3.8) at a substrate temperature of 600°C, in an atmosphere containing O 2 (20%). Vapor CdCl 2 treatments were conducted at temperatures ranging from 400°C to 460°C in a CSS chamber for 10min in an oxygen containing ambient. The full set of annealing temperatures is reported in Table 3.1. The fabrication process was similar to the one fully described in [6]. Then CdTe top surfaces were flattened by glancing-angle ion-milling (3kV, tilting angle of 5°) with an inert Ar + in a JEOL Cross-Section polisher tool to avoid topographic artifacts during CL experiments. The sample preparation is not believed to modify significantly nor the defects nature nor the defects density. This preparation is commonly performed prior to any measurements (CL, TRCL, EBIC, TRPL...) on CdTe samples. temperatures ranging from 400°C to 460°C (Table 3.1). Based on this set of sample we will investigate the impact of the annealing temperature on the passivation and formation of defects at the grain boundaries and in the grain interiors.

Analysis of RT CL maps

Evolution of the radiative efficiency with the CdCl 2 annealing temperature

High resolution hyperspectral CL maps (128×128 pixels) were recorded at RT over a 27×27 µm 2 sample surface area. Measurements for the whole set of samples were made the same day in order to keep the injection level constant. A 30µm wide image has been chosen in order to scan a large number of grains, of at least 50 grains, but still to possess a high spatial resolution for a reasonable acquisition time. The lower bound of the spatial resolution for the CL image has been estimated to be of a = 95 nm from CASINO simulations (see Subsection 2.3.1). With the number of pixels considered the acquisition times was of 54 min.

Standard injection conditions have been used: 6 kV and I = 2 nA. A corresponding excess minority carriers have been estimated to be of ∆n≈1.6×10 16 cm -3 . The high-injection conditions used are expected to screen the built-in electric fields associated with the GBs and the surface and to minimize charge separation. The impact of the pn-junction field is supposed to be minimized as the measurement is conducted on the back surface and the sample thickness of 4-5µm is more than one order of magnitude greater than the electron beam penetration depth (R e ).

First, no clear variations in the peak position or FWHM are observable (Figure 3.10), despite the very high signal/noise ratio obtained. But we can observe an increase of the CL intensity until an annealing temperature of 420°C and then a decrease for the 460°C annealed sample. No clear local variations of the peak position, the FWHM and of the Urbach tail has been observed in CL spectra at RT for the different samples. This is why, for RT CL maps on CdTe, we focused on the analysis of the panchromatic maps.

The panchromatic CL maps measured at room temperature on the 5 samples are plotted in (Figure 3.9) (a-e) using a common linear intensity scale. The polycrystalline nature of the CdTe thin films is clear. For the as-deposited CdTe Figure 3.9 a), the CL map shows small grains and very inhomogeneous CL intensities, with stronger emission from larger grains. For CdCl 2 treated films, the average grain size increases (see Figure 3.11). This progression correlates with an improved CL homogeneity. Interestingly, the highest CL intensity is observed for an annealing temperature of 420°C, rather than the larger grains obtained at 460°C. A first effect of the CdCl 2 annealing is to increase the grain size by recrystallisation [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][START_REF] Matthew | Characterization of Sputtered CdTe Thin Films with Electron Backscatter Diffraction and Correlation with Device Performance[END_REF]. In the following we show how to estimate the average grain size from CL panchromatic maps. High resolution hyperspectral (256×256 pixels) CL maps have been acquired over a 30 µm×30 µm sample surface area. Longer acquisition are needed for this method as we need a very high resolution, a very wide field of view and a good signal/noise ratio. Acquisition has been made over several days but it is not critical as we don't need to compare quantitatively the CL intensities. Besides, we choose to perform this estimation with RT maps as the distinction between GB and GI is sharper at RT compare to LT because. This is highlighted by comparing panchromatic CL maps at RT (Figure 3.9) and LT (Figure 3.17).

Estimation of the average grain size

Based on the CL contrast between GBs and GI, we defined a CL intensity threshold to define grains. Then image processing has been performed, with the ImageJ software, in order to measure the average area occupied by grains on each map. We assume each grain to be a cylinders. This assumption is supported by cross section images of CdTe samples (Figure3.6) which show columnar grains after CdCl 2 PDTs. We determine the diameter from disc with an area corresponding to the average grain area calculated before.

Grain diameter d (µm)

No Evolution of the grains diameter with the CdCl 2 annealing temperature has been plotted in Figure 3.11. We can notice a clear increase of the average grain size with the annealing temperature from a 1.45 µm for the untreated sample up to 3.19µm for an annealing temperature of 460°C. Besides, an increase of the standard deviation of the grain size can also be noticed and attributed to the grain-coalescence mechanism.

Statistical analysis of the luminescence intensity

A statistical analysis of the CL intensities measured on each map is presented in Figure 3.12a)d). The histograms (in grey) show the pixel counts as a function of the luminescence intensity. Interestingly, they exhibit two distinct families, which we identify as GBs (low-intensity peak) and GIs (high-intensity peak). In Figure 3.12 a)-c), the histograms are accurately fit with the sum of 2-Gaussian functions (red curves) defined by the mean intensities (I GI and I GB ), standard deviations (σ GI and σ GB ), and amplitudes (N GI and N GB ) as presented in Equation 3.1:

N (I CL ) = N GB × e - (I CL -I GB ) 2 2σ GB + N GI × e - (I CL -I GI ) 2 2σ GI (3.1)
In Figure3.12d) and Figure3.12e), the GI peaks are much narrower and the two families no longer overlap. The narrowness and the height of the GI is related to the luminescence homogeneity and the increase of the grain size. The upper and lower parts of the histograms are fit separately by 1-Gaussian functions, allowing for a third family with intermediate intensities. The fit parameters are given in (Table 3.3). They provide a convenient way to assess quantitatively the impact of the CdCl 2 annealing on the CdTe luminescence properties.

The evolution of the mean CL intensities I GI and I GB is shown in (Figure 3.13). The intensity I GB increases with the annealing temperature up to 420°C, indicating a partial passivation of the non-radiative grain boundary defects [START_REF] Kanevce | Quantitative Determination of Grain-Boundary Recombination Velocity in CdTe by Cathodoluminescence Measurements and Numerical Simulations[END_REF][START_REF] Moseley | Luminescence methodology to determine grain-boundary, graininterior, and surface recombination in thin-film solar cells[END_REF]. Passivation is explained by an accumulation of Cl atoms at the GB [START_REF] Barnard | 3D Lifetime Tomography Reveals How CdCl 2 Improves Recombination Throughout CdTe Solar Cells[END_REF][START_REF] Consonni | Correlated structural reordering and dopant redistribution in annealed polycrystalline CdTe[END_REF][START_REF] Thomas | 3D Distributions of Chlorine and Sulphur Impurities in a Thin-Film Cadmium Telluride Solar Cell[END_REF][START_REF] Li | Grain-Boundary-Enhanced Carrier Collection in CdTe Solar Cells[END_REF][START_REF] Zhang | Effect of Copassivation of Cl and Cu on CdTe Grain Boundaries[END_REF], which form complex Cl Te -V Cd (shallow acceptor) instead of V Cd (deep level) [START_REF] Hofmann | Identification of the chlorine A center in CdTe[END_REF][START_REF] Persson | Anomalous Grain Boundary Physics in Polycrystalline C u I n S e 2 : The Existence of a Hole Barrier[END_REF][START_REF] Zhang | Effect of Copassivation of Cl and Cu on CdTe Grain Boundaries[END_REF]. No further improvement is seen for higher temperatures.

The same trend is found for I GI , with the brightest grains observed at 420°C. For higher annealing temperatures, the luminescence is more homogeneous in the GIs Figure 3.9d) and Figure 3.9e), which correlates with narrower intensity distributions Figure 3.12d) and Figure 3.12e), narrower error bars in Figure 3.13 and a slight decrease in I GI . The behavior is related to the appearance of regions with intermediate luminescence intensity. In Figure 3.14, the CL maps are plotted with false colors to show the spatial distribution of each family: low intensity at GBs (blue), high intensity in the GIs (green), and intermediate intensity (red). For the higher annealing temperatures (440°C and 460°C), this third family emerges in grain interiors. It is correlated with the appearance of dark spots and lines within the biggest grains which can be attributed to dislocations or twin boundaries [START_REF] Moseley | Explanation of red spectral shifts at CdTe grain boundaries[END_REF]. Overall, these results are consistent with a decrease in the defect density in grain interiors for annealing temperatures up to 420°C and the appearance of new defective regions for higher temperatures.

Local analysis: characteristic length

The decrease of the CL intensity close to GBs provides insight into the impact of the defect densities on carrier diffusion lengths. Figure 3.15 shows a detailed region of a CL map Figure 3.15 a) and a CL profile around a GB Figure 3.15 b). It can be fit with a simple one-dimensional diffusion model (Equation 3.2) developed by Mendis [START_REF] Mendis | A contactless method for measuring the recombination velocity of an individual grain boundary in thin-film photovoltaics[END_REF], based on Van Roosbroeck's derivation [173] of the steady state carrier density distribution normal to a free surface, in a semi-infinite solid that does not contain any other free surfaces. In our case, the plane generation is the CL excitation and the free surface is the grain boundary. It describes the decay of the luminescence intensity from the grain interior (x GI >0) towards the grain boundary (x=0). The model considers a characteristic length L charac and a reduced surface recombination velocity S at the grain boundary (S=s×τ GI /L charac with s the recombination velocity at GB and τ GI the intragrain lifetime):

ln(1 - I(x) I(x GI ) ) = S S + 1 - x L charac (3.2)
L charac equals the diffusion length if recombination at the top and bottom surfaces is negligible. Equation 3.2 is used to fit the decay of the luminescence close to GBs-see Figure 3.15. Extraction of at least eight CL profiles has been done for the different samples. The results obtained for the full set of samples are shown in Figure 3c). L charac is found to increase with the annealing temperature up to 420°C and to decrease for higher temperature. This is consistent with the evolution of the grain interior intensity I GI . The larger standard deviation obtained for 420°C reflects the heterogeneous CL intensity between grains observed in Figure 3.9. The longer L charac corresponds to brighter grains, with values above 0.6µm. I x GI is defined as the CL intensity in the grain where no carriers recombine at the free surface. Plateaus in intensity in the grain are needed in order to well define I x GI . But this condition was not verified in all grains for each samples. This can be explained by a small grain size compared to the diffusion length (high quality grains). This deviation from the model induces high errors in the S values, with negative values for instance. Room-temperature panchromatic CL maps show the positive effect of the annealing for temperatures up to 420°C, with an increase of the grain size, a partial passivation of grain boundary defects (increased CL intensity I GB ) and an improvement of the radiative efficiency in grain interiors (I GI ). At higher temperatures, the detrimental impact of the annealing under CdCl 2 is observed with the decreased luminescence I GI and a lower L charac , which may be due to an increased density of defects in grain interiors.

Defects nature and distribution in grains

To gain insights into the role of defects on both the luminescence and characteristic length, we have collected low-temperature (holder temperature 10 K) CL maps (128×128 pixels) on the same areas as the room-temperature maps. Same injection conditions have been used for the acquisition.

Defects in as-deposited CdTe thin-film

We first present results for the as-deposited sample in (Figure 3.16). The CL spectrum averaged over the full area is plotted in Figure 3.16. Three spectral bands are clearly identified: the excitonic band with a peak at 1.596 eV [START_REF] Halliday | A photoluminescence study of polycrystalline thin-film CdTe/CdS solar cells[END_REF], and two defect bands at lower energies. The band at 1.45-1.5eV is attributed to A-center defects induced by complexes formed between V Cd and shallow donors [START_REF] Brunthaler | An A centre in CdTe[END_REF][START_REF] Consonni | Correlated structural reordering and dopant redistribution in annealed polycrystalline CdTe[END_REF][START_REF] Halliday | A photoluminescence study of polycrystalline thin-film CdTe/CdS solar cells[END_REF][START_REF] Halliday | A study of the depth dependence of photoluminescence from thin film CdS/CdTe solar cells using bevel etched samples[END_REF][START_REF] Hofmann | Identification of the chlorine A center in CdTe[END_REF][START_REF] Kraft | Comprehensive photoluminescence study of chlorine activated polycrystalline cadmium telluride layers[END_REF][START_REF] Mendis | Long Lifetime Hole Traps at Grain Boundaries in CdTe Thin-Film Photovoltaics[END_REF][START_REF] Persson | Anomalous Grain Boundary Physics in Polycrystalline C u I n S e 2 : The Existence of a Hole Barrier[END_REF]. The presence of A-centers on the as-deposited sample may be explained by the presence of impurities such as O or Cu. The low-energy band around 1.25-1.4eV is the "Z-band" defect that is believed to be due to donor-acceptor pair recombination between shallow donors and deep acceptors (V Cd vacancies) [START_REF] Kraft | Comprehensive photoluminescence study of chlorine activated polycrystalline cadmium telluride layers[END_REF]. These defect bands have been widely studied in low-temperature photoluminescence experiments. Here, we investigate their spatial distribution over a large number of grains with high-resolution, as shown in Figure 3.16. False colors are attributed to the three spectral bands as defined in Equations 3.3-3.5, and the color indicates the predominant recombination process, either through the excitonic transition (green) or a defect level (red and blue):

I green (x, y) = I Exciton (x, y) I Full spectrum (x, y) (3.

b) False color (128x128 pixels) CL map where contributions for each spectra bands Z-defects (blue) A-defects (red) exciton (green). Contribution corresponds to ratio of the integrated intensity of the spectral band over the integrated intensity of the full spectrum.

Interestingly, A-and Z-bands are found in different grains with no or very little overlap. This is consistent with the origin of these defects related to cadmium vacancies alone (Z-band), or that form complexes with shallow donors (A-band). Only one of the two luminescence bands is observed in each grain, and in most grains, the predominant emission is uniformly distributed in the grain interior.

of the excitonic band increases monotonously with T, with a slight red-shift of the peak energy, as shown in Figure 3.17 a) and Figure 3.17 b). On the contrary, we observe a decrease of the emission through the defect band by a factor of more than two when T is increased from 400°C to T=420°C. The spatial distribution of these defects is shown in Figure 3.17 k) and Figure 3.17 l), respectively. At T=400°C, the defect density seems to be homogeneous within each grain, and we observe strong variations from one grain to another with no obvious correlation with the grain size. At T=420°C, the emission is strongly reduced and is much more homogeneous (see Figure 3.17 l)). This is consistent with the higher luminescence at room temperature, Figure 3.17 l).

For annealing at higher temperatures, the LT CL maps reveal an increase in the defect density (see Figure 3.17 m)-n) and the integrated CL intensity in Figure 3.17 b)). We observe a higher luminescence of the defect band in the grain interiors, in particular in larger grains. This phenomenon may be due to the diffusion of more Cl into grain interiors at the higher annealing temperatures, forming Cl Te defects. For the highest annealing temperature of T=460°C, the growth of large grains results in the addition of structural defects due to twin boundaries and dislocations. They are evidenced by gray regions in the LT CL maps (Figure 3.17j),n)), which can be distinguished from grain boundaries with lower luminescence. These structural defects are also responsible for regions with lower luminescence at room-temperature [START_REF] Moutinho | Grain boundary character and recombination properties in CdTe thin films[END_REF], as observed in Figure 3.17 f) and Figure 3.14 (red regions in grain interiors corresponding to the intermediate family of pixels in Figure 3.17 d),e)).

Conclusion of the section

We presented high-resolution CL maps measured at both room-and low-temperature on the same microscopic area for as-deposited samples and samples annealed at different temperatures. The results reveal some of the roles of the CdCl 2 annealing, and the mechanisms that lead to an optimal annealing temperature. Ideally, the monotonous increase of the grain size (Figure 3.11) with the annealing temperature should help decrease nonradiative recombination at the GBs and GIs, and increase solar cell performance. This is not the case. We studied the impact of the annealing at different temperatures on the nature, density and spatial distribution of the defects We showed that the CdCl 2 annealing contributes to the passivation of both GBs and GIs for temperatures up to T opt = 420°C. This corresponds to an increase of the overall luminescence and GI diffusion lengths. Beyond T opt , the grain size continues to increase, with no visible effect on the luminescence intensity at GB. In contrast, the higher annealing temperature is responsible for Cl diffusion in the GI and an increase of the bulk defect density. Hence, the optimal temperature is the result of a trade-off between GB passivation, grain size, and GI defect density.

Investigation of CdSe x Te 1-x alloys properties

Recent improvements in both modules and cells efficiencies have been made adding Se in the absorber. Alloying CdTe with Se decreases the bandgap down to 1.4 eV, thereby increasing photocurrent in the red region of the solar spectrum. Indeed, the bandgaps of CdSe x Te 1-x thin films deduced from optical characterization have been shown to decrease from E g (CdTe)=1.51eV down to E g (CdSe 0.4 T e 0.6 ) = 1.41 eV and then to sharply increase up to E g (CdSe) = 1.7eV . The quadratic evolution of the bandgap is called the bowing effect. It is attributed to a change in the crystalline structure from zinc-blende to wurtzite for x Se ≈ 0.4 [START_REF] Brill | Molecular beam epitaxial growth and characterization of Cd-based II-VI wide-bandgap compounds on Si substrates[END_REF]. It is described by Equation 3.6 [START_REF] Brill | Molecular beam epitaxial growth and characterization of Cd-based II-VI wide-bandgap compounds on Si substrates[END_REF][START_REF] Lane | A review of the optical band gap of thin film CdSxTe1-x[END_REF][START_REF] Poon | Relativistic band structure of ternary II-VI semiconductor alloys containing Cd, Zn, Se and Te[END_REF][START_REF] Swanson | Co-sublimation of CdSexTe1-x layers for CdTe solar cells[END_REF], where b is the bowing factor. The bandgap evolution of the CdSe x Te 1-x has been studied for different deposition techniques and characterization methods and bowing factors of 0.90 [START_REF] Poon | Relativistic band structure of ternary II-VI semiconductor alloys containing Cd, Zn, Se and Te[END_REF], 0.996 [START_REF] Brill | Molecular beam epitaxial growth and characterization of Cd-based II-VI wide-bandgap compounds on Si substrates[END_REF] and 1.88 [START_REF] Lane | A review of the optical band gap of thin film CdSxTe1-x[END_REF] have been found. [START_REF] Brill | Molecular beam epitaxial growth and characterization of Cd-based II-VI wide-bandgap compounds on Si substrates[END_REF] Aside from the lowering of the bandgap, higher content of Se at the front of the device creates a grading in electron affinity that helps repulsing electrons to the back contacts [START_REF] Thomas | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF][START_REF] Metzger | Exceeding 20% efficiency with in situ group V doping in polycrystalline CdTe solar cells[END_REF]. This bandgap grading, as presented in Figure 3.18, can be fabricated by the deposition of a CdSe x Te 1-x layer followed by the desposition of a CdTe layer and by a CdCl 2 annealing causing a Se diffusion within the CdTe film. This type of structures are called CdSeTe/CdTe layers. The bandgap grading therefore allows passivating back and front interfaces by repulsing minority carriers away from the interfaces. Critically, V oc does not decrease commensurately with E g . Recent studies have shown that this is associated with improved minority carrier lifetimes in the absorber up to hundreds of nanoseconds [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF][START_REF] Thomas | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF][START_REF] Kuciauskas | Radiative Efficiency and Charge-Carrier Lifetimes and Diffusion Length in Polycrystalline CdSeTe Heterostructures[END_REF][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].

In the following Section 3.3, we performed room-and low-temperature high resolution hyperspectral CL maps on CdSeTe/CdTe layers for the investigation of the bandgap grading and the impact on the Se content on the radiative efficiencies. First, the fabrication of the series on CdSeTe/CdTe samples as well as the bevels will be described in Subsubsection 3.3.2.1. Then results on the in-depth analysis made by room-temperature hyperspectral high resolution CL mapping will be shown in Subsubsection 3.3.2.2. Passivation of defects and their nature will be explored in Subsubsection 3.3.2.3 and related to cells parameters as the lifetimes and the V oc .

Prior to this, first results on the experiments conducted on CdSeTe thin films will be presented in Subsection 3.3.1.

Notable results on CdSeTe/CdTe layers

It includes an outline where the SIMS and CL measurement area have been performed (dotted black line). b) A SIMS elemental map of the selenium distribution on the beveled surface of the CdCl 2 -treated device with the 'skeletonized' chlorine signal overlaid in white to help delineate the grain boundaries. c) Panchromatic CL map taken on the same area as the SIMS selenium map. d) A high-magnification Se SIMS and panchromatic CL maps of a region at the top of the bevel. Similarity in the positioning of the selenium and CL signals is outlined by the dashed lines in d) and e). f) A scatter plot of the Se concentration versus CL counts (for equivalent regions in b) and c). Figure reprinted from [53]

Zheng et al. [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF] and Fiducia et al. [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF][START_REF] Thomas | Three-Dimensional Imaging of Selenium and Chlorine Distributions in Highly Efficient Selenium-Graded Cadmium Telluride Solar Cells[END_REF] have performed CL experiments on beveled CdSeTe/CdTe samples (see Figure 3.19). Such sample preparation allows probing the optical properties of the CdSeTe/CdTe sample along its thickness. Besides, direct comparison between CdTe and CdSeTe can be made. Fiducia et al. [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF] performed a complementary nano-secondary ion mass spectroscopy (SIMS) map (see Figure 3.19) at the same location than the CL map. The SIMS is a characterization technique which allows for probing at the nanoscale the composition by eroding the sample and then analyzing the extracted ions with a mass spectrometer. Therefore, the Se content has been correlated with the CL intensity at the nanoscale level. It has first been shown that the Se atoms diffuses through the film during the CdCl 2 annealing via the GB, which act as diffusion channels. Besides, a red-shift of the CL maximum intensity peak has been observed in the Se rich region, attesting for a lower bandgap. An increase of the radiative efficiencies with the Se content is observed, by an increase of the CL intensity, and correlated with an improvement of the diffusion lengths (estimated using CL profiles as described by Equation 3.2). This increase of the diffusion lengths is suggested by a passivation of the defects at both the GB and in the GI.

Through the use of EBSD, Auger-electron spectroscopy profile and CL Zheng et al. [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF] characterized the structural and chemical changes in the CdSeTe/CdTe compared to CdTe-only layers, and throughout the device depth. A bilayer structure in terms of Se composition and grain sizes was demonstrated in the CdSeTe sample. The CL map corroborates with Auger electron spectroscopy compositional changes and shows a four time enhancement of the CL intensity in the CdSeTe region. Besides, lifetimes in the CdSeTe/CdTe and in CdTe-only devices have been measured using TRPL. Longer lifetimes, by orders of magnitudes, have been obtained in the CdSeTe region compared to CdTe and to the CdTe-only devices.

A passivation of defects is observed with the addition of Se in the structure. But, no evidence of the nature and distribution of defects which have been passivated is given. Besides, the dependence of the passivation effects with the CdSe x Te 1-x layer composition (before CdCl 2 annealing) has not been also analyzed as much as the evolution of the grain size as a function of the Se content.

Impact of the Se composition on the bandgap and the defects passivation within CdSeTe/CdTe layers

Description of the series of CdSe x Te 1-x beveled samples

Five samples were prepared at the NREL with five different Se contents in order to study the impact of the additions of Se on the optoelectronic properties and structural changes. The layers were deposited on commercial soda lime glass coated with fluorine-doped tin oxide (SnO 2 :F) and intrinsic tin oxide (i-SnO 2 ). MgZnO was deposited on the i-SnO 2 layer by RF sputtering from a compressed target of ZnO and MgO. CdSeTe and CdTe were then evaporated sequentially without breaking vacuum. The CdSe x Te 1-x source materials had Se compositions ranging from x=0 to 0.4. The chamber base pressure was ≈10 -6 Torr. The source and substrate temperatures were approximately 680°C and 400°C, respectively. The CdSeTe layers were nominally 1 µm thick. After CdSeTe and CdTe evaporation, the samples were CdCl 2 annealed at 420°C for 10 minutes. Residual CdCl 2 was rinsed off with DI water. The samples were then dipped in a 0.1 mMol CuCl 2 solution in water and annealed in a tube furnace between 180°C and 200°C for 30 minutes. Finally, a 100-nm Au contact was deposited by evaporation.

After the sample preparation, a 20°angle between the ion beam and the surface has been made by tilting the samples. Regarding the bevel preparation, the whole sample process has been described by Moseley et al. [START_REF] Moseley | Cathodoluminescence spectrum imaging analysis of CdTe thin-film bevels[END_REF]. The milling allows exposing the whole absorber thickness. As presented in Figure 3.20 b), we can observe on the SEM image that the whole stack has been exposed thanks to the ion milling treatment. Three bevels for each sample have been prepared. CdSe x Te 1-x samples with different x Se content have been prepared at the NREL with x Se ranging from 0 to 0.4. Bevels have also been prepared (Figure 3.20) in order to perform in-depth analysis.

In depth analysis of CdSeTe thin films

High resolution hyperspectral RT CL maps (256 × 256 pixels) have been performed on the five different samples. The sample area which has been scanned has been set in order to probe the whole sample thickness. The sample thickness slightly varies between the samples so the CL maps were not exactly the same and varies from 15.2 × 15.2 µm 2 to 32 × 32 µm 2 . A standard acceleration voltage of 6 keV has been used, but a slightly higher injection current of I≈4 nA has been set. The whole acquisition time was about one hour and a half for one sample. It was not possible to put all samples at the same time on the sample holder because of their dimensions.

We aim at comparing quantitatively the CL intensity between the CdSeTe and the CdTe part of the films. To do so, we checked that the collection efficiency was not affected by the variation of height due to the bevel. The focus and the collection optics position were optimized for the top of the film, near the back contact. In Figure 3.20 we can notice grains distinctly both near the MZO front contact and near the back of the cell. This suggests that both the collection efficiency and the focus are not affected by the change in sample height. Besides, we checked that the zero order signal intensity was optimized at any point of the bevel prior to the map acquisition. Figure 3.21 shows RT CL panchromatic maps under same excitation conditions. The grayscale has been set in order to optimize the contrast for each map, but with a common scale in intensity which make the quantitative comparison possible. Because of a drift caused by charge effect induced by the electron beam scanning on insulator layers, some CL intensity maps (3.21b) d) f)) have been realigned. The realignment procedure is based on the CL intensity contrast and has been checked using the corresponding SEM image after realignment.

In all CL maps (Figure 3.21 a)-d)) we can notice a bilayer structure for all samples, the only-CdTe and the CdSe x Te 1-x /CdTe samples, where smaller grains are observed near the front interface (bottom of the CL maps). Such a bilayer structure has already been observed for CdSeTe/CdTe [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF] but also for CdTe [6,7] structures, for non-optimized CdCl 2 annealing temperatures. Smaller grains, in comparison with the CdTe region, are observed in Figure 3.21 b)-e) near the MZO interface when adding a CdSeTe layer. Besides, an increase of the grain size is observed from x Se =0.1 to x Se =0.3 samples. But for higher Se content (x Se =0.4), the grain sizes decrease.

A clear contrast between the top and the bottom of the CL intensity maps is observed on the CdSeTe/CdTe films (Figure 3.21b)-e)). Such a contrast is not seen on the only-CdTe films. Indeed, despite smaller grains, CL intensity is higher in the Se rich region. Such a contrast, between the Se rich and poor regions, has already been observed to some extends by Fiducia et al. [START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF], and explained by a passivation of the grain boundaries and the bulk defects in the Se rich region. Aside from the decrease in CL intensity within the film thickness, we can notice an increase of the CL intensity as we increase the Se content in the film. Besides, the brightest grains are observed in Figure 3.21 e). From the grayscale bars, we can notice that the CL intensity has been increased by a factor ≈10 from the only-CdTe sample to the CdSe 0.4 Te 0.6 film. The increase of the CL intensity, despite the reduction of the grain size, for the x Se =0.4 attests for an efficient passivation of defects in the grain interiors or/and at the grain boundaries as suggested before.

Thirty average spectra have been determined by averaging over the y-direction (horizontal, see Figure 3.20). For each spectrum the maximum CL intensity and the peak energy position have been assessed. Profiles of the evolution of the CL intensity and the peak energy position have been plotted in Figure 3.22. The abscissa in Figure 3.22 corresponds to the approximate distance (y-axis) normal to the film thickness (z-axis). The origin of the distance abscissa (dis-tance=0 µm) has been set to be at the MZO/CdSeTe interface. The same trend is observed for all the CL intensity profiles: the CL intensity increases as we move away from the MZO/CdSeTe interface, and then decreases as we get closer to the back interface. As already emphasized in RT CL maps, the CL intensity is higher in the CdSeTe part (orange) than in the CdTe part (blue) of the film. Besides, as we increase the Se content, from x=0 to x=0.4, the CL intensity increases. We can also notice a slight increase of the mean CL intensity in the CdTe region (blue) when Se is added to the structure, compared to the only-CdTe sample. This behavior may be explained by the diffusion of Se ions along the GB during the CdCl 2 annealing thus passivating the grain boundaries [START_REF] Zheng | Recombination and bandgap engineering in CdSeTe/CdTe solar cells[END_REF]. Despite the PDT annealing the transition between the two parts of the film is abrupt. After 4 µm, all CL intensities decrease and reach a plateau after a few microns. This effect is also observed in the evolution of the peak position as a function of the distance (Figure 3.22 b) ). For the x = 0.3 and x = 0.4 samples, the energy peak position plateau in the CdTe region reaches higher values. This behavior is not fully understood. Far from the CdSeTe part of the film (blue area), the peak energy reaches a plateau at 1.51eV.

Both peak positions and FWHMs of the CL peaks in the CdSeTe regions have been summarized in Table 3.5. We can notice a red-shift of the peak position as we increase the Se content and a slight decrease in the FWHM. As a first order approximation we can associate the CL peak position to the bandgap. We plot the evolution of the CL peak position with the Se content. We approximate that, despite Se diffusion, the Se content in the CdSeTe region was unchanged. A fit of the CL peak position of the CdSeTe with Equation 3.6 is shown in Figure 3.23, and a bowing factor of b = 0.68 is found.

We observed a red-shift of the peak position as a function of x Se . This is consistent with the fact that we expected a red-shift of the CdSe x Te 1-x bandgap from x=0 to x=0. 4 underestimation of the bowing factor may be partly explained by the overestimation of the bandgap that we made by associating it to the CL peak position. Still, the estimation of the Se content from the position of the CL peak at RT seems to be a good approximation as the FWHMs vary only slightly with the composition. So, if there is an error on the Se composition estimation, from the CL peak position approximation, it would be constant.

In-depth analysis has been performed by RT high resolution CL maps on beveled CdSe x Te 1-x samples with x ranging from 0 to 0.4. Small and bright grains are observed in the CdSeTe region. The small grain size is similar to the bilayer structures already observed for CdTe layer where the CdCl 2 annealing process was not optimized. But the fact that, despite the small grain sizes, a higher CL intensity has been observed in the Se rich region for all samples attests for greater radiative efficiencies. This greater radiative efficiency can be explained by a passivation of grain boundaries and bulk defects. Besides, an estimation of the Se content with the CL peak position has been shown to be a approximation (Figure 3.23).

Defects passivation with Se in CdSeTe layers

In the previous Subsubsection 3.3.2.2, we discussed for the evolution of the CL intensity and peak position within the films and between the samples. We conclude on the increase of the CL intensity and a red-shift of the CL peak position with the Se content. In order to go a step further we plotted the evolution of the RT CL intensity with CL peak position in Figure 3.25, for CdSeTe/CdTe samples. The set of data which have been used have been determined from RT CL maps (Figure 3.21). For each photon energy, an average CL intensity has been determined for each sample. Indeed, we showed in Figure 3.23 that the CL peak position is closely linked to the Se content through the bandgap evolution of CdSe x Te 1-x with x. In Figure 3.24 we can notice that, as the peak position is red-shifted (i.e. the Se proportion is increased) the CL intensity is increased by more than one order of magnitude between 1.41 eV and 1.51 eV. Besides, a common slope can be observed for the evolution of the CL intensity as a function of the CL peak position. This evolution is independent of the grain size and of the sample. A sharp decrease of the CL intensity is observed for the lowest energy points for each sample. This behavior can be attributed to the fact that these points correspond to the part of the film near the front interface where more non-radiative recombinations occur.

In order to connect the evolution of radiative efficiencies with the presence of defects, LT CL maps have been performed. LT hyperspectral CL maps (256×256 pixels) have been recorded for each sample on the same bevels, as the ones already analyzed with RT study, and same injection conditions have been used (6kV, I=4nA). Higher intensities have been obtained at LT, allowing to decrease the acquisition time. As the acquisition time was greatly reduced (from 100ms down to 5ms), no charge effects was observed and so no realignment of the CL map was performed.

Average CL spectra were determined by averaging over the 256×256 pixels of the LT CL maps. The results are shown in Figure 3.25 As expected, a characteristic CdTe LT spectrum is found for the Se free sample (Figure 3.25 a)). The CL spectrum is composed of a sharp excitonic peak (1.6 eV) and a broad low energy band corresponding to the Z-band defects, as described in Section 3.2.3. Adding Se in the structure induces the apparition of a second sharp peak with a lower energy, which is attributed to CdSeTe band edge transition. Furthermore, as we increase the Se content the excitonic peak red-shifts (Figure 3.25 b)-e)). Besides we can

Spectra are extracted from LT CL maps (256×256pixels). f)-o) CL false color maps (256×256 pixels) at low-temperature (20K) where normalized luminescence for each spectra bands defects (blue) CdSeTe-exciton (red) CdTe-exciton (green) are plotted. CdSeTe-exciton and CdTe-exciton maps were superposed for a better comparison of the change in composition within the film. Normalization factors have been determined from the maximum integrated luminescence from each spectral bands. White scale bar: 2 µm

notice a second defect broad band which first overlaps with the CdTe related defect band and then starts dissociating at lower energies for higher Se content.

In order to analyze the distribution of each band, integrated intensities over each spectral region were calculated at each pixel and attributed to a color: defects (blue), CdSeTe exciton (red) and CdTe exciton (green). The maximum integrated intensities have been determined for each spectral bands from all samples, then false color CL maps were normalized by these factors. Due to a red-shift of the peak, the spectral band for the CdSeTe excitonic emission was adjusted for each sample. The CL maps of the defect-related transition have been plotted in Figure 3.25 f)-j) and the scale has been adjusted for each map. The CL maps of the excitonic emissions have been superimposed and plotted for each sample in Figure 3.25k)-o) and the scale has been adjusted for each map and for each excitonic emission.

We can first notice a continuous increase of the CdSeTe excitonic intensity (red) with x (Figure 3.25 l)-o)), and the CdTe excitonic intensity (blue) increases from x=0.1 to x=0.3 and then decreases for x=0.4. Besides, we can notice that the intensity is more and more homogeneous in the CdTe part. As already observed on the evolution of both the CL intensity and the peak energy position at RT, the transition is abrupt between the red and the green parts of the film. Grains are mainly red or green, the fact that some grains are both red and green can be explained by carriers diffusion or a slight Se diffusion within a grain. Despite the carriers diffusion, the transition is steep which attests for a low diffusion of Se within grains but also of a low diffusion length (at LT) and a high spatial resolution. We can notice that the defect band intensity becomes more and more intense and homogeneous within the film from x= 0.1 to x= 0.3 and then decreases for x=0.4. For all samples, the highest defect band intensity is observed at the interfaces between the CdTe and the CdSeTe part of the film. No clear correlation between the low radiative efficiencies at RT and LT CL maps of the defect band has been observed. We observed an increase of the RT CL intensity as a function of the CL peak position. However we did not observe any clear correlation between this increase of the radiative efficiency and local variation of the defects distribution. But, a clear distinction between the different excitonic emissions (red and green) have been made attesting for a low Se diffusion within grains.

TRPL measurements have been performed on the similar CdSeTe/CdTe and only-CdTe samples at the NREL (Figure 3.26). The time decays were measured by time-correlated single photon counting using a laser pulse at 640 nm incident through the glass and absorbed primarily in the CdSeTe absorber region. Time decays were fitted with a two-exponential function. Only the second component has been reported in Figure 3.26 b). Figure 3.26 a) illustrates a representative curve with very long lifetimes exceeding 250 ns. Figure 3.26 b) illustrates that the open-circuit voltage is correlated to the logarithm of the measured carrier lifetime. This is consistent with the fact that the excess minority carriers quantity are linked to the lifetimes by ∆n = G × τ , where G is the generation rate. First, we can notice that the V oc of the samples which were analyzed are state of the art V oc values. It also shows that the defects passivation that we observed for high Se content is consistent with the V oc and so the cells performances. Despite all that, we observed higher RT CL intensities for the x=0.4 and yet the estimated lifetimes and the V oc values are not improved and even slightly decreases. This behavior outlines the fact that there are phenomena which are not fully understood for the defects passivation with Se in the CdSdTe/CdTe cells. Higher RT CL intensities and lower defect band emission CL intensities are observed for x=0.4 but are not related with higher V oc values. Besides, the short lifetime behavior has not been explored in this study.

Conclusion of the Section

Samples with different Se composition (x Se =0.1,0.2,0.3,0.4) and a only-CdTe sample have been fabricated with the same process. Besides, bevels have been prepared on these samples allowing for in-depth investigations. High-resolution RT hyperspectral CL maps have been performed on CdSeTe/CdTe and only-CdTe for quantitative comparison of the CL intensities. Panchromatic CL maps showed that small grains are obtained in the CdSeTe region, near the front contact. Despite the small grain size, brighter grains are obtained in this region with a higher Se content. Then, evolution of the CL spectra over the thickness have been studied with CL intensity and peak energy profiles. Besides, the CL peak position has been related to the Se content in Figure 3.23. We conclude on the improvement of the radiative efficiency with the Se content. In Subsubsection 3.3.2.3, we analyzed the evolution of the RT CL intensity as a function of the Se content. A passivation of the bulk defects has been suggested in Figure 3.24. Complementary LT hyperspectral CL maps over the same surface than the RT maps have been recorded. No clear correlation between the evolution of the radiative efficiency observed and the defects distribution within the films have been observed. TRPL experiments have been performed on similar CdSeTe/CdTe and CdTe samples to determine the lifetimes. An improvement of the V oc and the minority carriers lifetimes is observed when increasing Se from x=0 to x=0.3. Then, no improvement in the V oc and even a slight decrease can be noticed for x=0.4, despite higher CL intensities. The nature of the defects which are passivated and a quantitative determination of τ GI and S GB are still under study.

Lifetime assessment in CdTe grains

Lifetimes measurements is a key point for materials improvement as it is directly linked to the cells performances through Voc. Indeed a logarithmic evolution of the V oc as a function of the minority carrier lifetime is predicted and observed [START_REF] Chantana | Urbach energy of Cu(In,Ga)Se2 and Cu(In,Ga)(S,Se)2 absorbers prepared by various methods: Indicator of their quality[END_REF][START_REF] Kanevce | Impact of interface recombination on time resolved photoluminescence (TRPL) decays in CdTe solar cells (numerical simulation analysis)[END_REF][START_REF] Kuciauskas | Optical-fiber-based, time-resolved photoluminescence spectrometer for thin-film absorber characterization and analysis of TRPL data for CdS/CdTe interface[END_REF][START_REF] Repins | Correlation Between Measured Minority-Carrier Lifetime and Cu(In, Ga)Se2 Device Performance[END_REF]. Besides, theoretical models require a reliable set of parameters such as the minority carrier lifetime.

Difficulty for the lifetime estimation in polycrystalline materials relies on the fact that it is complicated to deconvolute GB, GI and interfaces effects. Kuciauskas et al [START_REF] Kuciauskas | Optical-fiber-based, time-resolved photoluminescence spectrometer for thin-film absorber characterization and analysis of TRPL data for CdS/CdTe interface[END_REF]95] demonstrated using 2PE that it is possible to perform lifetime profiles within the absorber thickness. Through the use of this technique, they demonstrate an efficient surface passivation of the CdTeSe/Al 2 O 3 interface. But due to the low spatial resolution of the excitation source, no distinction of the GB and GI is possible as only average lifetimes over multiple grains are measured.

In this Section we explore time-resolved CL (TRCL) experiments in order to measure carrier lifetimes in single grains. Based on the lifetimes assessment in several grains with various grain size, a method to quantitatively measure the surface recombination velocity S GB and the grain interior lifetime τ GI is proposed (Subsection 3.4.1). But, we will outline the fact that RT lifetimes measured on state of the art cells are too long for our acquisition system. Still, feasibility of the methodology for quantitative assessment of S GB and τ GI have been demonstrated by LT TRCL measurements in single CdTe grains in Subsection 3.4.2.

Quantitative estimation of τ bulk and S GB by effective lifetimes measurements in CdTe grains

I already showed that, based on the acquisition of a very high resolution RT CL map (Subsubsection 3.2.2.2), we are able to measure the grain size diameter. Then after this acquisition of the map, by switching into pulsed conditions we can measure time decays by focusing the electron beam spot at the center of a grain that we observed in the RT map. We can assume that the drift of the stage (v drif t =250nm/hr at LT) is low compared to the grain diameter (d≈1µm) so that the positioning of the stage during the RT CL map, recorded before switching into pulsed, still corresponds to the actual stage positioning during the time decay acquisition. Such an assumption is required as a precise positioning of the SEM spot only made using the SEM image under pulsed conditions is difficult because of a very noisy image.

Effective lifetime can be separated between two terms, a grain interior lifetime τ GI and a lifetime associated to grain boundaries states τ GB , as described in Equation 3.7.

1 τ eff = 1 τ GI + 1 τ GB (3.7)
A grain can be described as a cylinder with a diameter d, a surface recombination velocity (τ GB ) and a interior lifetimes (τ GI ), as presented in Figure 3.27 a). A model has been developed by Dan et al. [40] to quantify the effect of surface recombination velocity in a cylinder (Equation 3.8). We assumed that no carriers go through the grain boundary. As shown in Figure 3.16, even at low temperature it seems that there is no intermixing of the luminescence of the grains.

1 τ eff = 1 τ GI + 4 × S GB d (3.8)
Theoretical (Equation 3.8) effective lifetimes as a function of both the grain diameter and the grain interior lifetime (τ eff ) for two grain boundary recombination velocity (S GB = 10 4 cm.s -1 and S GB = 10 5 cm.s -1 ) have been plotted in Figure 3.27 b)c). Orders of magnitude for the effective lifetimes which have been determined for CdTe grains are consistent with lifetimes measurements presented in the literature for CdTe [START_REF] Kanevce | Impact of interface recombination on time resolved photoluminescence (TRPL) decays in CdTe solar cells (numerical simulation analysis)[END_REF][START_REF] Kuciauskas | Optical-fiber-based, time-resolved photoluminescence spectrometer for thin-film absorber characterization and analysis of TRPL data for CdS/CdTe interface[END_REF].

We can observe that, for S GB =10 5 cm.s -1 and τ GI ≈100ns, effective lifetime is limited to 4.5 ns even for high grain sizes (Figure 3.27c)). This lifetime leads to a L dif f <2µm, assuming D=7.84cm.s -1 , which is lower than the film thickness of about 4-5µm. But, by decreasing the S GB down to 10 4 cm.s -1 , the lifetimes are greatly increased up to 30ns. Such lifetimes allow to reach L dif f >5µm which are sufficient in the case of a film thickness of 5µm. Equation 3.8 is commonly used to determine τ GI and S GB in nanowires [30,40,47] by fitting the evolution of the τ ef f as a function of the nanowire diameter. By fitting with Equation 3.8 the evolution of the effective lifetime as a function of the grain diameter an estimation of τ GI and S GB is possible. Average values for these two crucial characteristic will be estimated. We observed in RT panchromatic CL maps (Figure 3.9) relatively homogeneous CL intensities in GI and at GB, characterized by σ GI and σ GB , for CdCl 2 annealed samples. This result support the fact that an estimation of the average τ GI and S GB is consistent.

Proof of feasibility: low temperature lifetime measurement in a CdTe grain

Minority carriers lifetimes in the order of tens to hundreds of nanoseconds have been measured in CdSeTe/CdTe cells at RT (Figure 3.26). In concrete terms, we can observe that the signal (Figure 3.26 a)) decreases by one order of magnitude after about 50 ns. But, as described in Subsection 2.2, the repetition rate of our LASER is of one pulse each 12 ns. This means that we expect not to be able to discern a time-decay with our system as the temporal window of 12 ns is not wide enough to see a decrease of the signal. But by implementing a pulse peaker, which allows to reduce the repetition rate by deflecting LASER pulses, we may be able to measure lifetimes. However, by reducing the number of pulses both the CL intensity and the SEM signals are expected to be reduced. Assuming that we want to measure a decay as the one showed in Figure 3.26 a) with a one order of magnitude signal/noise ratio for a CdSeTe/CdTe sample, with a similar time decays as the one presented in Figure 3.26 a), we will need to peak only one over five LASER pulses and so reducing the signals of a factor 5. This important fact may have a very negative impact on the TRCL measurements.

Nevertheless, in Subsection 3.4.2 I wanted to show that despite this technical issue we are able to measure lifetimes in CdTe single grains of the as-deposited sample (see Subsubsection 3.1). To do so we first measured a RT CL map under continuous excitation. This first step is necessary in order to localize grains and estimate their diameter afterwards (Figure 3.28 a)).

Then we placed the sample under low-temperature (20 K) conditions and a hyperspectral LT CL map (256×256 pixels) has been recorded over the same area than the RT CL map in order to recognize the grains. A LT CL map (50×50 pixels) under pulsed conditions has also been performed and the normalized average spectra, over the whole maps, are compared in Figure 3.29 b). Under pulsed excitation, the defects peak is higher than the excitonic peak. This is not the case for the steady-state measurements, probably because of the very different injection levels (and the saturation of defects in the continuous mode).

Time decays were measured for the exciton and the defects transitions by acquiring streak images (Figure 3.4) with different central wavelength. Time decays have then been determined by integrating on a spectral ranges in streak images. The spectral ranges I considered for the decays are displayed in Figure 3 Almost no time decay is observed for the defects, this phenomenon is explained by a long lifetime (τ def >12 ns) which has already been measured with TRPL at 80 K to be about τ = 25 ns [START_REF] Alberi | Measuring long-range carrier diffusion across multiple grains in polycrystalline semiconductors by photoluminescence imaging[END_REF]. The time-decay observed for the exciton has been fitted with a single-exponential model. A lifetime of τ =35 ps has been measured at 20K. Short excitonic lifetimes have already been measured in polycrystalline CdTe [START_REF] Alberi | Measuring long-range carrier diffusion across multiple grains in polycrystalline semiconductors by photoluminescence imaging[END_REF] and in monocrystalline CdTe [START_REF] Polhmann | Exciton lifetimes in CdTe/CdMnTe single quantum wells[END_REF].A very short excitonic lifetimes and a long lifetimes for defects at LT may suggest that there are unresolved fast decay mechanisms which are not fully understood in CdTe, as hopping mechanisms between traps [START_REF] Alberi | Measuring long-range carrier diffusion across multiple grains in polycrystalline semiconductors by photoluminescence imaging[END_REF]

Conclusion of the Section

TRCL measurements have been explored as a solution for the quantitative assessment of τ GI and S GB in CdTe. The protocol and methods to do so have been first explained in Subsection 3.4. But we also discussed the technological issues for RT lifetimes estimation with our TRCL setup as it is now. Indeed, lifetimes of tens to hundreds of nanoseconds are expected but the repetition rate of the LASER used in our TRCL setup is of 12 ns. We highlighted the need for a pulse peaker in order to measure a RT time decay on CdTe samples. However, we demonstrated the feasibility of the measurements of lifetimes within CdTe grains with LT TRCL. A very long lifetime τ def >12 ns for defects and a short lifetime of τ ef f =35 ps for the exciton have been found.

Summary

In this chapter, we focused on the characterization polycrystalline CdTe materials. We first reminded that polycrystalline materials suffer from a high V oc deficit which hinders the efficiencies. We also present the factors which impact the V oc values.

• In section 3.2, I analyzed a set of CdTe samples with different CdCl 2 annealing temperature (no-CdCl 2 , 400 °C, 420 °C, 440°C and 460 °C). First, high resolution hyperspectral RT CL maps have been recorded for each sample in order to study the impact of the CdCl 2 annealing temperature on CdTe grains. The brightest grains have been observed for the 420°C annealing temperature. This result does not coincide with the larger grains, which are obtained for the highest CdCl 2 temperature (460°C). To explain this behavior, both a statistical analysis of the RT CL intensities distribution (Subsubsection 3.2.2.3) and a local analysis of the evolution of the CL intensity near a GB (Subsubsection 3.2.2.4) have been performed. Low temperature (10K) hyperspectral CL map measurements have then been performed on the same surface sample areas. They have been used to evidence the defects in the GI and to correlate their density with the radiative efficiency. We showed that the CdCl 2 annealing contributes to the passivation of both GBs and GIs for temperatures up to T opt = 420 °C. Beyond T opt , the grain size continues to increase, with no visible effect on the luminescence intensity at GB. The higher annealing temperature is responsible for an increase of the bulk defect density. Hence, the optimal temperature is the result of a trade-off between GB passivation, grain size, and GI defect density.

• In section 3.3, we analyzed CdSe x Te 1-x /CdTe samples with different Se contents (x=0.1, 0.2, 0.3, 0.4) and a only-CdTe sample. RT hyperspectral CL maps have been recorded on the beveled samples in order to perform in-depth analyses. Despite the small grain sizes, brighter grains are obtained in the CdSeTe region (Figure 3.21). Besides, we noticed that the brightest grains are observed for the sample with the highest Se content (x=0.4). Despite the CdCl 2 annealing, a very sharp evolution of the CL peak position is observed, indicating a low Se diffusion. A common evolution of the CL intensity as a function of the peak position, and so the Se content, has been found for all the samples attesting for a passivation of defects with Se. Defect passivation and radiative efficiency of CdSe x Te 1-x thin films have been investigated and correlated to macroscopic measurements (Voc, lifetimes).

• In section 3.4, I explored TRCL measurements as a solution for the lifetimes assessment in single CdTe grains. By coupling lifetime measurements with grain sizes estimation (Subsubsection 3.2.2.2) we presented a method to determine quantitatively S GB and τ GI . First the model for the evolution of the effective lifetime as a function of the grain size has been explained in Subsection 3.4.1. Proof of feasibility of TRCL measurements in a single grain were presented in Subsection 3.4.2. A long lifetimes τ def >12 ps and a short lifetime of τ X =35 ps were measured for defect band transition and the exciton, respectively, in a single CdTe grain at LT using TRCL.

As described in Subsection 1.3.3, nanowire (NW) based solar cells present an interesting opportunity in order to fabricate tandem III-V on Si solar cells. But, their efficiencies are hindered by low V oc . These low V oc can be explained by the high defect density introduced during the NW growth, by a low doping level or by high surface recombination velocities.

In this Chapter I will first describe briefly (Subsection 4.1.1) the growth technique which has been used to grow NWs at C2N. During the growth, defects but also change in the crystalline phase can be introduced in the NW. Wurtzite (WZ) and Zinc-Blende (ZB) polytypism is often observed in III-V NWs. In Section 4.2, their optical properties such as the polarization, the bandgap and their lifetimes will be compared in GaAs NWs. Dopant incorporation and the doping assessment in NWs are challenging. A method has been developed in the laboratory to assess quantitatively the doping level in single NWs, based on the luminescence spectra recorded with the CL. I applied this method in Section 4.3 to determine the doping level induced by the incorporation of Te in GaAs NW, and its homogeneity along the NW. Because of a high surface/volume ratio, lifetimes in NWs are very dependent of the surface passivation. In Section 4.4, the effectiveness of surface passivation with high bandgap materials will be investigated using TRCL. Besides, the impact the possible stress induced by the passivation shell will be studied.

Background and nanowires growth

The GaAs NW characterized and analyzed in this chapter were grown by MBE on a prepatterned Si substrate using the Vapor-Liquid-Solid technique. In Subsection 4.1.1, I will describe the growth process used in order to obtain the NW arrays. In Subsection 4.1.2, I will describe relevant properties which need to be characterized for solar cells improvement. Then I will end by describing CL experiments on single NWs, and I will demonstrate that it allows probing important opto-electronical properties (bandgap, doping levels, defects with CL spectrum analysis) and the carriers dynamic (lifetimes, surface recombination velocities by TRCL) and their homogeneity along the NW at the nanoscale.

Description of the process for the GaAs nanowire growth by MBE

The oldest growth technique used in order to grow a NW is the catalyst assisted technique [START_REF] Lapierre | III-V nanowire photovoltaics: Review of design for high efficiency: III-V nanowire photovoltaics: Review of design for high efficiency[END_REF] Vapor-Liquid-Solid (VLS). This bottom-up technique is based on the use of a metal droplet, mostly Au, which contributes to collection of species and catalyze the growth by MBE or MOVPE. The supersaturation of the droplet allows the nucleation to start and then the epitaxial growth to start. But, the use of Au as a catalyst is not suitable (defects incorporation in the NW, Au diffusion on and in the Si substrate). This is why in 2008, Fontcuberta i Morral et al. [START_REF] Fontcuberta | Nucleation mechanism of gallium-assisted molecular beam epitaxy growth of gallium arsenide nanowires[END_REF] reported the growth of NWs with a Ga particle as the catalyst. Few years later this technique, named self-catalyzed growth, has been used for the growth of NW arrays [START_REF] Plissard | Gold-free growth of GaAs nanowires on silicon: arrays and polytypism[END_REF]. The main advantage of the use of a self-catalyzed process is the catalyst consumption in the MBE chamber via crystallization under As flux at the end of the NW growth. Then a shell can be grown using a VS technique on the sides. This technique has been chosen as it offers the opportunity to create a very uniform array of NWs [104].

Fabrication processes have already been fully described by De Lépinau [104], in this thesis we only aim at giving a brief description of the fabrication processes of the GaAs NWs. Ga droplets are deposited on a patterned Si substrate. This pattern can be achieved by the use of electron beam lithography. The Ga atoms will diffuse towards the holes in the SiO 2 . An array of Ga droplets is then formed as described in Figure 4.1 a). Then, the growth is initiated by providing As an Ga fluxes. Dopant are incorporated by providing a dopant flux (Te, Si, Be...) during the growth.

Using VLS, two types of structures for NWs are possible: radial or vertical p-n junction. The vertical junction corresponds to the case of a vertical superposition of the p-n junction (Figure 4.2 a)). On the other hand, radial junction (Figure 4.2 b)) corresponds to a core-shell structure where the shell is n-doped and the core is p-doped, or vice versa. Lapierre et al. has demonstrated that radial junctions are predicted to allow more efficient carrier collection [100]. Using VLS self-catalyzed process, Dastjerbi at al. [41] and Boulanger et al. [19] reached respectively 4.1% and 3.3% cells efficiency for a GaAs NW array grown by MBE on a Si substrate.

In this thesis we considered GaAs NWs grown by MBE using a self-catalyzed Vapor-Liquid-Solid growth process on a prepatterned Si substrate. Using this method, Dastjerbi at al. [41] and Boulanger et al. [19] have reached respectively 4.1% and 3.3% cell efficiencies.

Material limitations in nanowires

Low V oc values (0.45 V [41] and 0.39 V [19]) are obtained for the GaAs NW-based solar cells which have been developed so far. At C2N, a GaAs NW-based solar cell has been produced recently with 3.7% efficiency, on a prepatterned Si substrate using self-catalyzed VLS by MBE, presenting a record V oc = 0.65V . Still, this V oc value is far from its theoretical limit (record V oc = 1.1V for GaAs under 1-sun). This large room for improvement can be explained by many factors such as:

b) The radial p-n junction is characterized by both n-(orange) and p-(green) type doped materials. c) A passivation shell can be grown in order to reduce the surface recombination velocity (S) and so to increase the carriers lifetime within the NW.

• Crystal quality: GaAs NWs grown on Si substrate have the potential to form perfectly mono-crystalline GaAs absorber for PV application. However, the growth technique allows polytypism between Zinc-Blende and Wurtzite. Indeed, most III-V materials grow preferentially in the Zinc-Blende phase, but it has been shown that Wurtzite phase can also be obtained by controlling the growth parameters [41,[START_REF] Glas | Why Does Wurtzite Form in Nanowires of III-V Zinc Blende Semiconductors?[END_REF][START_REF] Hiruma | Growth and optical properties of nanometer-scale GaAs and InAs whiskers[END_REF][START_REF] Koguchi | Crystal Structure Change of GaAs and InAs Whiskers from Zinc-Blende to Wurtzite Type[END_REF]. This change in the crystal phase can be seen as an opportunity to create quantum dots within a NW [START_REF] Spirkoska | Structural and optical properties of high quality zinc-blende/wurtzite GaAs nanowire heterostructures[END_REF][START_REF] Vainorius | Observation of type-II recombination in single wurtzite/zincblende GaAs heterojunction nanowires[END_REF][START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF]. But uncontrolled phase modification can also lead to staking faults, non-radiative recombination centers and type II interfaces, which may limit lifetimes and carrier separation in the NW structure. I will investigate ZB and WZ phases optical and transport properties in a single NW in Section 4.2.

• Doping level: Most common dopants (Si [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF]34], C, Be [34,[START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF], Te [START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF]) have been shown to incorporate differently for NW and for thin-films. Indeed, for planar structures Si is commonly used as a n-type dopant in MBE leading to n-type doping higher than 5×10 18 cm -3 [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF]34,[START_REF] Ketterer | Compensation mechanism in silicon-doped gallium arsenide nanowires[END_REF]. But several mechanisms limit its incorporation or the effective n-type doping in the case of NWs: Amphoteric behavior, complexes formation and supersaturation of the Ga droplet. In order to increase the n-type doping level Te has been used and high doping levels have been measured in NWs [73]. But it has been observed that Te dopant induces also a change in the NW morphology by modifying the Ga contact angle [START_REF] Czaban | GaAs Core-Shell Nanowires for Photovoltaic Applications[END_REF][START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF]73]. In addition to the complex dopant incorporation mechanisms during the NW growth, the doping assessment in single NW is not straightforward because of the difficulty to establish metal contacts. However, new methods based on CL [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF] or Raman spectroscopy [73] have shown to be very promising for doping assessment in single NW. A study of the doping in GaAs using Te has been conducted in Section 4.3

• Surface passivation: In the case of NWs, a very high surface to volume ratio is obtained. This is why carrier lifetimes are expected to be highly dependent of the surface recombination velocity (SRV) [40]. In addition, GaAs suffers from a very high SRV (>10 5 cm.s -1 ). Assuming S≈10 5 cm.s -1 and a bulk lifetimes of τ bulk =100ns, carriers lifetimes are predicted to be below tens of picoseconds in GaAs NW with a diameter of about 300nm. Besides, Lapierre et al. [START_REF] Andrew | Analytical model of surface depletion in GaAs nanowires[END_REF] have demonstrated that poorly doped GaAs NWs (1×10 17 cm -3 ), with radii lower than 100nm are expected to be fully depleted. Surface passivation can be performed using chemical treatment [START_REF] Czaban | GaAs Core-Shell Nanowires for Photovoltaic Applications[END_REF] or by growing a higher bandgap material on the sides of the NWs [START_REF] Breuer | Suitability of Au-and Self-Assisted GaAs Nanowires for Optoelectronic Applications[END_REF]40,91,111,[START_REF] Perera | Nearly intrinsic exciton lifetimes in single twin-free GaAs????AlGaAs core-shell nanowire heterostructures[END_REF]. The wide bandgap material acts as a barrier for one type of carriers, reducing the surface recombination velocity. AlGaAs [4,91,[START_REF] Perera | Nearly intrinsic exciton lifetimes in single twin-free GaAs????AlGaAs core-shell nanowire heterostructures[END_REF] is widely used as a passivation layer for ease to obtain a lattice-matched shell but InGaP [111] has also been explored. In Section 4.4, I will investigate the surface passivation efficiency of both AlGaAs and InGaP. In order to the characterize properties of NWs along the NW length, a nanoscale spatial resolution characterization technique is needed. CL offers the opportunity to probe the optical properties of the NW with a high spatial resolution. In order to perform CL experiments, NWs are spread on a Si substrate (which is doped in order to avoid charge effects). A SEM image acquired simultaneously with the hyperspectral CL map is plotted in Figure 4.4 a). The panchromatic CL map (Figure 4.4 b)) demonstrates that the radiative efficiency can be probed in the NW with a high spatial resolution. Spectral information of the luminescence emitted at each point are also resolved using the spectrometer (see Figure 4.4 c)). This method has been used in this thesis in order to compare and assess the optical properties GaAs NWs .As we can notice in Figure 4.4, NWs present inhomogeneities in radiative properties. Using the very high spatial resolution of the CL in the pulsed excitation conditions we are able to measure lifetimes at different location of a NW. This method has been used in order to measure lifetimes in single NWs at both room and low temperature.

CL experiments on GaAs nanowires

Previous CL investigations of single GaAs NWs have been published. In particular, Spirkoska et al. [START_REF] Spirkoska | Structural and optical properties of high quality zinc-blende/wurtzite GaAs nanowire heterostructures[END_REF] have studied the emission patterns of a pure ZB GaAs NW and a GaAs NW presenting different sections with wurtzite and zinc-blende structures. A series of monochromatic CL images of several nanowires was recorded, with a STEM system, localized emission at energies between 1.51 and 1.46 eV. The appearance of low energy peaks has been interpreted as the evidence for type II band alignment at the wurtzite/zinc-blende interface. Using the growth substrate as a reference in terms of luminescence, Bolinsson et al. [START_REF] Bolinsson | GaAs/AlGaAs heterostructure nanowires studied by cathodoluminescence[END_REF] have realized a systematic investigation of the impact of the nanowire growth-temperature on the total photon emission, its variations in the emission energy and intensity along the length of the nanowires. Using a combination of CL and TRCL Liu et al. [START_REF] Liu | Spatially dependent carrier dynamics in single InGaN/GaN core-shell microrod by time-resolved cathodoluminescence[END_REF] have studied the influence of the inhomogeneities in a single InGaN/GaN NW on both radiative and SRH recombinations. The latter are studied by recording simultaneously position dependent cathodoluminescence intensities and time-decays from 4 K to 300 K. From simple assumption, they estimated the radiative and non radiative 100 lifetimes and their evolution with the temperature in a single NW. In Subsection 4.3, a method based on the evolution of the luminescence lineshape, recorded with CL, as a function of the doping level and type is presented.

In the following, high resolution hyperspectral CL mapping will be used to determine the evolution of the optical properties, doping levels and defects density along the nanowire. Direct quantitative comparison of luminescence has also been conducted between the samples to compare the radiative efficiencies. Besides innovative techniques such as CL polarimetry allowed evidencing unambiguously the different phases in a single nanowire.

Conclusion of the section

At the C2N laboratory GaAs NW based solar cell with a 3.7% efficiency has been fabricated. This result has been obtained with NWs grown by MBE on a prepatterned Si substrate with a self-catalyzed vapor-liquid-solid technique. First characterization showed that this cell presents a record V oc = 0.65 V . Despite this achievement, this value is still low compared to its theoretical limit showing room for improvements. In Subsection 4.1.2, crystal quality, low doping levels and high surface recombination velocities have been described as the main issues which need to be tackled. These issues will be investigated in the following using high resolution CL mapping for single NW characterization.

Wurtzite and Zinc-Blende polytypism in GaAs NWs

In order to improve cells efficiencies it is of great importance to grow NWs with as few defects as possible. In the last decade, many efforts have been made in order to control the NW phase and defects density. Polytipsm between ZB and WZ within NWs is made possible thanks to the high surface/volume ratio of the NW. An uncontrolled switch between WZ and ZB may have a detrimental impact on the device performances. Besides, heterostructures may be formed due to a band offset and/or a difference of the bandgap. But the WZ band structure is not perfectly known and is still under debate.

Main research topics on the characterization of WZ optical properties rely on the bandgap and the carriers dynamics. Ketterer et al. [91] compared optical properties of the WZ and the ZB GaAs NW using resonant Raman spectroscopy. A very small difference in the bandgap was found (E g (WZ)= 1.5165 eV and E g (ZB)=1.5162 eV). A strong polarization of the light absorbed by the WZ has been found, compared to the ZB. Indeed, because of a break in the crystallographic symmetry from ZB (cubic) to WZ (hexagonal) the light absorption is no longer isotropic in the case of the WZ (further explanations are given in Subsection 4.2.3). Besides, Ahtapdov et al. [4] analyzed WZ optical properties through the use of micro-PL and STEM experiments. First they are able to obtain pure WZ NW, then micro-PL experiments were performed in order to characterize the polarization of the light emitted by the WZ and the exciton lifetime was also assessed. Polarization of the light emitted has been shown to be strongly polarized perpendicularly to the NW axis for pure WZ NWs. Besides, very long lifetimes (3 ns) were observed for the WZ NW at low-temperature compared to bulk ZB.

In this Section 4.2, we will focus on the comparison of the optical properties of the WZ and the ZB in a single GaAs/AlGaAs/GaAs NW. NW with two distinct segments of the ZB and WZ were prepared at C2N using self-catalyzed VLS for a direct comparison. TEM experiements showed a high crystal quality of both ZB and WZ segments with no stacking faults (SF). The high spatial resolution of CL was used in order to compare optical properties of the two segments in single NWs. Room-and low-temperature CL maps were performed in order to investigate the band structure of the WZ phase. Complementary studies were explored to study both the polarization of light emitted and the lifetimes of the two phases (Subsection 4.2.3) using original processes.

Samples preparation

III-V compounds mostly adopt a cubic zinc blende (ZB) crystal structure under bulk or thinfilms forms. In the case of NWs, a Wurtzite phase has been observed [START_REF] Glas | Why Does Wurtzite Form in Nanowires of III-V Zinc Blende Semiconductors?[END_REF][START_REF] Hiruma | Growth and optical properties of nanometer-scale GaAs and InAs whiskers[END_REF][START_REF] Koguchi | Crystal Structure Change of GaAs and InAs Whiskers from Zinc-Blende to Wurtzite Type[END_REF]. Appearance of a new stable morphology is explained by a difference in the interface energy [START_REF] Bechstedt | Structure, energetics, and electronic states of III-V compound polytypes[END_REF]. The high surfaceto-volume ratio of the NWs leads to stabilize crystal phases with lower surface energy [START_REF] De | Addressing the Fundamental Electronic Properties of Wurtzite GaAs Nanowires by High-Field Magneto-Photoluminescence Spectroscopy[END_REF]. Glas et al. [START_REF] Glas | Why Does Wurtzite Form in Nanowires of III-V Zinc Blende Semiconductors?[END_REF] and Dastjerdi et al. [41] have shown that, when the super-saturation of the droplet is high enough, it is possible to control the phase with the contact angle of the Ga droplet β as presented in Figure 4.5. Predominance diagram as a function of the contact angle and the super saturation has been determined in [41] (see Figure 4.5 b). In this sense, Vainorius et al. [START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF] and Lehmann et al. [START_REF] Lehmann | High crystal quality wurtzite-zinc blende heterostructures in metal-organic vapor phase epitaxy-grown GaAs nanowires[END_REF] have performed crystal phase engineering with a Au catalyst in MOVPE.

The coexistence of the two phases (WZ/ZB) leads to stacking faults which has a detrimental impact on the lifetimes [41,[START_REF] Perera | Nearly intrinsic exciton lifetimes in single twin-free GaAs????AlGaAs core-shell nanowire heterostructures[END_REF][START_REF] Woo | Effect of Twinning on the Photoluminescence and Photoelectrochemical Properties of Indium Phosphide Nanowires Grown on Silicon (111)[END_REF]. Figure 4.5 shows GaAs/AlGaAs NWs (76715), grown at the C2N in an MBE using a VLS self-catalyzed process, with two distinct segments of about 800 nm and 350 nm. These segments can be distinguished by the STEM contrast in Figure 4.5c). Besides, TEM images (Figure 1.25) on the two segments show characteristic WZ and ZB diffrac-tion patterns. The change in the crystallographic phase at the top of the NW is attributed to Ga droplet re-crystallization which induces a change in the droplet contact angle and the III-V flux. A transition phase is also observed between the two segments which corresponds to the transition from the ZB stable growth conditions to the WZ ones. We don't expect a type II behavior characteristic of WZ and ZB interface [START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF].A lattice matched AlGaAS shell has been grown afterwards. This shell was grown to passivate the GaAs core (see Section 4.4).

Figure 4.5: a) Scheme of a VLS growth with the Ga droplet (blue) and the growing NW (red). Relevant parameters such as the supersaturation of the droplet (∆µ) and the contact angle β have been plotted. b) Supersaturation of the Ga droplet as a function of the contact angle. The black curve shows the predominance diagram between WZ and ZB phases, characterized by a supersaturation limit. The red horizontal arrow indicates the crystallization phase of the droplet under high group V flux. c)d) STEM images of GaAs NWs grown by VLS with a self-catalyzed process. Contrast has been set in order to distinguish the two WZ and ZB segments. Figure b) is reprinted from [41] (Sample 76715).

As presented in Figure 4.5, GaAs NWs have been grown with two distinct segments of WZ and ZB. Segments have been measured to be at least 350 nm (WZ) and 800 nm (ZB) with no stacking faults. So, direct comparison of the ZB and the WZ optical properties and transport properties is possible.

Band structure investigation of WZ and ZB

Many publications agree on the fact that polytypism in GaAs leads to a type II interface [START_REF] Spirkoska | Structural and optical properties of high quality zinc-blende/wurtzite GaAs nanowire heterostructures[END_REF][START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF]. Apart from the fact that type II interface could lead to trap carriers and give rise to indirect recombination, it highlights differences in the band structures from the two phases. ZB band structure has been well studied for a long time. But for the WZ phase this task has proved to be complex and works published have shown contradictory results [91,[START_REF] Martelli | Photoluminescence of GaAs nanowires at an energy larger than the zincblende band-gap: dependence on growth parameters[END_REF][START_REF] Spirkoska | Structural and optical properties of high quality zinc-blende/wurtzite GaAs nanowire heterostructures[END_REF][START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF]. The bandgap of WZ GaAs (or low temperature exciton transition) was reported to be either similar (within 5 meV difference) [91,[START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF] or 20-40 meV higher [START_REF] Ba | Resonant Excitation and Imaging of Nonequilibrium Exciton Spins in Single Core-Shell GaAs-AlGaAs Nanowires[END_REF][START_REF] Martelli | Photoluminescence of GaAs nanowires at an energy larger than the zincblende band-gap: dependence on growth parameters[END_REF] than that of ZB GaAs. In this work, a direct comparison of the optoelectronic properties between WZ and ZB can be realized within single NWs containing extended regions of the two phases. The top and the bottom part of the NW can be distinguished as the top has a tapered shape. On the integrated intensity map (Figure 4.6 b)), the top of the NW appears to be brighter, and the intensity rapidly decreases after hundreds of nanometers. The size of the bright part of the NW appears to be about 350 nm. This length is very similar to the WZ length measured on STEM images (Figure 4.5). The bright area can be attributed to WZ phase from the correspondence between STEM and panchromatic CL map. Spectra determined by averaging 3×3 pixels areas at different locations of the NW are plotted in Figure 4.6 c). We compared the spectra from the WZ, the ZB and the intermediary phase. The peak energy has been assessed for each spectrum using a 1-Gaussian function. A similar peak position is observed between the ZB and the WZ (1.513 eV). The intermediary phase spectrum (green) presents an intermediary peak position between WZ and ZB. A low energy peak at 1.469 eV, related to defects, is observed for all the spectra. A supplementary peak is observed on the WZ and intermediary segments spectra at 1.495eV attributed to to recombination involving a free electron and C acceptor [32]. Besides, we have seen that despite possible carrier diffusion, WZ and ZB have very distinct spectra. Average spectra of both WZ and ZB areas under different injection levels for one NW (different than the one presented in Figure 4.6) are plotted in Figure 4.7. The pulsed injection corresponds to an excitation with pulses of electrons (see Subsection 2.2), which is expected to be the lowest average injection level. Apart from the variation in the CL intensity no clear change of the shape of the CL spectra of the ZB phase is observed. A slight blue-shift from 1.513 eV to 1.515 eV and a broadening from 19 meV to 25 meV of the ZB main peak are observed from low to high injection levels. A similar trend is observed for the WZ spectra with a shift from 1.513 eV to 1.515 eV for the energy peak and a broadening from 24meV to 38meV for the FWHM. Besides, we can notice a clear evolution of the ratio between the main peak and the low energy defect peaks as we change the injection level. Under high injection, the broadening of the main peak can be attributed to hot carriers. When the incident current is decreased two other low energy peaks emerge, one at ≈1.465 eV and an other one at 1.495eV for the WZ spectra. From the evolution in the peak ratio, both the shoulder at 1.465 eV and at 1.495 eV are attributed to defects transitions. Under pulsed injection the WZ spectrum is governed by defects related peaks.

We showed that the peak observed at 1.513 eV for WZ, in Figure 4.6, is expected to be an excitonic transition. The energy peak position under low injection levels has been selected be-cause the blue-shift observed under higher injection levels is believed to caused by a band filling effect. No artefacts due to quantum confinement effects [START_REF] Hiruma | Growth and optical properties of nanometer-scale GaAs and InAs whiskers[END_REF][START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF] or strain [START_REF] Signorello | Tuning the Light Emission from GaAs Nanowires over 290 meV with Uniaxial Strain[END_REF] are expected due to sufficiently large NWs.

Vainorius et al. [START_REF] Vainorius | Observation of type-II recombination in single wurtzite/zincblende GaAs heterojunction nanowires[END_REF][START_REF] Vainorius | Confinement in Thickness-Controlled GaAs Polytype Nanodots[END_REF] have evidenced a type II interface due to a band offset of about 100 meV between the ZB and the WZ. A filling of the state at the interfaces with the injection level has been shown to cause a blue-shift of the transitions across the type II heterojunction. But, we did not observe such an effect during our experiments (see Figure 4.7). Besides, no clear WZ/ZB interface is expected as the intermediary segment is believed to be composed of a random alternation of WZ and ZB.

LT (10K) CL maps have been performed on a GaAs NW which presents distinct segments of pure WZ and ZB phases. Through CL analysis, we have seen that we were able to distinguish luminescence from the two different segments. The WZ phase has been observed to have a similar bandgap to the ZB at 10K (1.513 eV). This result is in agreement with results from literature [91]. Injection resolved CL has also been performed in order to highlight low energy defect transitions at 1.46-1.47 eV and at 1.495 eV in the GaAs WZ phase.

CL polarimetry and TRCL

GaAs NWs with distinct WZ and ZB segments have been prepared (Section 4.2.1), and studied through LT (10K) CL maps 4.2.2. Distinct regions have been observed in the hyperspectral CL maps consistently with the TEM images. In this Subsection we will first demonstrate that it is possible to distinguish unambiguously the luminescence emitted by the WZ from the one emitted by the ZB using CL polarimetry. Then time-resolved CL will be performed at two different locations of a single NW for direct comparison of carriers dynamics. We already discussed that GaAs nanowires can be grown with segments of the WZ and the ZB crystal phase. ZB belongs to the cubic hexagonal crystal family whereas WZ possesses a hexagonal symmetry (see Figure 4.8). The only difference between ZB and WZ structures is the placement of the third nearest neighbors. This difference leads to, for the WZ structure, a doubled unit cell along the c-axis (0001) compared to the one of ZB along the equivalent [111] direction. This break in the symmetry along the NW axis (c-axis) of the WZ semiconductor gives rise to remarkable differences in the band structures with the ZB counterparts. WZ and ZB band structures at the Γ point have been predicted and are depicted in Figure 4.9 a). In Figure 4.9, a schematic of the WZ band structure together with the dipole-allowed optical transitions are shown. We can notice that the dipole transition between Γ 9V -Γ 7C and Γ 9V -Γ 8C are forbidden in the direction perpendicular to the NW axis. This selection rule predicts the absorption, and so the emission, of photons is polarized perpendicular to the NW c-axis in the WZ phase [32,[START_REF] Ketterer | Raman Spectroscopy of GaAs Nanowires:Doping Mechanisms and Fundamental Properties[END_REF]. In Subsection 4.2.2 we showed that there is no intermediary between luminescence of the WZ and the ZB phases at LT. This is why we expect a strong linear polarization of the light emitted by the top of the NW (WZ) in the direction perpendicular to the NW. We characterized the polarization of light emitted by the two segments of the same NW using CL polarimetry. As described in Figure 4.2, a linear polarizer has been added to the CL setup in order to perform CL polarimetry. The 0°orientation was defined parallel to the NW. The orientation of the polarizer has been changed manually for each map. Prior to these experiments LT CL maps of a GaAs thin film, whose emission is supposed to be non polarized, have been measured with different polarizer orientations to measure the system response. Then average CL intensities have been determined and normalized with the average CL intensity measured for a LT CL map without a polarizer. In the following the CL panchromatic intensities have been corrected by the polarization of the system (Figure 4.11). Then LT CL maps with polarizer orientation ranging from 0°to 105°have been recorded. Panchromatic CL intensities have been determined for the different segments of the NW. In Figure 4.12, evolution of the CL intensities as a function of the linear polarizer orientation has been plotted. CL intensities from 120°to 180°have been extrapolated by symmetry, for instance:

CL polarimetry on GaAs nanowires

I CL (60°) = I CL (120°). I(θ) = I // × sin 2 (θ) + I ⊥ × cos 2 (θ) (4.1)
P = I // -I ⊥ I // + I ⊥ (4.2)
The CL intensity evolution with the polarizer orientation can be described by Equation 4.1, where I ⊥ and I // correspond to the component of the light polarized perpendicular and parallel to the NW orientation, respectively. In the case of a perfectly circular polarization we expect I ⊥ =I // , and for a perfectly linear polarization I ⊥ =1 and I // =0 or vice versa. Fits of Equation 4.1of the CL intensity evolution with the polarizer orientation are shown in Figure 4.12 c) (solid lines). From the fits I ⊥ and I // have been determined for the ZB and the WZ. The degree of polarization has then been calculated (Equation 4.2). No polarization was expected from the ZB segment but a slight 5% polarization degree is still observed after all. A strong polarization of -74 %is found for the WZ. The minus sign of the polarization degree (see definition Equation 4.2) is in agreement with the selection rules as we expected a linear polarization perpendicular to the NW axis. The evolution of the CL intensity with the polarizer orientation is plotted in Figure 4.13, it evidences a strong polarization of the light emitted by the WZ part of the NW. The polarization of the luminescence can be used as a tool in order to distinguish unambiguously the two segments of the NW.

The results regarding the linear polarization of light emitted by the WZ are close to results found by Ahtapodov et al. [4]. They found a degree of polarization of a WZ NW P=-71% using PL. Besides, the polar plot of the WZ and the ZB (Figure 4.13) is very similar to the one obtained by PL [4,[START_REF] Ba | Resonant Excitation and Imaging of Nonequilibrium Exciton Spins in Single Core-Shell GaAs-AlGaAs Nanowires[END_REF] and resonant Raman spectroscopy [91]. Still, a non-perfect linear polarization and a slight polarization of the light emitted by the ZB are obtained. Hoang et al. [START_REF] Ba | Resonant Excitation and Imaging of Nonequilibrium Exciton Spins in Single Core-Shell GaAs-AlGaAs Nanowires[END_REF] showed that polarization can also be induced by the difference in dielectric constant between the vacuum and the NW. Besides, due to finite size of the NW, an antenna effect [START_REF] Benjamin | Directional Emission from Leaky and Guided Modes in GaAs Nanowires Measured by Cathodoluminescence[END_REF][START_REF] Grzela | Nanowire Antenna Emission[END_REF][START_REF] Paniagua-Domínguez | Enhanced and directional emission of semiconductor nanowires tailored through leaky/guided modes[END_REF] can impact also its polarization.

CL polarimetry has been performed at LT on a NW in order to compare the polarization degree of both the WZ and the ZB segment. The strong perpendicular linear polarization (-74 %) of the WZ is in agreement with the predictions. The fact that the intensity has been shown to be polarized allows to attribute unambiguously this peak emission to the band to band transition of the WZ.

Lifetimes measurements in the WZ and ZB segments of a NW

Two CL maps at LT under pulsed excitation with two different polarizer orientations are shown in Figure 4.14. On these maps, we can observe that even under very low excitation power, the light emitted by the WZ is still linearly polarized. Indeed, in Figure 4.14 a) and b) we can see that the ratio between the WZ and the ZB intensities is strongly modified by the polarizer orientation. Despite the very low CL signal, as shown in Figure 4.14, it is still possible by acquiring two CL maps to distinguish unambiguously the WZ part from the ZB part. Using CL polarimetry we first identified the WZ and the ZB segment. Once they have been identified, the electron beam pulses are focused on one segment for streak image acquisition. Spectral integration have been made in the streak images (represented by the rectangles in Figure 4.15 a) and b)) in order to determine the decays which have been plotted in Figure 4.16 c). Lifetimes have been determined by a one-Exponential fits of the time-decays (exciton and defect) of the ZB (red) and the WZ (blue). A longer exciton lifetime is observed for the WZ τ X (W Z) = 0.63ns than ZB phase τ X (W Z) = 0.45 ns at LT. Besides, a longer lifetime has been determined for the defect transition in the WZ than in the ZB phase: τ def (W Z) = 1.95 ns and τ X (W Z) = 0.98ns.

We obtained GaAs NW lifetimes of ≈1 ns which is comparable to state of the art GaAs NW lifetimes at LT [4,[START_REF] Demichel | Impact of surfaces on the optical properties of GaAs nanowires[END_REF][START_REF] Perera | Nearly intrinsic exciton lifetimes in single twin-free GaAs????AlGaAs core-shell nanowire heterostructures[END_REF]. Besides, longer lifetimes are observed for the defect and the exciton transitions in the WZ. A longer defects lifetime for the WZ phase is consistent with a higher radiative efficiency observed on the pulsed LT CL panchromatic map (Figure 4

Conclusion of the Section

III-V compounds usually form a cubic structure in bulk or thin films but they may adopt a hexagonal structure when grown in the NW form. Because of the difficulty to control the phase during the NW growth, ZB/WZ polytypism coexists in single NWs. The switch in the phase may have a detrimental impact on the device performances. This is why, is this section we focused on the investigations of the WZ and the ZB phases optical and transport properties. To this end, undoped GaAs NWs have been prepared with two distinct segments of WZ and ZB pure phases where no stacking faults have been found. First, the band structure have been investigated thought LT CL maps. A very similar bandgap (E g = 1.513 eV ) between the two phases has been evidenced. Then, CL polarimetry has demonstrated a strong linear polarization of the WZ (P=-74 %), which is in agreement with symmetry considerations. Complementary TRCL experiments have been performed to determine lifetimes in the ZB and the WZ phase of a single NW. A longer exciton lifetime has been found for the WZ (0.63 ns) than for the ZB (0.45 ns). But a clear variation of the defects density between the NWs has been observed from the variation in the CL lineshape at low temperature. A statistical analysis of the fluctuations between the NWs of the defects density in the phases may be the next step to analyze the impact of the presence of polytypism in the NWs.

Te incorporation in GaAs nanowires for n-type doping

Doping management is a prerequisite for the realization of a high efficiency PV cells. Considering NWs, VLS-mediated growth results in doping incorporation are quite different from the thin film case [START_REF] Lapierre | III-V nanowire photovoltaics: Review of design for high efficiency: III-V nanowire photovoltaics: Review of design for high efficiency[END_REF]. Various challenges have been identified such as the amphoteric behavior of the Si dopant which has been shown [START_REF] Ketterer | Compensation mechanism in silicon-doped gallium arsenide nanowires[END_REF] to be a real problem in the case of NW. Besides, doping assessment in single NWs is very challenging. Indeed main characterization techniques (Hall effect, C-V...) are based on electrical contacts and highly challenging lithography are needed for their application on NWs [167]. This is why contactless characterization techniques are desirable.

Optical techniques such as resonant Raman spectroscopy [START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF]73,[START_REF] Hilse | Incorporation of the dopants Si and Be into GaAs nanowires[END_REF][START_REF] Ketterer | Compensation mechanism in silicon-doped gallium arsenide nanowires[END_REF] and PL [START_REF] Arab | Doping concentration dependence of the photoluminescence spectra of n-type GaAs nanowires[END_REF][START_REF] Borghs | Band-gap narrowing in highly doped n -and p -type GaAs studied by photoluminescence spectroscopy[END_REF][START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF][START_REF] Orrù | A Roadmap for Controlled and Efficient n-Type Doping of Self-Assisted GaAs Nanowires Grown by Molecular Beam Epitaxy[END_REF] have been used for the doping assessment in NW arrays or single NWs. Hakkarainen et al. [73] studied the incorporation of Te, as a type n dopant, within the GaAs NW using a set of characterization techniques on several NWs: Atom probe tomography, Raman spectroscopy and electron holography. A good incorporation of Te impurities has been found (4×10 18 cm -3 ) using APT. Homogeneity along the NW axis has been investigated with relative evolution of the raman spectra. A shift in the coupled plasmon phonon peak has been attributed to an increase of the doping level along the NW axis. In our team at C2N, Chen (a previous PhD student) et al. [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF][34][START_REF] Chen | Quantitative Assessment of Carrier Density by Cathodoluminescence (2): GaAs nanowires[END_REF] have implemented a method to assess the doping level (p or n type) in single NWs based on CL. This technique relies on the evolution of the lineshape of the luminescence spectrum with the doping level and type, probed at the nanoscale. This method has been shown to be able to assess quantitatively both p-type (Be) and n-type (Si) doping in NWs, using the fit of the complete RT CL spectra with the generalized Planck law and determination of the FWHM in LT CL spectra [START_REF] Chen | Quantitative Assessment of Carrier Density by Cathodoluminescence (2): GaAs nanowires[END_REF].

In this section we will first explain the method developed at the C2N by chen et al. [34] for the quantitative assessment of the carrier density. As a result, n-type doping of 6 × 10 17 cm -3 to 10 18 cm -3 in n-GaAs:Si NWs have been inferred in [34]. In order to increase the electron concentration the n-type doping using Te has been explored in this thesis. For the purpose of extending the doping assessment method to the Te doping in GaAs, RT and LT CL maps were measured on Te doped thin-films and will be presented in the following Section. They have been compared to Hall-effect measurements in order to serve as a reference. Then LT and RT CL maps have been performed on Te doped GaAs NWs for doping assessment and investigation of its homogeneity.

Quantitative n-type doping assessment method based on CL

Owing to the nanoscale of the NWs, doping assessment through electric measurements is a complex task. Optical contactless characterization techniques such as Raman spectroscopy or PL [START_REF] Borghs | Band-gap narrowing in highly doped n -and p -type GaAs studied by photoluminescence spectroscopy[END_REF][8] [START_REF] Haufe | The luminescence lineshape of highly doped direct-gap III-V compounds[END_REF] have been explored in order to assess the doping level III-V materials. For instance, De sheng [START_REF] De-Sheng | Electrical properties and photoluminescence of Te-doped GaAs grown by molecular beam epitaxy[END_REF] has analyzed the lineshape evolution with the doping level (Te dopant) of luminescence spectra at LT based on the Burnstein Moss effect. However, due to their low spatial resolution these techniques are not suitable to characterize doping gradient at the nanoscale. Chen et al. [33][34][35] presented in 2017 an effective method of determining the doping level in n-type GaAs at the nanoscale using CL measurements and performed on single Si-doped GaAS NWs. In this Subsection details of this method, developed at the C2N, will be given. Then I will present and analyze CL epxeriments at RT and LT performed on Te doped GaAs thin films in order to extend the results obtained for the Si doping to the case of a Te dopant.

Modeling of CL spectra for quantitative doping assessment in thin-films

Si doped GaAs thin-films preparation GaAs thin films were grown by MBE on semiinsulating GaAs (001) substrates. The thicknesses of the films have been measured from crosssection SEM images of cleaved samples and estimated to range from 320 nm to 850 nm. Si was used for n-type doping of GaAs. No additional surface passivation layer was grown on the top of the doped GaAs layers. Carrier mobility and concentration were measured afterwards by Hall effect measurements. The effective doping range of the Si doped GaAs thin films which have been prepared and presented in [34] varies from 2.2 × 10 17 cm -3 to 7.6 × 10 18 cm -3 .

Room temperature measurements RT CL spectra recorded on Si doped GaAs thin films are plotted in Figure 4.17 a). A blue-shift and a broadening of the CL peak when increasing the carrier concentration is observed. The peak energy shifts from 1.42eV for low electron concentration up to 1.52eV for high concentrations. This behavior is attributed to a band filling (Burnstein Moss effect). As described in Figure 4.17 c), band filling is characterized by a Fermi level of electrons within the conduction band. In degenerate GaAs, E fc is expected to be above the CB minimum (Figure 4.17), and free electrons fill the conduction band. N-type GaAs is degenerate for doping concentration above 5×10 17 cm -3 , hence the absorption edge and luminescence peak shift to higher photon energy (blueshift) due to filling of the conduction band [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF]. [34].

A fit of the luminescence spectrum has been made in order to assess for carrier concentration in the thin films. The generalized Planck law has already been shown to describe the luminescence spectrum Φ SC ( ω) as a function of the absorptivity A( ω), and quasi-Fermi levels for electrons (E fc ) and holes (E fv ), at a given photon energy ω. We reminded below a simplified version of the generalized Planck law (Equation 4.3):

Φ SC ( ω) = A( ω) × Φ bb ( ω) × exp( E fc -E fv kT ) (4.3)
Doping impacts significantly the absorption of the SC near the bandgap. The parabolic ideal absorption model (Equation 2.14) for GaAs has been estimated using α 0 (1.6eV ) = 14800cm -1 . Then the convolution of the ideal absorption α ideal with an Urbach's tail sub-bandgap absorption model has been used (Equation 2.13). Correction of the absorption has also been made with the occupation probability (see Subsubsection 2.4.1.2). Fits of RT CL spectra have been performed and plotted in 4.17 a). The fitting parameters which have been used were E g , the Urbach's tail parameter γ and E fc and the E fv has been set to be above the valence band.

The evolution of the Fermi level as a function of the doping level has been plotted in Figure 4.18 (blue squares), using prior doping assessment with Hall effect measurements. Theoretical evolution (Equation 4.4) has also been plotted for two different absorption models (blue solid and green dashed lines). The ideal absorption can be corrected from its non-parabolicity using Equation 4.4, where F j is the Fermi integral and N c the conduction band density of state. The non-parabolicity factor which has been used in [34] is β=-0.83. As we can observe a slightly lower E f c is obtained. [34].

n = N c F 1/2 ( E fc -E fv kT ) - 15βkT 4E g F 3/2 ( E fc -E fv kT ) (4.4)
Comparison between experimental and theoretical data has been plotted in Figure 4.18. We can observe a relatively good agreement between the theoretical and the experimental evolution of the Fermi level. It is especially the case for the high electron concentrations (n>10 18 cm -3 ). In the following, electrons concentration at RT will be estimated using Equation 4.4 based on the Fermi level estimation with the generalized Planck law fit.

Low temperature measurements LT CL spectra have also been measured for n-doped GaAs thin films and plotted in Figure 4.19. As for RT measurements, a blue-shift and a broadening of the CL peak is observed when increasing the n doping level. A shoulder near 1.48 eV, observed for the highest n-doped sample, may be attributed to the recombination through a Si acceptor (Si As ). Besides, a broad peak at around 1.29 eV can be noticed and attributed to recombination involving the complex formed by a Si donor (Si Ga ) and a Ga vacancy [34].

De Sheng et al. [START_REF] De-Sheng | Electrical properties and photoluminescence of Te-doped GaAs grown by molecular beam epitaxy[END_REF] have developed a semi-empirical model based on the evolution FWHM with the electron concentration of PL spectra for Te doped thin films. In a first order approximation the FWHM is expected to be proportional to the electron Fermi level. Then a 3D electron gas is considered to occupy a Fermi sphere of radius k F = (3πn) 1/3 . E fc is proportional to k F 2 for parabolic conduction band and is thus proportional to n 2/3 .

Based on this reasoning, the FWHM as a function of the electron concentration (Figure 4.20) has been fitted with a 2/3 power law, leading to Equation 4.5. We can notice a good agreement between the 2/3 power law and the experimental data plotted in Figure 4.20. Besides, this law seems to be even relevant for electron concentration as low as 5 × 10 17 cm -3 . F W HM (eV ) = 3.75 × 10 -14 × n 2/3 (4.5)

Shifts and broadening in the luminescence spectra were observed and modeled by Chen et al. (at the C2N) in order to measure quantitatively the electron concentration in n-type GaAs:Si thin films. Shifts at RT are explained by the Burnstein-Moss effect. RT CL spectra have been fully fitted using Equation 4.3. The determination of the electron concentration from the fitting parameters has been shown to be consistent with Hall effect measurements for n > 10 18 cm -3 . At lower doping levels, the FWHM of LT CL spectra can be used to determine accurately the doping level down to 5 × 10 17 cm -3 , using an empirical law (Equation 4.5). By measuring the electron concentration in Si doped GaAs NWs using this method, the difficulty to increase the carrier concentration with the Si doping has been highlighted.

Te doped GaAs thin-films quantitative doping assessment

Sample preparation Te-doped GaAs thin-films were grown on semi-insulating GaAs(100) substrates, at the University of Tampere (Finland). Te was provided from a GaTe cell with a nominal dopant concentration of 2.0 × 10 Room temperature measurements RT hyperspectral CL maps (64×64 pixels) were measured for all TF, identified in Table 4.1, over a 6×6 µm 2 area. CL spectra were determined by averaging the 64×64 spectra in order to improve the signal/noise ratio. Standard injection conditions were set (E = 6keV, I = 2nA). Spectra were normalized and shifted vertically with increasing electron concentration values and plotted in Figure 4.21.

We observe a blue-shift and a broadening of the CL peak from 1.428 eV to 1.55 eV with increasing n-doping levels. Besides, for high doping (D1) an other peak at 1.28 eV is observed. The peak is attributed to deep level states [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF]. A very good agreement with fits (red lines in Figure 4.21) with the generalized Planck law (Equation 4.3) is found, except regarding the appearance of a low energy defects peak.

Fermi levels and the other fitting parameters are given in Table 4.2. E fc is the electron Fermi level taking the minimum of the conduction band as the reference. No band filling 4.2) together with the electron concentration determined with the hall effect (summarized in Table 4. 1).

(E fc -E CB = 0eV ) is observed for A1 and B1 samples. A non-zero E fc has been obtained for C1, C3, D1 and A3. We observed an increase of E fc consistent with the increase of the electron concentration measured with Hall effect. We can notice a red-shift of the Bandgap (E g ) with the doping level correlated with the overlapping of the donor levels with the conduction band. But because of the low energy peaks, inconsistent value of the bandgap is fitted for the highest doping level. Besides, as we increase the dopant concentration, a greater sub-bandgap defects density is anticipated and correlated with the Urbach's tail parameter (γ) value. Very low E fc have been obtained for A3 and C1 so a low reliability is given to the electron concentration determined from these values. Indeed, below a certain value the E fc no longer impacts the shape of the CL spectrum. Electron Fermi level obtained from the fit of CL spectra and theoretical curves calculated Low temperature measurements LT (10K) CL measurements have also been performed on the Te doped GaAs thin films (TFs), over a 6×6 µm 2 area. Average LT CL spectra are plotted in Figure 4.22 a). Then they have been normalized and shifted vertically. A clear broadening and a blue-shift of the CL peaks is observed. Peak position and FWHM for each peak have been determined using a one-Gaussian function and are reported in Table 4.3.

Sample name

Experimental (black dots) and empirical relations have been compared in Figure 4.22 b). Electron concentrations determined with the FWHM of the CL peak at LT are compared with ones measured with the Hall effect measurements in Table 4.3. A very good agreement between the two methods is observed. We can notice that the sample for which we have the highest difference between electron concentrations determined by CL and Hall effect measurements is C1. Such a discrepancy has also been observed with the complete fit of the RT CL spectrum. By excluding this value, a mean deviation of ≈ 9% is found. This result allows us to put a high level of reliability in our measurements and to suggest redoing the Hall effect measurements for this sample. Besides, the evolution of the mobility determined with the Hall effect measurements (Table 4.1) is not consistent with the carrier concentration. Deviation from the empirical models may originate from the FWHM estimation. In the C3 and D1 spectra, a shoulder at about 1.48 which corresponds to donor related peaks [START_REF] Arab | Doping concentration dependence of the photoluminescence spectra of n-type GaAs nanowires[END_REF]34,[START_REF] Isik Goktas | Doping assessment in GaAs nanowires[END_REF][START_REF] De-Sheng | Electrical properties and photoluminescence of Te-doped GaAs grown by molecular beam epitaxy[END_REF]. Appearance of defect related shoulders on the CL spectra may lead to an error in the estimation of the FWHM even if this effect is smoothed by the estimation of the FWHM with a Gaussian function. In this Subsection we first described the method developed by Chen et al. [START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF]34] (at C2N) in order to determine the electron concentration in Si doped GaAs NWs. Based on this method, the difficulty of reaching carrier concentration higher than 1 × 10 18 cm -3 has been evidenced. This is why as a second step we tried to extend this method to the case of the Te doping. To do this end, we measured at low-and room-temperature CL spectra for Te doped GaAs thin-films grown at the Tampere University. A very good agreement between the Hall effect measurements of the electron concentration and the estimation from CL spectra has been found at both RT and LT. These results attest that the quantitative determination of the carrier concentration implemented for the n-type Si doping can be applied in the case of the Te doping.

Sample name

Quantitative Te doping assessment in GaAs nanowires

In the previous Subsection we attested that the method for quantitative determination of the electron concentration in Si doped GaAs thin-films can be transposed to the case of the Te doping. In this Subsection we will apply this method to determine quantitatively the carrier concentration and its homogeneity within a Te doped GaAs NW. Self-catalyzed GaAs NWs were grown on p-doped Si (111) wafers by MBE using a selfassembled oxide patterning method. After a predeposition of Ga, the growth was started by providing both Ga and As fluxes. As for the thin-films, Te was provided from a GaTe cell with a nominal dopant concentration of 2.0 × 10 19 cm -3 . No passivation shell has been grown after the growth of the GaAs:Te core. Geometries and relevant growth parameters have been resumed in Table 4. 4. Other details on the growth process are given in [START_REF] T V Hakkarainen | Lithography-free oxide patterns as templates for self-catalyzed growth of highly uniform GaAs nanowires on Si(111)[END_REF]73]. Subsequently, the NWs have been dispersed on a Si substrate for CL analysis. This GaAs:Te NW sample is provided by Tampere University who already published in [73] an investigation of the electron concentration and its homogeneity along the NW. It has been estimated using atom probe tomography, assuming a 100% activation of the Te dopant, that the top of the NW presents a electron concentration of n ≈ 4 × 10 18 cm -3 . Besides, a gradient of the free-carrier concentration which increases in the axial direction from the bottom to the tip of the NW has been evidenced with off-axis electron holography and Raman results.

High resolution hyperspectral RT (Figure 4.24) and LT (Figure 4.23) CL maps have been measured with similar injection conditions than the one used during the acquisition of the CL maps on thin-films (E=6 keV, I=2 nA). This has been made in order to avoid band filling artifacts due to high injection levels. During CL experiments, charge effects have been observed which may be induced by surface states. We can observe that the live-SEM image is drifted due to beam deflection in Figure 4.23 a). For some maps (not shown) we were not able to measure the whole NW as it "moved out" of the scanning area. This drift has been observed to be very dependent of the exposure time, the pixel size of the maps and even of the sample. We optimized the spatial resolution by increasing the number of pixel but still being able to acquire the whole NW. In order to correct this effect, an image processing has been performed and presented as "Corrected-SEM" image.

LT hyperspectral CL map measurements of Te NW1 are shown in Figure 4.23. Average CL spectra have been determined by averaging the CL spectra over the areas described by the white squares along the NW axis. Then spectra have been normalized and plotted in Figure 4.23 d). At the bottom of the NW (spectrum (8)), luminescence is mainly due to recombination through the donor related peak at E= 1.335 eV. Then, when moving towards the top of the NW, peaks at E= 1.575 eV increase in intensity, attributed to band to band. The band to band transition intensity increases and even slightly blue-shifts from spectrum (1) to [START_REF] Arab | Doping concentration dependence of the photoluminescence spectra of n-type GaAs nanowires[END_REF]. Still, LT maps can not be used for the carrier concentration determination due to intense low-energy emissions (defects) screening the band to band emission peak. A deconvolution of the luminescence spectra is needed in order to determine the FWHM of the band to band transition CL peak.

A RT hyperspectral CL map on Te NW1 is shown in Figure 4.24. Nine CL spectra have been determined by averaging over 4 × 4 pixels areas along the NW axis. On these spectra, CL peaks located around 1.42 eV corresponding to the bandgap of GaAs can be observed. Besides, a broadening and a blue-shift of the CL peak can be noticed from the bottom (spectrum (9), Figure 4.24) to the tip (spectrum (1), Figure 4.24) of the NW. This result suggests an increase of the electron concentration along the NW axis from the bottom to the top. This evolution is consistent with the observed increase of the CL intensity along the NW axis which may be explained by an increasing of the doping level inducing a reduction of the depletion layer [START_REF] Chen | Quantitative Assessment of Carrier Density by Cathodoluminescence (2): GaAs nanowires[END_REF][START_REF] Andrew | Analytical model of surface depletion in GaAs nanowires[END_REF].

We have selected small rectangles along the nanowire axis with a length corresponding to 300-350 nm, and we have integrated the CL intensity over these areas. Then, we have fitted the resulting CL spectra with the generalized Planck law. The experimental CL spectra (black dots) and the fitted curves (red solid lines) are plotted for four rectangles in Figure 4.26. The fit was applied only on the main CL peak emission. We can observe a very good agreement between the fit and the RT CL spectra.

Electron quasi-Fermi (E f c ) levels have been determined from nine RT CL spectra along the NW axis as presented in Figure 4.26. From the E f c values, an electron concentration has been estimated. We found a high doping gradient along the top ≈ 3µm part of the NW, with electron concentration ranging from 3.3 × 10 18 cm -3 down to 1.6 × 10 18 cm -3 . In the bottom part, we see a continuous shrinking of the luminescence peak width, but we are not able to fit the spectra and to provide a reliable assessment of the carrier concentration. It is interesting to note that the low-energy (sub-bandgap) luminescence seems to follow the inverse tendency of the estimated Te doping: it is increasing towards the bottom of the NW. High electron concentration (> 10 18 cm -3 ) have been assessed in the top part of the NW using Te as a dopant, which are close to the dopant incorporation. It is particularly promising because a nearly 100% activation ratio of Te dopants is expected leading to a low defects incorporation, due to the dopant, in the NW. Still, a variation by more than a factor two of the doping has been measured in the NW. As a consequence, a careful control of the growth conditions is needed to obtain a uniform doping along the NW.

Conclusion of the section

In this section we discussed of the quantitative doping assessment in n-type GaAs NWs. Because of their nanoscale, carrier concentration estimation in a NW is a complex task. However, the control of the doping in a structure is a prerequisite for PV applications. Chen et al. [34] developed at C2N a method in order to determine the electron concentration in Si doped GaAs NWs. We first explained this method based on the evolution of the CL spectra with the carrier concentration. Then we demonstrated that this method can be transferred to the estimation of the doping in case of Te as a dopant. To this end, we measured RT and LT CL spectra on Te-doped GaAs thin-films. A very good agreement between the electron concentration obtained with CL and the ones from Hall effect measurements has been found. Based this result, we measured the doping level in a n-type GaAs:Te NW. A doping of n = 3.3 × 10 18 cm -3 has been determined at the top of the NW. This result shows that high electron concentration can be achieved using Te dopants, which is extremely promising for device application. We have also evidenced a doping gradient with a spatial resolution as low as about 300 nm. This issue still needs to be addressed in order to use Te as a dopant in NW based solar cells.

Impact of passivation shells on GaAs nanowires

Due to a large surface/volume ratio, NWs lifetimes are highly dependent on surface recombination. Chia et al. [START_REF] Andrew | Analytical model of surface depletion in GaAs nanowires[END_REF] has determined analytically that a lightly doped GaAs NW, with a radius of 100 nm, would be fully depleted because of a high surface recombination velocity (> 10 5 cm.s -1 ). Surface passivation can be achieved by growing a thin layer of a high bandgap materials (such as InGaP of AlGaAs) acting as a barrier for carriers.

Demichel et al. [47] have investigated surface passivation through the comparison of AlGaAs capped and uncapped samples using low temperature µ-PL. Surface passivation has been shown to be able to increase minority carrier lifetimes by two orders of magnitude (for ≈100nm NWs diameter). Skold et al. [START_REF] Sköld | Growth and Optical Properties of Strained GaAs-Ga x In 1 -x P Core-Shell Nanowires[END_REF] have investigated surface passivation using an InGaP passivation shell. It has been demonstrated that growing a InGaP shell allows to improve the luminescence intensity by a 3 to 4 orders of magnitude compared to unpassivated GaAs NWs. Besides strain effects, imposed by a possible lattice mismtach between the InGaP shell and the GaAs core, on the band gap of the core were measured for different shell compositions using photoluminescence.

In this Section we investigate the impact of a high bandgap passivation shell on the radiative efficiency on GaAs NWs, lifetimes or bandgap using room-and low-temperature CL and TRCL measurements. First, in Subsection 4.4.1 a simple model (Equation 4.6) is applied to assess the impact of the surface passivation on the NWs lifetimes and the cells parameters. Then, in Subsection 4.4.2 the GaAs-AlGaAs and the GaAs-InGaP core-shell NWs structures will be compared. Low-temperature CL and TRCL measurements have been performed to analyze their defect densities, lifetimes and bandgaps, and then their room-temperature radiative will be compared. Results on the lifetimes measurements in single InGaP-GaAs NWs are presented in Subsection 4.4.3. These results, coupled with the model explained in Subsection 4.4.1 are used to give an upper limit to the surface recombination velocity in InGaP-GaAs NWs. The strain induced by the lattice mismatch between the GaAs core and InGaP the shell has been estimated quantitatively from the shift of the luminescence spectra. I exploit the strain, localized at the top of the NW, to estimate the diffusion coefficient in an InGaP-GaAs NW.

Impact of the surface passivation

For homoepitaxial GaAs thin films, minority carrier lifetimes of about 20 ns are obtained [START_REF] Sieg | High minority-carrier lifetimes in GaAs grown on low-defect-density Ge/GeSi/Si substrates[END_REF]. Despite relatively high dislocation density, Sieg et al. [START_REF] Sieg | High minority-carrier lifetimes in GaAs grown on low-defect-density Ge/GeSi/Si substrates[END_REF] have obtained bulk lifetimes of 7.7 ns for GaAs epitaxial growth on Si. In the case of NWs, lifetimes are still limited around 1.2 ns at LT [4,30,40,[START_REF] Demichel | Impact of surfaces on the optical properties of GaAs nanowires[END_REF][START_REF] Jiang | Enhanced Minority Carrier Lifetimes in GaAs/AlGaAs Core-Shell Nanowires through Shell Growth Optimization[END_REF]. These low lifetimes can be explained by the role played by surfaces.

The effective lifetime can be decomposed between a bulk (τ bulk ) and a surface component. In the case of a NW, Dan et al. [40] have shown that the effective lifetime can be written as in Equation 4.6. This Equation details the influence of the surface recombination velocity (S) and the NW diameter on the effective lifetime (τ ef f ). Indeed, from Equation 4.6, if we increase S we decrease τ ef f for a given diameter. Besides, by decreasing the diameter (d) of the NW, the surface component increases and so the effective lifetimes decrease.

1 τ ef f = 1 τ bulk + 4 × S d (4.6)
In order to emphasize the impact of S, we plotted in Figure 4.27 τ ef f as a function of the NW diameter for two values of S. In order to estimate the evolution of the effective lifetimes we assumed a bulk lifetime τ bulk = 20 ns [START_REF] Sieg | High minority-carrier lifetimes in GaAs grown on low-defect-density Ge/GeSi/Si substrates[END_REF]. As we can observe in Figure 4.27, reducing S from S = 1 × 10 5 cm.s -1 to S = 2.5 × 10 3 cm.s -1 leads to an increase of τ ef f by two orders of magnitude. Such a surface passivation has already been realized using an AlGaAs passivation shell [30,[START_REF] Demichel | Impact of surfaces on the optical properties of GaAs nanowires[END_REF]. Considering a NW diameter of d = 200 nm, τ ef f has been increased from 5 ps to 1.8 ns by reducing S of about two orders of magnitude. 4.6. We considered a τ bulk =20 ns.

Using an InGaP passivation layer, Marianni et al. [111] demonstrated a good surface passivation by increasing the PCE from 1.02 % to 6.63 % of a GaAs nanopillar-array solar cell. Indeed, by decreasing non-radiative recombinations, the carrier concentration is increased and so the voltage within the cell. On the basis of a simple model, we can estimate the gain in V oc made by a surface passivation. The link between the carrier density and the voltage within the cell is given by Equation 4.7:

n × p = n i 2 × exp( qV kT ) (4.7)
The excess minority carriers generated through a light excitation is then given by Equation 4.8, where G is the generation rate:

∆n = G × τ ef f (4.8)
Assuming a low injection regime and a residual p doping (N a ) Equation 4.7 can be rewritten as Equation 4.9. The gain in voltage (∆V ) induced by an increased lifetime of minority carriers can then be calculated with Equation 4.10, only depending on the carrier lifetimes τ 1 and τ 2 . In this model, the voltage gain is expected to evolve as a logarithm of the ratio between τ 1 and τ 2 .

∆n × N a = n 2 i × exp( qV kT ) (4.9) ∆V = V 1 -V 2 = kT q log( τ 1 τ 2 ) (4.10)
In the case of a reduction of surface recombination velocity of about two orders of magnitude, we showed that the effective lifetime can be improved from 5 ps up to 1.8 ns (example above). Thus the gain in voltage (∆V ) can be estimated to be of about 150 meV with Equation 4.10.

Using a simple model we estimated the impact of the surface passivation on effective lifetime in NWs. A surface passivation based on the growth of an AlGaAs shell has been shown to be able to reduce the SRV down to about S= 2.5×10 3 cm.s -1 [START_REF] Demichel | Impact of surfaces on the optical properties of GaAs nanowires[END_REF] [30]. Based on this result, we gave an example of the increase of the effective lifetimes achieved by surface passivation for GaAs NW with a diameter d = 200 nm using Equation 4.6. Then the gain in voltage (∆V ) induced by an increased lifetime of minority carriers was estimated to be of about 150 meV. This estimation provides a motivation in order to realize a surface passivation of GaAs NW for PV applications.

Comparison between AlGaAs and InGaP passivation shells

Efficient surface passivation has already been shown using an AlGaAs passivation shell. But InGaP has also been shown to be a promising material in order to form a passivation shell [START_REF] Pearton | Comparison of surface recombination velocities in InGaP and AlGaAs mesa diodes[END_REF]. In this Subsection we present low-and room-temperature measurements performed on GaAs-InGaP and GaAs-AlGaAs core-shell NWs to compare the two passivation shells.

Low-temperature comparison

Two types of NWs structures, an undoped GaAs-InGaP NW (76851) with core diameter of 200nm and an InGaP shell of 35nm and an undoped GaAs-AlGaAs NW (76715) with a core diameter of about 120nm and an AlGaAs shell of about 35nm, and a thin GaAs shell was added to avoid oxidation of AlGaAs. Low-temperature CL maps have been recorded on these two structures. Both CL maps are measured in the very same excitation and detection conditions and can be compared quantitatively. In Figure 4.28 is presented CL intensity maps as well as SEM images. Besides, the energy peak position (Figure 4.28c) and h)) and the FWHM (Figure 4.28d) and i)) have been assessed at each point of the CL map for the two samples.

Homogeneous low temperature CL intensities can be noticed for the two samples throughout the NWs. Moreover, a four times higher CL intensity is observed for the InGaP shell structure as compared to the AlGaAs. In Figures 4. 28 c) and h) we can notice that the peak position in the NW is not perfectly homogeneous along the NW. For the GaAs-AlGaAs NW, this inhomogeneity in the peak position has already been discussed in Section 4.2. Besides, a clear difference between the two samples of the peak position can be remarked. At 10K, the GaAs bandgap is expected to be of 1.51eV and so the peak emission to be close to this value. It is the case for the GaAs-AlGaAs structure, but not for the GaAs-InGaP one. This shift can be explained by the strain induced by a lattice mismatched InGaP shell [START_REF] Signorello | Tuning the Light Emission from GaAs Nanowires over 290 meV with Uniaxial Strain[END_REF][START_REF] Sköld | Growth and Optical Properties of Strained GaAs-Ga x In 1 -x P Core-Shell Nanowires[END_REF]. This effect will be further analyzed in Subsection 4.4.5. In addition to a shift in the peak, a broadening of the CL peak can be observed, in Figures 4.28d images near the excitonic peak (between white lines in Figure 4.28a) and b)). Fits with 1exponential functions have been performed in order to assess for the lifetimes. Fitting has been made only on a portion of the time decay.

Lifetimes, estimated with the fits of the time decays, have been plotted in Figure 4.29. Direct correlation with the radiative efficiency of each sample is not obvious. Because of a higher CL intensity estimated on the InGaP sample, a longer lifetime was expected. Indeed a higher CL intensity goes with a higher radiative efficiency and so a lower rate of non-radiative recombinations η ∝ τ ef f /τ r . More insights are needed to fully understand this behavior at LT.

Room-temperature measurements: radiative efficiency comparison

Room-temperature CL maps have also been conducted on the same samples (but a different NW in the case of the GaAs-InGaP structure). In Figure 4.30 we compared room-temperature CL intensity maps for the two structures. We observed an almost two times higher CL intensity for the GaAs-InGaP structure than for the GaAs-AlGaAs structure. But, the CL intensity is less homogeneous in the case of the InGaP shell. A higher radiative efficiency at RT is therefore expected for the InGaP shell structure compared to the AlGaAs.

Lifetimes determination in GaAs/InGaP core-shell at room-temperature

In this subsection we present room-temperature measurements conducted on single GaAs-InGaP nanowires in order to assess for the effectiveness of the surface passivation. After the measurements presented in Figure 4.30, we switched to pulsed excitation conditions. First, spectra (under pulsed conditions) were recorded for three different NWs with the CCD camera. The CL spectra were recorded on the highest CL intensity area of the NW. NW1 has already been presented and compared to an AlGaAs passivated GaAs NW in the previous Subsubsection at both RT and LT. No measurements were conducted on the GaAs-AlGaAs because no CL signal has been obtained at RT under pulsed conditions.

A spreading of the peak position can be observed between the NWs. A red-shift of the peak seems to be followed with a decrease in the CL intensity, under pulsed conditions. Besides, two peaks seem to be distinguishable: A first one at λ 1 ≈ 870 nm corresponding to a classic GaAs luminescence peak, and one located at λ 2 ≈ 890 nm corresponding to a strained GaAs phase. Variation of a factor 2 between the highest CL intensity (NW1) and the lowest CL intensity (NW3) can also be noticed.

Streak images were recorded for the three NWs. Integration was performed for 20 min, until the TRCL signal was too weak. The disappearance of the CL signal can be attributed to two reasons. The first one is the drift of the stage, so that the beam is no longer on the NW. An other factor is the drift of the LASER on the tip which induces a weaker electron beam. After At least one order of magnitude for the signal/noise ratio was obtained for each time decay.

Lifetimes have been determined from the time decays with a 1-exponential function, from 190 ps (NW3) up to 400 ps (NW1). The trend in the lifetimes at RT between the NWs is consistent with the CL intensities measured under pulsed conditions and presented in Figure 4.31. Indeed, a factor two is observed between the CL intensities of NW1 and NW3. The same factor two is obtained for the NW1 and NW3 lifetimes.

The measured lifetimes can be compared to the bulk lifetime of defect-free GaAs grown on Si which are of about 5 ns [START_REF] Sieg | High minority-carrier lifetimes in GaAs grown on low-defect-density Ge/GeSi/Si substrates[END_REF]. Two orders of magnitude lower lifetimes have been obtained on our NWs. This result shows that the NW effective lifetimes can still be improved. Room-temperature carriers lifetimes have been measured using TRCL in single GaAs-InGaP core-shell NWs. Lifetimes ranging from 190 ps to 400 ps have been measured. To my knowledge we presented the first reported time decays measured on single III-V nanowires by TRCL at RT.

Upper bound for the surface recombination velocity

In the previous Subsection 4.4.3 we presented lifetimes measurements on single NWs at RT. In this Subsection we will discuss these results. By coupling these measurements with a simple model we can estimate the upper bound for the surface recombination velocity (S). regime where non-radiative recombinations at the NW surface dominate (high S values).

We plotted in

Lifetimes ranging from 0.19 ns to 0.4 ns have been measured previously and an average lifetime of 0.31 ns can be estimated. τ ef f = 0.31 ns corresponds to the light blue region/area in Figure 4.34. Asymptotic values for S = 2 × 10 4 cm.s -1 and τ bulk = 0.3 ns have been plotted as vertical and horizontal black dashed lines. These two asymptotes correspond to the two extreme behaviors when the τ ef f is governed by 1) bulk defects and 2) surface recombination. An upper bound, corresponding to the horizontal asymptote, can be given by S = 1.7 × 10 4 cm.s -1 , considering 1/τ bulk 4 × S/d. But complementary experiments need to be performed in order to conclude on whether the lifetimes are governed by defects within the core or by an inefficient surface passivation by the InGaP shell.

Demichel et al. [START_REF] Demichel | Impact of surfaces on the optical properties of GaAs nanowires[END_REF] has estimated the SRV of a AlGaAs/GaAs NWs to be reduced down to 2.5 × 10 3 cm.s -1 . In our case, even if we largely overestimate the S value, we still show that we succeeded in reducing by at least one order of magnitude the SRV by coupling TRCL measurements at RT and a simple model for τ ef f .

Strain induced by a lattice mismatch InGaP shell

Strain estimation

In Subsection 4.4.5, we showed that the growth of an InGaP passivation shell allows to reduce the SRV by at least one order of magnitude. But we also observed a clear shift of the CL peak at low temperature (Figure 4.28) of the GaAs-InGaP structure. This shift in the CL peak has been attributed to the strain imposed by a lattice mismatch InGaP shell. Indeed, growing a lattice mismatched shell causes a strain on the core and may change its bandgap [START_REF] Sköld | Growth and Optical Properties of Strained GaAs-Ga x In 1 -x P Core-Shell Nanowires[END_REF]. TEM experiments have been performed on the tips of GaAs-InGaP core-shell NWs. A quantitative EDX analysis have been performed on the shell in order to assess for the composition evolution along the NW axis. The chemical composition is assessed from local averaged EDX maps. We determined the quantity of Ga/In/P from the resulting EDX spectrograms. Then the lattice mismatch (∆) has been calculated from Equation 4.11 and plotted in Figure 4. 35 c). An important average lattice mismatch, over 1 %, has been determined determined for two NWs so we can observe two peaks whose ratio varies along the NW axis. A peak located at 1.435 eV, attributed to band to band transition, and a second peak shifted by 40 meV at 1.395 eV. Complementary temperature resolved CL experiments have been performed (not shown in this thesis). We observe that the difference in energy between the two peaks (estimated at 30 meV at LT) is constant as a function of the temperature. It suggests that deformation potentials and elastic constants exhibit no substantial temperature dependence. This behavior is consistent with the attribution [START_REF] Signorello | Tuning the Light Emission from GaAs Nanowires over 290 meV with Uniaxial Strain[END_REF].

Fits of the CL spectra using a two-Gaussian function have been performed for each CL spectrum of Figure 4.38c). The Gaussian energy position and σ have been fixed (un-strained phase:(E 1 = 1.435 eV , σ 1 = 22 meV ), strained phase:(E 2 = 1.395 eV , σ 2 = 17 meV )) in order to let only the intensities of the two 1-Gaussian functions as fitting parameters. Then the fitted intensity has been normalized by the maximum CL intensity for each spectrum. The evolution of the normalized low energy Gaussian intensity has been plotted in Figure 4.38 d). The evolution of the low energy contribution along the NW axis is plotted in Figure 4.38 d). We can observe that the intensity increases until 1.5 µm and then reaches a plateau. The trend is consistent with recombinations localized at the top of the NW for the low energy peak. A diffusion length of L n = 370 nm has been determined from a one-Exponential fit of the low energy peak intensity evolution.

As we measured a lifetime at RT (τ = 0.4 ns) for the exact same NW we can estimate the diffusion coefficient, from Equation 4.12, in the NW core: D n = 3.3 m 2 .s -1 . By assuming that Equation 4.13 can also be defined [START_REF] Sze | Physics of semiconductor devices[END_REF], where q is the elementary charge k the Boltzmann constant and µ n the mobility of electrons, in the case of a NW we can estimate the mobility in a NW. We can assume that our NW is lightly p-doped so that the determined lifetime and diffusion length correspond to the electrons lifetime (τ n ) and diffusion length (Ł n ).:

D n = L 2 n τ n (4.12) D n = k × q µ n (4.13)
Using Equation 4.13 and the experimental results we found µ n = 130 cm 2 .V.s -1 . The same order of magnitude for the carrier mobility in a passivated GaAs has been found by Chang et al. [30]. This value is one order of magnitude lower than GaAs bulk mobilities [START_REF] Blakemore | Semiconducting and other major properties of gallium arsenide[END_REF][START_REF] Stillman | Hall coefficient factor for polar mode scattering in n-type GaAs[END_REF].

From the evolution of CL spectra along the NW axis we measured a diffusion length in a GaAs-InGaP core-shell NW. We found a diffusion length of L n = 370 nm. By combining this result with the lifetime measurement performed by TRCL on the exact same NW we determined value for the diffusion coefficient in a single GaAs NW: D n = 3.34 m 2 .s -1 . From this value we estimated an electron mobility µ n = 130 cm 2 .V.s -1 . But because of lower dimensions, lower diffusion lengths and carrier mobilities are needed for a complete carrier collection in the case of the NW based solar cells.

General conclusion

In this thesis, I investigated the characterization of nanomaterials at the nanoscale using cathodoluminescence. Indeed, as introduced in Chapter 1, materials based on micro-or nano-structures such as grains in polycrystalline thin films and nanowire-based solar cells present an interesting solution for low-cost or high efficiency solar cells. We described in Section 3.1 and Subsection Background and Nanowires growth that their efficiencies are hindered by the very nature of nanomaterials: their small size, the detrimental role of surfaces, the presence of bulk defects, and difficulty of controlling the carrier concentration by doping. This work investigates these properties using the C2N cathodoluminescence setup. In Chapter 2 we have presented the specific features of the tool in order to probe the optical and transport properties at the nanoscale. In this thesis, this cathodoluminescence setup is used to develop novel measurement and analysis methods applied to CdTe and CdSeTe polycrystalline thin films and GaAs nanowires.

CdTe and CdSeTe polycrystalline thin films

In Chapter 3, I presented CL measurements of CdTe thin films annealed under CdCl 2 at different temperatures. I developed a statistical analysis in order to fully exploit the high-resolution panchromatic CL maps, together with a local analysis close to the grain boundaries. Hyperspectral low-temperature high-resolution CL maps were also measured on the same areas. As a result, we were able to explain the origin of the optimal CdCl 2 annealing temperature. This process induces a passivation of GBs, an increase of the grain size and an improvement of the radiative efficiency of grains for temperatures up to T opt = 420°C. Beyond this temperature, despite an increase of the grain size, lower radiative efficiencies are observed. This behavior was correlated with an increase of the defect density in GIs. Hence, the optimal temperature is the result of a trade-off between GB passivation, grain size, and the defect density in GI. Diffusion lengths were also estimated for each sample, and first results of time-resolved CL decays in single grains were measured at low temperature.

The introduction of Se in CdSeTe thin films provides a way to create bandgap gradients and to improve carriers collection in photovoltaic devices. It has also been shown to decrease the Voc deficit and increase the carrier lifetime. We have studied a series of CdSeTe samples prepared at NREL, with well-controled Se concentrations. I recorded high resolution RT and LT CL maps in bevels that allow an in-depth profile of the properties of CdSeTe. The analysis of hyperspectral maps allowed us to link the Se content to the radiative efficiency. We have shown the efficient passivation of defects with Se, resulting in an improvement of the radiative efficiency of one order of magnitude from CdTe to CdSe 0.4 Te 0.6 . These measurements were also correlated with macroscopic ones of solar cells (V oc and lifetimes).

Optical and transport properties in single GaAs NWs

In Chapter 4, I presented CL measurements of single GaAs NWs with distinct WZ and ZB segments, and the investigation of their optical and transport properties. LT CL maps under different injection levels and CL polarimetry were first performed in order to analyze the optical properties of these two crystalline phases. We evidence a strong polarization of light emitted by the WZ phase (degree of polarization P = -74%), in contrast with ZB (P = 5 %), enabling to distinguish unambiguously the two crystalline phases. We found a similar bandgap for both crystal phases, with a LT CL peak located at 1.513 eV, and we measured time-decays and found similar lifetimes for WZ (τ = 0.62 ns) and ZB (τ = 0.55 ns).

I also investigated the incorporation of Te as a dopant in GaAs NWs. We adapted a method for the quantitative determination of the electron concentration in Si-doped GaAs NWs developed at C2N. CL spectra at both RT and LT were measured and fitted with the generalized Planck law (explained in Chapter 2) on n-type reference thin films doped with Te. We confirmed the accuracy of the method in the case of the Te-doped GaAs layers. Then I applied this method to the determination of the electron concentration in Te doped NWs. An electron concentration up to n = 3.3 × 10 18 cm -3 was found. It confirms that the limitation of Si doping in GaAs nanowires, <1018 cm-3, can be overcome with Te doping. However, we also evidence a gradient of the effective Te doping along the nanowire.

Finally, we compared the passivation of GaAs NWs realized with high bandgap III-V shells: GaAs-AlGaAs and GaAs-InGaP. We first used a simple model to estimate the gain in voltage which can be achieved with surface passivation. Subsequently, RT and LT CL maps were performed in order to compare the radiative efficiency and the optical properties of the two core-shell structures, and time-resolved cathodoluminescence measurements provided additional information on the dynamics of charges and luminescence decay. We conclude on a higher RT radiative efficiency in the case of an GaAs-InGaP structure compared to an GaAs-AlGaAs structure. In this case of GaAs-InGaP NWs, we used the residual strain to determine the diffusion length (L = 370 nm), and time-decays measurements resulted in an average lifetime of about τ average = 0.31 ns. By coupling these measurements with a simple model, we estimated an upper value for the surface recombination velocity of S < 1.7 10 4 cm.s -1 , and approximated assessment of the diffusion coefficient (D = 3.34 m 2 .s -1 ) and an electron mobility (ţ = 130 cm 2 .V.s -1 ).

Conclusion and perspectives

The quantitative comparison of the luminescence emitted by nanostructures (grains or nanowires) has been performed in order to compare the evolution of the radiative efficiency as a function of the annealing CdCl 2 temperature for CdTe grains, or the passivation of GaAs NWs with high-bandgap shell. Additional transport properties measurements (diffusion or characteristic lengths) have been performed on these nanostructures, and correlated with the trend in the radiative efficiency. An absolute calibration of the CL setup, coupled with an estimation of the injection levels, is a perspective in order to determine an absolute value of the radiative efficiency, ratio of the radiative recombination rate to the total recombination rate.These measurements are currently under study at C2N by Bérengère Frouin (PhD started in 2019). Besides, this methodology can be transferred to other polycrystalline materials such as CIGS or perovskite, as implemented in a project started in 2019 (PhD of Salim Mejaouri).

In this thesis, I demonstrated that it was possible to measure time-decays in nanostructures using time-resolved CL. These measurements were performed at different locations of single NWs but also in single grains of a polycrystalline semiconductor, in order to determine the carrier lifetime. The combination of the lifetime estimation with the diffusion length enabled the determination of a diffusion coefficient and the electron mobility in at the NW scale. A long terms perspective would be to measure time-decays in single CdTe grains at room temperature. These measurements, coupled with grain size estimation would allow determining quantitatively both the grain boundary recombination velocity and the bulk lifetime. However, our TRCL setup 4.5. Summary is currently constrained by the repetition rate of pulses, limiting lifetime measurements below 10 ns, much below typical lifetimes measured by time-resolved photoluminescence on macroscopic thin films made of CdTe and CIGS (several 100 ns).

Hyperspectral maps at both RT and LT have been performed in order to study the evolution of the nature and distribution of defects within nanostructures. These measurements have been accompanied by injection-resolved CL and CL polarimetry. We have shown that a thorough analysis of the full luminescence spectrum can be used to determine quantitatively the density of defects in CdTe, or electron concentration (effective n-type doping) in single GaAs NWs, with a spatial resolution of the order of 100-300 nm.

Despite recent improvements, CdTe and CdSeTe thin films still suffer from low (p-type) doping levels. The method developed in GaAs can hardly be extended to CdTe due to the low concentration of holes. Injection-resolved CL coupled with simulation could provide an alternative way to a determine the concentration of majority carriers, using the different behavior expected at low and high injection levels. Besides, doping assessment in NWs is part of the continuous improvement process in order to develop efficient NW-based solar cells at C2N (PhD of Capucine Tong started in 2019). The CL polarimetry study performed on GaAs nanowires could also be extended to study the crystal properties (orientation, strain) in polycrystalline semiconductor layers such as CdTe, CdSeTe or CIGS. are seen as perspectives for this thesis.

The methodologies and techniques based on CL and TRCL developed during my thesis aim at being used to investigate, optimize and accelerate the understanding and development of photovoltaic devices based on nanomaterials.

Resumé en français Introduction

En 2016, 195 des 197 pays de l'ONU se sont rassemblés pour ratifier les accords de Paris dans le but de réduire le réchauffement climatique à 1.5°C d'ici 2100. Pour cela de nouvelles politiques énergétiques doivent être mises en place pour répondre à une demande croissante, notamment pour accorder une plus grande part aux énergies renouvelables. Ces dernières qui ne représentent que 26% de la production d'électricité en 2018. Parmi elles le photovoltaïque (PV), qui permet de produire de l'électricité à partir d'énergie lumineuse, a le potentiel pour jouer un rôle majeur dans la production d'électricité dans le futur. Cet engouement peut s'expliquer par une forte abondance de l'énergie solaire, un faible coût de production et une bonne prévisibilité de la production. Fin 2018 la capacité de production PV a dépassé 500 GW et 500 GW supplémentaires devraient être installés d'ici l'horizon 2022-2023 amenant la capacité de production mondiale d'électricité par le PV au delà du TW.

La technologie PV faite à partir de matériau Si est la plus utilisée, représentant plus de 95% de l'électricité produite par l'ensemble des modules PV. Leur efficacité de conversion avoisine 20% et leur prix décroît au cours du temps. Mais l'efficacité de ces cellules solaire en Si est limitée théoriquement en dessous de la barre des 30%. C'est pour cela que de nouveaux types de matériaux sont actuellement à l'étude pour produire des cellules solaires à plus hauts rendements. Parmi ces alternatives : les technologies dites de couches minces et les cellules tandem.

• Les dispositifs à couches minces sont faits à partir de matériaux semi-conducteurs fortement absorbant (CdTe, CIGS, perovskite...), permettant la réduction de l'épaisseur de la couche active. Ces matériaux sont crus avec des procédés peu onéreux grâce à leur faible température et à une haute pression de croissance. Du fait de ces conditions de croissance, des matériaux polycristallins sont obtenus présentant des grains de taille submicrométrique.

• Les cellules tandem correspondent à la juxtaposition de deux cellules solaires permettant l'optimisation de l'utilisation du spectre solaire. Ces technologies devraient se baser, à priori, sur la combinaison d'une cellule solaire avec une cellule Si, pour pouvoir bénéficier de la maturité technologique et industrielle de cette dernière. Les matériaux III-V semblent être de bons candidats pour l'élaboration de telles cellules, du fait notamment de leur très bonne qualité cristalline et de leur forte absorbance. Pour limiter les coûts de production, une croissance directe de la couche III-V sur Si doit être envisagée, mais cela est rendu complexe du fait du grand désaccord de maille entre les deux matériaux. Pour s'accommoder de cette contrainte la croissance sous forme de nanofils est explorée, ce qui permettrait la croissance directe de III-V sur Si avec une très bonne qualité cristalline.

Ces deux technologies sont des pistes prometteuses pour produire des dispositifs PV à hauts rendements mais leur efficacité est limitée par la faible qualité des nanostructures (nanofils et grains) qui cause un fort déficit en V oc . Or la caractérisation de celles-ci est complexifiée du fait de leur taille nanoscopique. Ce travail de thèse a pour objectif de caractériser les grains de CdTe et les nanofils en GaAs en utilisant la cathodoluminescence et de développer des méthodes d'analyse basées sur cet outil. La cathodoluminescence est l'émission de lumière par un solide soumis à un bombardement d'électrons. Ce phénomène permet de mettre en évidence les variations spatiales des propriétés optiques locales d'un matériau grâce à la taille nanoscopique de la sonde électronique. Dans cette thèse, nous avons exploité les performances du système de cathodoluminescence du C2N pour caractériser les nanostructures que sont les grains de CdTe et les nanofils en GaAs.

Nous détaillerons dans la première partie le contexte de cette thèse et les objectifs poursuivis (Chapitre 1), puis nous présenterons les spécifications et les principaux atouts de l'installation de cathodoluminescence au laboratoire C2N (Chapitre 2), auquel a été ajouté un protocole de calibration de la machine. Nous avons ensuite utilisé les atouts de cette machine pour caractériser dans un premier temps des films fins polycristallins en CdTe (Chapitre 3) et dans un second temps des nanofils en GaAs (Chapitre 4).

Films fins polycristallins en CdTe

Dans ce Chapitre nous avons étudié des couches minces en CdTe qui ont été préparés au NREL. L'amélioration continue de ces films fins polycristallins en CdTe a mené à l'établissement d'un record de 22.1%. Cependant, ceux-ci présentent toujours des efficacités qui sont assez éloignées de leur limite théorique avec notamment un déficit en V oc de 630 meV, comparé à un déficit de 280 meV dans le cas des meilleures cellules en GaAs. Cela peut être expliqué par la nature polycristalline de ces films où les joints de grains (zone à forte densité de défauts séparant deux grains) jouent un rôle encore mal compris. La complexité de telles nanostrucutures ainsi que leurs tailles rendent difficiles leur caractérisation et leur modélisation.

Le recuit en CdCl 2 est utilisé par l'industrie car il permet d'améliorer les efficacités des films fins en CdTe de manière drastique. Ce procédé est couramment utilisé et a été optimisé de manière phénoménologique. Nous avons donc tenté d'expliquer l'effet exact de la température de ce recuit sur les grains et joints de grains de CdTe. Pour cela nous avons comparé quantitativement l'intensité de CL de divers films fins recuits sous atmosphère de CdCl 2 à différentes températures (Figure 4.39). On a pu, dans un premier temps, remarquer une augmentation de la taille de grains avec la température de recuit. Or les grains les plus brillants sont observés pour l'échantillon à 420 °C, qui ne correspond pas à l'échantillon avec les plus gros grains. Une diminution de l'efficacité radiative des intérieurs de grains s'opère donc pour de trop grandes températures de recuit. Une analyse statistique des cartographies panchromatiques à température ambiante nous a permis de mettre en lumière le rôle joué par la température de recuit sous CdCl 2 , et de séparer son impact sur les joints de grains et sur les intérieurs de grains. Nous avons donc pu montrer que le recuit sous CdCl 2 permet une passivation des joints de grains jusqu'à une température de 420 °C. Au dessus de cette température, la passivation des joints n'est pas améliorée mais l'efficacité radiative des grains semblent diminuer. Cette évolution de l'efficacité radiative des intérieurs de grains a été étudiée par une étude locale de l'évolution de l'intensité de CL aux abords d'un joint de grains. Cette étude a permis d'estimer des longueurs caractéristiques de diffusion (Figure 4.39 k)), pour chaque échantillon, que l'on a corrélées avec l'évolution de l'efficacité radiative observée via l'étude statistique.

Le comportement de l'efficacité radiative de films fins de CdTe en fonction de la température de recuit en CdCl 2 a été expliqué par des mesures de cartographies de CL hyperspectrales à basse température. Ces cartographies ont été acquises aux mêmes endroits où ont été mesurées les cartographies à température ambiante. Cela a permis de démontrer une diminution de la densité de défauts, jusqu'à 420 °C, puis une augmentation de celle-ci. Nous avons pu par ailleurs montrer que la densité de défauts augmentait dans les intérieurs de grains. La corrélation des données à température ambiante (statistique et locale) et basse température a donc montré que le recuit sous CdCl 2 contribue à la passivation des défauts aux joints de grains et intérieurs de grains jusqu'à une température de 420 °C. Au delà de cette température la taille de grains continue d'augmenter, mais une diminution de la luminescence est causée par la formation de défauts dans l'intérieur de grains.

Ensuite nous nous sommes intéressés à la seconde problématique qui est l'ajout de Se dans leur structure. Cet ajout cause un décalage vers le rouge de la bande interdite du CdSe x Te 1-x . Cela a permis d'augmenter le J sc de manière significative, mais de manière surprenante cela n'a pas diminué le V oc . Cet effet a été expliqué par une passivation des défauts dans le volume et aux des joints de grains. Pour analyser ce phénomène nous avons étudié des films fins en CdSe x Te 1-x avec différentes compositions en Se, fabriqués au NREL. Des biseaux ont été préparés pour l'analyse de l'évolution, dans l'épaisseur du film, des efficacités radiatives et des propriétés optiques. Des mesures de CL hyperspectrales hautes résolutions à température ambiante ont été réalisées et ont permis d'établir des profils d'évolution de l'intensité et de la position du pic de CL le long de l'épaisseur du film (Figure 4.40 a)b)). On a pu mettre en évidence un décalage vers le rouge de ce dernier avec l'augmentation de la proportion de Se, ainsi qu'une augmentation de l'efficacité radiative avec cette même proportion. Une corrélation entre la position et l'intensité du pic de CL a été établie (Figure 4.40 c)) pour chacun des échantillons. Une évolution commune, de l'intensité de CL en fonction de la position du pic, pour tous les échantillons a été mies en évidence, et ce malgré les variations de tailles de grains entre les échantillons. Par ailleurs, des mesures complémentaires de temps de vie ont été réalisées et ont démontré les très longs temps de vie obtenus sur ces structures CdSe x Te (1-x) (τ (x Se = 0.4)) > 100ns). Cela a permis de conclure quant à la passivation effective des défauts dans les intérieurs de grains avec l'introduction de Se dans les structures de CdTe.

La mesure des temps de vie effectifs couplée avec une mesure de la taille de grain permettrait en effet la détermination de la vitesse de recombinaison aux joints de grains et du temps de vie d'intérieur de grains. Des mesures préliminaires de cathodoluminescence résolue en temps (TRCL) dans des grains uniques de CdTe ont par ailleurs été effectuées durant cette thèse. Ces mesures ont été réalisées à basse température (10K) et ont permis de mettre en évidence un faible temps de vie excitonique (τ X = 35 ps) et un long temps de vie des défauts (τ def = 12 ns). Du fait du très long temps de vie à température ambiante de ces structures, la mesure de temps de vie effectif à température ambiante n'est pas réalisable avec notre système de mesure.

Nanofils en GaAs

Dans ce chapitre nous avons étudié des nanofils (NFs) de GaAs synthétisés par épitaxie par jet moléculaire (MBE) sur substrat de Si pré-patterné via la technique de croissance vapeur-liquidesolide (VLS) auto-catalysée. Malgré les nombreux travaux sur l'optimisation des paramètres de croissance de ces NFs, les performances de ces dispositifs sont notamment limitées par: la qualité cristalline, le contrôle du dopage et la passivation de la surface de ces NFs. Pour réaliser des dispositifs efficaces il est nécessaire de concevoir des NFs avec de forts taux de dopage, or l'incorporation de dopants dans le cas des NFs diffère du cas des films fins. Des mécanismes de compensation dans le cas de la croissance de NFs ont notamment été mis en évidence qui n'étaient pas observés dans le cas de la croissance de films fins. Par ailleurs la mesure du dopage est complexifiée du fait de la difficulté de la fabrication de contacts électriques dans le cas de nanostructure, rendant difficile son amélioration continue. Une méthode de mesure quantitative par CL du dopage au sein d'un NF de GaAs dopé n (Si) a été développée par notre équipe au C2N. Cette méthode se base sur l'analyse des spectres CL ajustés à la loi de Planck généralisée. Des relations entre les caractéristiques des spectres de luminescence (E g ,E f c ...) et les concentrations de porteurs ont été établies en utilisant des couches minces de références. Dans cette thèse nous avons dans un premier temps détaillé cette méthode puis montré qu'il était possible de la transposer au dopage n par un dopant Te. Dans un premier temps les relations entre les différents paramètres ont été vérifiées, dans le cas du dopant Te, par la mesure de spectres CL haute et basse température sur des couches minces en GaAs dopées crus à l'Université de Tampere (Finlande). Puis, une fois l'applicabilité de la méthode démontrée, nous avons utilisé ces résultats pour mesurer à l'échelle nanométrique la concentration en électrons dans un nanofil unique et son homogénéité au sein de ce même NF (Figure4.42 e)). Une concentration en électrons de n = 3.3 × 10 18 cm -3 a été mesurée dans le haut du nanofil par cette méthode A cause de leurs faibles dimensions, les dispositifs à NFs sont très dépendants des effets de surface. Dans le cas de NFs de GaAs une forte vitesse de recombinaison de surface est observée (≈ 10 5 cm.s -1 ). On peut montrer qu'une réduction de deux ordres de grandeurs de la vitesse de recombinaison de surface S permet d'allonger le temps de vie des porteurs de plusieurs ordres de grandeurs (Figure 4.43 a)). En se basant sur ce résultat on peut estimer qu'une telle passivation de la surface du NF entraine un gain en tension de ∆V = 150 meV . Pour réaliser cette passivation de surface, des coquilles à large bande interdite peuvent être utilisées. Nous avons donc comparé deux coquilles de passivation, l'une réalisée avec une couche de AlGaAs ( . Ces mesures ont été complétées par des mesures résolues en temps à basse température. Un plus long temps de vie a été mesuré pour la structure AlGaAs, en désaccord avec l'intensité de luminescence. Mais du fait d'une luminescence plus faible à température ambiante il n'a pas été possible de réaliser des mesures de temps de vie sur la structure GaAs-AlGaAs, au contraire de celle en GaAs-InGaP pour laquelle des mesures de temps de vie entre 190 ps et 400 ps ont été mesurés. Grâce à ces mesures nous avons pu estimer une borne supérieure de la vitesse recombinaison de surface qui est de S < 1.7 × 10 4 cm.s -1 . Grâce à l'analyse hyperspectrale de cartographies de CL mesurées sur les NFs GaAs-InGaP, une mesure de la longueur de diffusion en complément d'une mesure du temps de vie a été réalisée, aboutissant à L n = 370 nm. En supposant un comportement volumique des NFs nous avons pu estimer un coefficient de diffusion D = 3.3 m 2 .s -1 et une valeur de la mobilité des électrons µ n = 130 cm 2 .V.s -1 .

Conclusion

Les outils et méthodes mis en place lors de nos diverses études, présentées dans cette thèse, ont été basées sur les spécificités de la machine de CL installée au C2N. Nous avons montré comment grâce au développement de méthodes de mesure basées sur la CL il était possible de mesurer les propriétés optiques et de transport dans des nanostructures comme les grains de CdTe ou les nanofils en GaAs. L'affinement au cours de la thèse des techniques de mesure a permis d'améliorer notre compréhension des effets qui ont lieu dans ces nanostructures. Ces méthodes et techniques ont aussi pour but d'être adaptables à l'étude de divers matériaux pour accélérer la compréhension et l'amélioration continue de dispositifs photovoltaïques à haut rendement.

Dans cette thèse les caractéristiques de volume ont été étudiées dans le cadre de l'évolution de celles-ci avec la température de recuit en CdCl 2 et de l'ajout de Se dans les grains de CdTe, ou de l'alternance des phases de ZB et de WZ dans les NFs en GaAs. De même les effets de surface ont été analysés via la comparaison quantitative de la luminescence émise entre différents films en CdTe ou dans des NFs de GaAs avec différentes coquilles de passivation. Par ailleurs ces mesures ont été complétées par des mesures résolues en temps permettant la mesure de temps de vie effectif dans des nanostructures. La haute résolution spatiale de telles mesures a aussi été exploitée dans le but de caractériser l'incorporation de dopants dans des NFs uniques.

Comme perspectives à court terme, l'analyse des mesures sur les couches de CdSeTe devrait être une priorité dans le but d'élucider la nature des défauts passivés notamment grâce à une meilleure compréhension des cartographies de CL hyperspectrales à basse température. Ces travaux font notamment parties de la thèse démarrée par Bérengère Frouin en 2019. L'amélioration de la concentration en porteurs libres dans les films fins en CdTe est une question majeure qui reste à explorer. Nous avons montré comment la cathodoluminescence pouvait permettre ce type d'étude dans le cas du GaAs. Par ailleurs, la mesure de temps de vie ainsi que des mesures de polarimétrie de CL sont envisagées sur les couches minces en CdTe pour mieux comprendre la nature des défauts dans de telles structures. Les méthodes de caractérisation mises en place pour les NFs seront intégrées dans l'amélioration continue de cellules solaires solaires à base de NFs, projet démarré par Capucine Tong en 2019. L'extrapolation de ces techniques sur d'autres matériaux, comme les perovskites, sont aussi l'objet d'études comme notamment le projet de thèse démarré par Salim Méjaouri en 2019.

Figure 1 . 1 :

 11 Figure 1.1: Observed monthly global mean surface temperature compared to 1850-1900 (gray solide line). The orange dashed arrow and horizontal orange error bar show respectively the central estimate and likely range of the time at which 1.5 °C is reached if the current rate of warming continues. The plumes represent the response in temperature following different pathways for CO 2 emissions reductions policies. Vertical bars on the right present the likely ranges of the estimated warming in 2100 for the three considered pathways. Figure reprinted from:[START_REF] Masson-Delmotte | Summary for Policymakers[END_REF]

Figure 1 . 2 :

 12 Figure 1.2: Reasons for concern summarizing the key impacts and risks across sectors and regions. It illustrates implications of the global warming for ecosystem, economy and people. Impacts and risks are leveled for different global mean surface temperature changes. Colors represent the degree of risks and impacts of the warming: White (undetectable), Yellow (moderate), Red (high) and Purple (very high).Figure reprinted from:[START_REF] Masson-Delmotte | Summary for Policymakers[END_REF]

Figure 1 . 3 :

 13 Figure 1.3: Estimated shares in the total energy consumption in 2017 according to OECD/IEA. Figure reprinted from:[START_REF]renewables 2019[END_REF]

Figure 1 . 4 :

 14 Figure 1.4: Shares of renewable energy production in net annual additions of power generating capacity between 2008 and 2018. Figure reprinted from:[START_REF]renewables 2019[END_REF]

Figure 1 . 5 :

 15 Figure 1.5: Evolution of the module price as a function of the cumulative PV power installed. Reprinted from:[START_REF]Photovoltaic report[END_REF]

Figure 1 . 6 :

 16 Figure 1.6: Repartition in percentage of the annual production of modules between the main technologies of the PV market between 1980 and 2017. Figure reprinted from[START_REF]Photovoltaic report[END_REF]

Figure 1 . 7 :

 17 Figure 1.7: Solar spectral irradiance (AM1.5G) after the transmission through the atmosphere at 45 °(gray curve). Part of the spectrum effectively collected Si solar cell (green curve). Figure reprinted from[104]

Figure 1 . 8 :

 18 Figure 1.8: a) Scheme of a simplified solar cell structure. b) Band diagram of a p-n junction solar cell. The carrier collection process consists of three steps: A) photon absorption B) the thermalization of carriers directly after the photogeneration and C) carrier collection at selective contact's. Figure b) is reprinted from: [69].

Figure 1 . 11 :

 111 Figure 1.11: Current density-voltage curves taking into account the successive loss process: (1,2) ideal collection and absorption but thermalization process occurs, (4) radiative losses with non ideal absorption (5) non radiative losses are taken into account. For each J-V curve, a rectangle represents the maximum power achievable. The final white square correspond to the model of an actual solar cell taking into account radiative and non-radiatives losses, non ideal absorption, thermalization processes and resistive losses. Figure reprinted from:[START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF]

Figure 1 . 12 :

 112 Figure 1.12: Equivalent electrical circuit model of a solar cell corresponding to Equation 1.4. Dark currents are represented using two diodes (two-diode model): J 02 corresponds to non radiative recombinations processes or SRH recombination and J 01 corresponds to radiative recombination processes. Two resistances are also added in order to take into account series resistance (R s ) and the shunt resistance (R p ). Figure reprinted from:[START_REF] Guillemoles | Guide for the perplexed to the Shockley-Queisser model for solar cells[END_REF]

Figure 1 . 13 :

 113 Figure 1.13: Energy band diagram showing a selection of radiative and non radiative recombination channels: 1) Band-to-band 2) mid-bandgap defects mediated non radiative recombination 3) 4) and 5) Defects mediated radiative recombinations.

Figure 1 . 14 :

 114 Figure 1.14: scheme of a recombination via surfaces states continuously distributed within the SC bandgap. Figure reprinted from[START_REF] Würfel | Physics of solar cells: from basic principles to advanced concepts. 2nd[END_REF]

Figure 1 . 15 :

 115 Figure 1.15: Cross section scheme of a Si solar cell with an heterojunction and an interdigitated back contact HJ-IBC. Figure reprinted from[START_REF] Yoshikawa | Silicon heterojunction solar cell with interdigitated back contacts for a photoconversion efficiency over 26%[END_REF]

Figure 1 . 16 :

 116 Figure 1.16: Scheme of classic stacks for thin-film technologies a) CdTe b) perovskite and d) CIGS.Reprinted from[151]

Figure 1 . 17 :

 117 Figure 1.17: PCE efficiency comparison of pc-TF and Si technologies for cells and modules record efficiency. Figure reprinted from[START_REF]Photovoltaic report[END_REF].

Figure 1 . 18 :

 118 Figure 1.18: Detailed balance calculated efficiency for a 2-junction tandem cell under 1-sun illumination as a function of the top and the bottom cell bandgap. Si bottom cell bandgap has been set and the interception line has been plot (white dashed line). The white circle presents the maximum efficiencywhich can be achieved with a bottom cell Si with a E g =1.15eV. This figure has been reprinted from[64]

Figure 1 .

 1 Figure 1.19: a) Bandgap of SC materials as a function of the lattice constant. Filled circles correspond to direct bandgap and Un-filled circles correspond to indirect bandgap. b) Calculated critical thickness for a disk materials with different lattice mismatch. This figure has been reprinted from[133].

Figure 1 .

 1 Figure 1.20: a) SEM image of a GaAs nanowire based solar cells on a GaAs bottom cell fabricated by Aberg et al. [1] with a 15.3 % efficiency b) scheme of the GaAs nanowires based structure developed by Yao et al. [180] with a 11.4 % efficiency.

Figure 1 .

 1 Figure 1.21: a) Schematic illustration of the Hall effect in a SC thin layer. b) Scheme of the four contacts required for measuring thin-film doping concentration by Hall effect. c) Illustation of the contact deposition achieved by Blömerset al. [51] for the hall effect measurement in single InAs NW. Figure reprinted from a)b)[32] and c)[START_REF] Ch | Hall effect measurements on InAs nanowires[END_REF]

Figure 1 .

 1 Figure 1.22: a) scheme of the characterization tool used by Ketterer et al. [91] where a light beam has been focused on different part of a single NW. The red arrow on the top of the light beam indicate the polarization of the light incoming on the NW. b) Raman spectrum, recorded on a single GaAs NW, exhibits multi longitudinal optical (LO) phonons peaks. The Raman signal is superimposed with a photoluminescence background. c) GaAs band diagram where the resonant raman LO processes are displayed. Figure reprinted from [91].

Figure 1 . 23 :

 123 Figure 1.23: Scheme of a hyperspectral PL setup with a LASER (λ=532nm) source (green area). A beam splitter is added in the light path in order to send the PL signal to a CCD camera and to remove the signal due to the LASER. Figure reprinted from [44].

Figure 1 .

 1 Figure 1.24: a) SEM image of a GaAs NW. b) Photoluminescence spectra recorded on NWs similar as the one presented in a) at different location of the NW. Figure reprinted from[30]

Figure 1 . 25 :

 125 Figure 1.25: Dark field TEM micrograph of a single GaAs-AlGaAs NW (76715). TEM diffraction patterns have been collected at the base (blue), the middle (green) and the top (red) showing respectively pure Zinc-Blende (ZB), a mixed and a pure Wurtzite (WZ) phases. Inset shows a high annular angle dark field image of the top interface between pure ZB and pure WZ phases. The lighter part of the NW, observed in the inset, corresponds to the AlGaAs shell.

Figure 1 . 26 :

 126 Figure 1.26: Scheme of the CL principle where an high energy incident electron generate 1) secondary electrons and/or 2) electron-hole pairs. 3) Luminescence generated by radiative recombinations between excess electron-hole pairs. 4) When contacted collection of excess carriers result in an Electron-Beam-Induced-Current signal (EBIC).

Figure 2 . 1 :

 21 Figure 2.1:Cathodoluminescence setup scheme where four ensembles have been distinguished: the emission system (grey), the light collection optics (yellow), the detection and dispersion system (red) and the time-resolved system (green). The different system have not been drawn at scale.

Figure 2 . 2 :

 22 Figure 2.2: Schematic of the Attolight Allalin 4027 Chronos set up. The setup scheme has been simplified, compared to Figure 2.1, with a single spectrometer and a more detailed sketch of the collection optics and the nanopositionning system (see Subsection 2.1.4). Figure reprinted from: [9].

Figure 2 . 3 :

 23 Figure 2.3: Evolution of the incident current as a function of the gun lens for different apertures: 100µm (red), 50µm (blue), 30µm (green). The current is measured by averaging EBIC maps of 128×128 pixels made with a Faraday cup. The acceleration voltage of the electrons during the experiment was 6kVand a 10 nm gun probe size.

Figure 2 . 4 :

 24 Figure 2.4: Spectral responses of the CCD Andor camera (blue) and the grating 150/500 blazed at λ=500nm (red) determined from products datasheets. The spectral response of the system (red) is associated to the product of the two spectral responses.

Figure 2 . 5 :

 25 Figure 2.5: Description of the scanning areas. The sample (red square) images are shown at a) the SEM plan, b) the entrance slit and c) the CCD plan. d) and e) represents CL spectra at different location O (black) M (blue) and M'(green) of the hyperspectral map when d) no correction is applied and e) if the spatial shift is perfectly corrected.

Figure 2 . 6 :

 26 Figure 2.6: Peak position maps of CL spectra of a GaAs thin film. a) As measured peak position map, no spatial correction (Equation 2.2) has been applied, a blue-shift of about 13.6 meV is observed. b) The same map as the one presented in a) when the spectral correction is applied.

Figure 2 . 7 :

 27 Figure 2.7: Data processing steps : a) raw CL spectrum where the level of the dark is determined by averaging over low or high wavelength spectral ranges. b) CL spectrum where the noise level has been removed. Then the CL signal is corrected by the spectral responses of the CCD and the grating. The different signals which have been obtained, the un-corrected (red line) and the corrected (black line), have been normalized and plotted as a function of the wavelength c) and of the energy d).

Figure 2 . 8 :

 28 Figure 2.8: Comparison of the different time-resolved cathodoluminescence configurations. The standard configuration corresponds to the continuous excitation and is compared to the fast blanker and the photoemission configuration. Figure reprinted from [124].

Figure 2 . 10 :

 210 Figure 2.10: Streak image collected on a reference sample (A2033) at 70K. Color scale has been adjusted to improve the contrast.

  Figure 2.11: TRCL decay measured on a short lifetime sample (A2033) with a streak camera (red curve). The decay has been fit (blue curve) using the convolution of a Gaussian function (FWHM = 12 ps) with an exponential decay (τ =8ps) in order to estimate the time response of the system. The time scale has been shifted by 400 ps compared to Figure2.10.

Figure 2 . 13 :

 213 Figure 2.13: Electrons trajectories estimated from CASINO simulations for a a) 3kV and b) 6kV acceleration voltage for a bilayer structure with a 75nm AlGaAs layer on a GaAs substrat. Red lines describe back-scattered electrons. The scale was kept constant to compare the penetration depth between the two acceleration voltages.

Figure 2 . 14 :

 214 Figure 2.14: Generation distributions (schematic) a) for point source, b) spherical generation and c) Gaussian distribution function.

Figure 2 . 15 :

 215 Figure 2.15: Normalized luminescence spectrum of a n-doped GaAs thin film (black dots) fitted (red curve) with the generalized Planck law (Equation2.12). The fitted bandgap is shown with a red circle. Besides, exponential evolution of the low and high energy sides of the spectrum have been evidenced.

Figure 2 .

 2 Figure 2.16: a) Cross section view of the full GaAs stack. b) Scheme of the patterns used for the fabrication process of the calibration diode. Red inset presents the full patterns where the diodes (as presented in the yellow inset) with different sizes appear. Back contact are also evidenced (green inset). c) Photo of the CL holder where the pattern and the grooves position have been informed by arrows.

  Calibration diodes have been elaborated. Those diodes were fabricated from ultrathin layers of GaAs furnished by the Fraunhofer (Figure 2.16 a)) thus emitting around 1.41 eV. To maximize the number of workable diodes several patterns were fabricated. Each pattern contains 12 diodes surrounded by a front contact and 2 back contacts (see Figure 2.16 b)). The diameter of the diodes vary from 30 µm to 90 µm. During the experiments a 30 µm diameter diode with a 300 µm wide front contact has been used (Yellow inset in Figure 2.16 b)).Then front and back contact are contacted by micro-bounding with the holder grooves. These grooves are connected to the electrical output on the front door of the CL chamber which have been connected to a Keithley.

Figure 2 .

 2 Figure 2.17: a) Panchromatic electroluminescence map and b) SEM image collected with the PL system for the R=15µm GaAs diode presented in Figure 2.16.

I

  0 cos(θ)sin(θ)dθdφ = πI 0(2.18)

Figure 2 .

 2 Figure 2.18: a)-g) EL image of the GaAs diode recorded with the CCD of the C2N CL setup, when the spectrometer central wavelength was set at λ=875nm, with different spectrometer entrance slit opening: a) 7000µm b) 5000µm c) 4000µm d) 3000µm e) 2000µm f) 1000µm g)800µm. h) Electroluminescence spectra collected by averaging the CCD images presented in a)-g).

Figure 2 . 19 :

 219 Figure 2.19: Total intensities collected by the CCD as the function of the diode area proportion intercepted (%) (blue dots). Data are fitted with a linear function (red line).

Figure 2 . 20 :

 220 Figure 2.20: scheme of a lambertian emitter characterized by a maximum intensity I 0 for deviation from the normal θ=0°. Collection efficiency is described by angles α 1 and α 2 .

Figure 2 . 21 :

 221 Figure 2.21: Total intensity collected by the CCD as a function of the LASER power(blue dots). Data have been fitted by a linear function (red line).

Figure 3 . 1 :

 31 Figure 3.1: Record efficiencies evolution over the years of polycrystalline (pc) CdTe, CIGS and Si solar cells since 1987. Figure reprinted from[151]

Figure 3 . 3 :

 33 Figure 3.3: 3D scheme of a polycrystalline layer where a grain (green) has been highlighted. Grain boundaries(red) have been represented as red lines. Besides, top and bottom interfaces (blue) have been also colorized for clarity. Characteristic material properties for each family have been adressed.

Figure 3 . 4 :

 34 Figure 3.4: Theoretical efficiency as a function of grain boundary recombination velocity (S GB ) and the bulk lifetimes,1 ns (circles), 10 ns (squares), and 100 ns(triangles),with S int =100 cm.s -1 and carrier concentration of p=2 ×10 16 cm.s -1 for various grain sizes: 1 µm (green), 3 µm (red), and 20 µm (black).To compare, S int =10 5 cm.s -1 and 20µm grains is shown with dashed lines. This is Figure is reprinted from[START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grainboundary recombination for 25% efficient CdTe solar cells[END_REF]

Figure 3 . 5 :

 35 Figure 3.5: a) 2D model of a polycrystalline cell with the various parameters used to characterize the different interfaces b) Efficiency as a function of the front surface recombination (S front ) and the back surface recombination velocity (S b ) for a CdTe thin film with a two doping levels of p= 2×10 14 cm -3 and p=2×10 16 cm -3 . b) is reprinted from [50].

Figure 3 . 6 :

 36 Figure 3.6: a) cross-section EBSD of CdTe thin films deposited on a MZO front contact layer with different CdCl 2 annealing temperature. b) contour plot of theoretical V oc and efficiency as a function of the grain boundary recombination velocity S GB and the grain size. Images are respectively reprinted from a) [7] and b)[START_REF] Kanevce | The roles of carrier concentration and interface, bulk, and grainboundary recombination for 25% efficient CdTe solar cells[END_REF]

Figure 3 . 7 :

 37 Figure 3.7: Cathodoluminescence (CL) and Electron Back Scaterring Diffraction (EBSD) maps on the same area of a as-deposited (first row) and CdCl 2 treated CdTe polycrystalline thin films. Grains boundaries on EBSD are represented by continuous lines. Both CL and EBSD maps are very alike and dark areas on CL maps correspond to EBSD continuous lines[START_REF] Moseley | Recombination by grain-boundary type in CdTe[END_REF].

Figure 3 . 8 :

 38 Figure 3.8: Scheme of a preparation method of CdTe : The close space sublimation (CSS). A CdTe crucible is heated in order to deposited a CdTe layer on the Glass substrate. Figure reprinted from[151]

Figure 3 . 9 :

 39 Figure 3.9: CL Panchromatic maps (128×128 pixels) of CdTe samples with different CdCl 2 post deposition treatment a) No CdCl 2 b) 400°C c) 420°C d) 440 °C e) 460°C. All images have been normalized on the same intensity scale. (scale bar: 5µm).

Figure 3 . 10 :

 310 Figure 3.10: Room temperature CL spectra determined by averaging the 128×128 pixels of CL maps for the No CdCl 2 , 400°C, 420°C, 440°C and 400°C samples.

Figure 3 . 11 :

 311 Figure 3.11: Average grain size calculated for each sample from CL maps at RT. An average grain area is first determined then the corresponding area is transposed to a circle then a grain size corresponding to the diameter of the circle is calculated. Error bars correspond to the standard deviation of the estimated diameters for each sample.

Figure 3 . 12 :

 312 Figure 3.12: Room temperature intensity histograms for each sample with a two-gaussian function fit for a) No CdCl 2 b) 400°C c) 420°C d) 440°C and e) 460°C. Histograms have been determined from panchromatic CL maps showed in Figure 3.9

Figure 3 . 13 :

 313 Figure 3.13: Evolution of the intensities mean values I GI and I GB of the grain boundaries and grain interiors as described in Table 3.3. Plotted errors bars correspond to standard deviations σ GI and σ GB .

Figure 3 . 14 :

 314 Figure 3.14: False color room temperature CL map where points were divided between three segments low intensity (blue) [0:I GB +σ GB ], high intensity (green) [I GI -σ GI :∞] and an intermediate family (red) ]I GB +σ GB :I GI -σ GI [

Figure 3 .

 3 Figure 3.15: a) Panchromatic CL intensity map extracted from the untreated sample. Grain 1 and Grain 2 correspond to randomnly chosen grains b) CL intensity plot as a function of the position with 1D fits c) L charac as a function of the CdCl 2 annealing temperature. Error bars correspond to standard deviation of the 8 L charac estimated for each sample.

3 )IFigure 3 . 16 :

 3316 Figure 3.16: a) low temperature CL spectrum determined by averaging the 128×128 pixels CL map for the No CdCl 2 sample.b) False color (128x128 pixels) CL map where contributions for each spectra bands Z-defects (blue) A-defects (red) exciton (green). Contribution corresponds to ratio of the integrated intensity of the spectral band over the integrated intensity of the full spectrum.

E 2 Se( 3 . 6 )Figure 3 .

 2363 Figure 3.18: a) typical CdSe x Te 1-x solar cell structure. The Se grading in the absorber is represented by a color grading from orange (Se rich) to gray (pure CdTe) region. b) Evolution of the bandgap with the Se content of thin films grown by MBE. Figure are reprinted from a) [131] b)[START_REF] Brill | Molecular beam epitaxial growth and characterization of Cd-based II-VI wide-bandgap compounds on Si substrates[END_REF]

Figure 3 .

 3 Figure 3.19: a) Scheme of the CdSeTe/CdTe device structure and bevel.It includes an outline where the SIMS and CL measurement area have been performed (dotted black line). b) A SIMS elemental map of the selenium distribution on the beveled surface of the CdCl 2 -treated device with the 'skeletonized' chlorine signal overlaid in white to help delineate the grain boundaries. c) Panchromatic CL map taken on the same area as the SIMS selenium map. d) A high-magnification Se SIMS and panchromatic CL maps of a region at the top of the bevel. Similarity in the positioning of the selenium and CL signals is outlined by the dashed lines in d) and e). f) A scatter plot of the Se concentration versus CL counts (for equivalent regions in b) and c). Figure reprinted from[START_REF] Thomas | Understanding the role of selenium in defect passivation for highly efficient selenium-alloyed cadmium telluride solar cells[END_REF]

Figure 3 .

 3 Figure 3.20: a) Cross section scheme of Glass/SnO 2 /MZO/CdTeSe/CdTe stack with a bevel b) Top view SEM image of CdSeTe Stack with a bevel, scheme on the right describes the correspondence of the difference SEM contrast value.

Figure 3 . 21 :

 321 Figure 3.21: Room temperature panchromatic CL maps (256 × 256 pixels) of beveled CdSe x Te 1-x with a) x=0 (only-CdTe) b) x=0.1 c) x=0.2 d) x=0.3 e) x=0.4. Because of drift artefacts, images b)d)e) have been realigned by aligning the CL intensity constrat. The grayscale bars have been set for each CL intensity map in order to optimize the images contrast. (white scale bar: 5 µm).

Figure 3 . 22 :

 322 Figure 3.22: Profiles of the CL intensity and the peak position along the CdSe x Te 1-x thickness with various x Se . The Se rich region (orange) and Se poor region (blue) are also plotted.

Figure 3 . 23 :

 323 Figure 3.23: Evolution of the CdSe x Te (1-x) bandgap as a function of the Se content (x Se ). CL peak position are assimilated as a first order approximation to the bandgap (gray dots). Error bars are estimated from the FWHMs summarized in Table 3.5. Fit of the evolution of the bandgap (orange solid line) with a Equation 3.6 with a bowing factor b=0.68.

Figure 3 . 24 :

 324 Figure 3.24: Room temperature CL intensity as a function of the CL peak position estimated from 128×128 pixels hyperspectral CL maps on CdSe x Te 1-x samples with Se content ranging from 0.1 to 0.4.

Figure 3 .

 3 Figure 3.25: a)-e) Low-temperature (20K) average CL spectra for CdSe x Te 1-x samples annealed with different Se compositions. Spectra are extracted from LT CL maps (256×256pixels). f)-o) CL false color maps (256×256 pixels) at low-temperature (20K) where normalized luminescence for each spectra bands defects (blue) CdSeTe-exciton (red) CdTe-exciton (green) are plotted. CdSeTe-exciton and CdTe-exciton maps were superposed for a better comparison of the change in composition within the film. Normalization factors have been determined from the maximum integrated luminescence from each spectral bands. White scale bar: 2 µm

Figure 3 .

 3 Figure 3.26: a) Characteristic single photon counting curve measured on a solar cell with 30% Se composition. b) The correlation of the effective lifetime versus V oc 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 3 .

 3 26 b) illustrates that the lifetime increases by orders of magnitude with increasing Se concentration until reaching a value of x=0.3. In average, the measured lifetime decreases as the Se content increases from x=0.3 to x=0.4.

Figure 3 .

 3 Figure 3.27: a) Scheme of a grain where appears the grain interiors lifetimes (τ GI ), the grain boundary recombination velocity (S GB ). Effective lifetimes estimation from model described by Equation 3.8 for grain boudary recombination velocity a) S GB = 10 4 cm.s -1 b) S GB = 10 5 cm.s -1

Figure 3 .

 3 Figure 3.28: a) Zoom of a RT panchromatic CL image in the targeted grain acquired under steadystate excitation. The SEM spot has been plotted as the intersection of the two red lines. b) Normalized CL spectra determined by averaging CL maps at low temperature under pulsed excitation (pink) and continuous excitation (black). Spectral ranges which will be used for the spectral integration of the streak images is presented as green: [1.425 eV; 1.49 eV] and red: [1.57 eV; 1.61 eV] colored rectangles.

Figure 3 . 29 :

 329 Figure 3.29: TRCL streak image acquired at LT in a CdTE grain with a spectrometer central wavelength of a) λ=750 nm and b) λ=850 nm. c) Decays determined from streak images a) and b), Ev1= [1.425 eV; 1.49 eV] corresponding to exciton and Ev2= [1.57 eV; 1.61 eV] corresponding to defects.

  .28 b): Defect transition [1.425 eV; 1.49 eV] (Green) and exciton related transition [1.57 eV; 1.61 eV] (Red).

Figure 4 . 1 :

 41 Figure 4.1: In a)-c) schemes of the NW growth using the VLS technique on a patterned Si substrate. a) Ga droplet catalyst deposition followed by VLS growth b). At the end of the growth c) a crystallization of the catalyst is achieved and then a shell is grown. d-f) present SEM images of the corresponding steps of the NW fabrication.

Figure 4 . 2 :

 42 Figure 4.2: Schemes of the different nanowire p-n junction architectures. a) The vertical junction corresponds to the vertical superposition of p-and n-type materials. b) The radial junction corresponds to a core-shell structure of p-and n-type materials.

Figure 4 . 3 :

 43 Figure 4.3: a) Scheme of the NW core characterized by a bulk lifetime (τ b ).b) The radial p-n junction is characterized by both n-(orange) and p-(green) type doped materials. c) A passivation shell can be grown in order to reduce the surface recombination velocity (S) and so to increase the carriers lifetime within the NW.

Figure 4 . 4 :

 44 Figure 4.4: a) SEM image of a GaAs Nanowire spread on Si (n-doped) substrate. b) RT CL panchromatic map. c) CL spectra extracted and average on 3×3 pixels from CL map presented in b) and marks are plotted on a). These measurements have been made at the C2N.

Figure 4 . 6 :

 46 Figure 4.6: a)Live SEM image of a GaAs NW. The 3 selected areas for CL spectra, plotted in c), are shown as Blue (WZ), Red (Inter) and Green (ZB) squares. b) Low temperature CL panchromatic map at 10K. c) Low temperature spectra determined by averaging the 3×3 pixels squares described in b). Peak energy has been determined for each peak using a one-Gaussian fit of the peak. Operating conditions were 6 kV and I=2 nA. (76715)

Figure 4 . 7 :

 47 Figure 4.7: Low temperature (10K) CL spectra determined by averaging 3×3 pixels from CL maps for a) WZ area and b) ZB area. CL maps were performed under different incident currents estimated afterwards with a Faraday cup.

Figure 4 . 8 :

 48 Figure 4.8: Schematic of the crystallographic structures of GaAS ZB and the WZ stackings. The ZB phase presents a ABCABC stacking and a ABAB in the case of the WZ. Figure is reprinted from [90].

Figure 4 .

 4 Figure 4.9: a) Band diagrams of the GaAs ZB and WZ phases. b) dipole allowed transitions in the WZ phase. Figure reprinted from [90].

Figure 4 . 10 :

 410 Figure 4.10: Scheme of the optical setup for CL polarimetry. This scheme corresponds to a zoom of the full C2N CL tool scheme (Figure 2.2). A linear polarizer on the optical path of the collimated light. The Linear polarizer orientation was changed manually. The 0°orientation of the polarizer has been defined to be parallel to the NW orientation.

Figure 4 . 11 :

 411 Figure 4.11: Evolution of the average LT (10 K) CL intensity, determined from LT CL maps of a GaAs thin film, as a function of the polarizer orientation. LT CL intensities have been normalized by the average intensity CL map measured without polarizer.

Figure 4 .

 4 Figure 4.12: a) SEM image of a GaAs NW. Figure b) and c) present panchromatic CL maps at low temperature (10K) of the same NW with different polarizer orientations: b) 0°and c) 90°. In Figure d) is plotted the evolution of the integrated CL intensity of WZ (blue-squares) and ZB (red-squares) as a function of the polarizer orientation. Evolution of the CL intensity is fitted with Equation 4.1 and the degree of polarization P is determined with Equation 4.2.

Figure 4 . 13 :

 413 Figure 4.13: Polar plot of the CL intensity at low temperature (10K) extracted from the pure ZB (red dots) and the pure WZ (blue dots) region of the NW as a function of the polarizer angle. Solide lines represent the evolution of the intensity of the NW as a function of the polarizer determined from Equation 4.1 for P(WZ)=-74 % (blue solid line) and for P(ZB)=5 %

Figure 4 . 14 :

 414 Figure 4.14: Low temperature panchromatic CL maps, of the same NW, under pulsed injection with different polarizer orientations a) 0°and b) 90°. Intensities have been normalized under different grayscales. WZ and ZB segments of the NW have been identified using the evolution of the intensity with the polarizer orientation and plotted as blue (WZ) and red (ZB) ellipses.

Figure 4 .

 4 15 shows the streak images acquired by focusing on the ZB segment (Figure 4.15a)) and WZ segment (Figure 4.15b)). In the streak image of the ZB (Figure 4.15 a)) we observe one decay centered around 1.495 eV. For the WZ (Figure 4.15 b)) two decays are observed: a short decay at 1.51eV and a long decay at 1.495 eV.

  .14).

Figure 4 . 15 :

 415 Figure 4.15: TRCL experiments at low temperature (10K) on the a) ZB and b) the WZ parts. White and red dashed indicate the spectral band which has been used for decay determination. c) low temperature CL spectra, recorded on the CCD camera, determined by averaging 3×3 pixels from WZ and ZB locations of CL maps under continuous and pulsed excitation. Polarizer was oriented at 90°.

Figure 4 . 16 :

 416 Figure 4.16: CL intensity decays determined from streak images presented in Figure 4.15a) and b). Spectral intervals which has been used for the decay determination is: ZB and WZ exciton [1.53 eV;1.50 eV], ZB and WZ defects [1.50 eV;1.48 eV].

Figure 4 .

 4 Figure 4.17: a) Room Temperature CL spectra ( dots) n-GaAs:Si thin films. The normalization factors are indicated. Generalized Planck law fits (Equation 4.3) have been plotted (colored lines), open circles and squares superimposed on the CL spectra mark E g and E fc determined from the fits. The CL normalization factors are indicated. b) Schematic band diagram in the case of a degenerate n-type GaAs. Energy levels of the Valence band (E v ), the Conduction band (E c ) and the quasi-Fermi levels for electrons (E fc ) and holes (E fv ) are shown. Figures are adapted from[34].

Figure 4 . 18 :

 418 Figure 4.18: Evolution of electron the Fermi level as a function of the electron concentration n. Blue squares represent the Fermi levels (at RT) estimated from the full fit of CL spectra with a parabolic absorption convoluted with an Urbach's tail. Theoretical evolution of the Fermi level as a function of the electron concentration from a parabolic absorption model (green dashed line) and from a corrected parabolic absorption model (blue solid line). Figure reprinted from[34].

Figure 4 .

 4 Figure 4.19: a) Low Temperature (20 K) CL spectra of n-GaAs:Si thin films. Spectra are normalized and shifted vertically for clarity. Figure is reprinted from [34].

Figure 4 . 20 :

 420 Figure 4.20: Evolution of the FWHM of the CL peak as a function of the electron concentration, measured by Hall effect measurements for GaAs:Si Thin films. The fit with a ∼n 2/3 law has been made (solid blue line) and compared with previously results of the literature [157] (black dashed line). Figure is reprinted from [33].

Figure 4 .

 4 Figure 4.21: a) RT CL spectra for n-type GaAs Thin films (Black dots). Spectra have been normalized and vertically shifted for clarity. Fits using Equation 4.3 are plotted in continuous colored line. Circles and squares mark the fitted E g and E fc , respectively.b) Evolution of the Fermi level as a function of the electron concentration for a parabolic model of absorption (dashed line), and with a non-parabolicity correction (blue solid). Circles correspond to the E fc estimated from the fits (summarized in Table 4.2) together with the electron concentration determined with the hall effect (summarized in Table 4.1).

Figure 4 .

 4 Figure 4.22: a) LT average CL spectra for n-type (Te) doped GaAs thin films. Spectra have been normalized and shifted vertically for clarity. b) FWHM evolution as a function of the electron concentration for two empirical models Chen et al.[START_REF] Chen | Determination of n-Type Doping Level in Single GaAs Nanowires by Cathodoluminescence[END_REF] 34] (red solid line) and De Sheng[START_REF] De-Sheng | Electrical properties and photoluminescence of Te-doped GaAs grown by molecular beam epitaxy[END_REF](blue solid line). FWHM experimental data, determined from the one-Gaussian fitting of the CL peak, are also plotted (black dots).

Figure 4 .

 4 Figure 4.23: a)Live-SEM recorded during the LT CL mapping of the Te NW1. Results of the image processing are presented through corrected SEM image b) and CL intensity c). d) Normalized LT average CL spectra have been determined by averaging over areas described by the white squares presented in the CL intensity map c).(White scale bar: 0.5 µm)

Figure 4 .

 4 Figure 4.24: a) Live-SEM recorded during the RT CL mapping of the Te NW1. Results of the image processing are presented through corrected SEM image b) and CL intensity c). d) Normalized RT CL spectra which have been determined by averaging over 4 × 4 pixels areas (white squares). (White scale bar: 0.5 µm)

Figure 4 .

 4 Figure 4.25: a)-d) RT CL spectra (black dots) integrated over the squares presented on the CL peak energy maps e)-g) and their fits (red solid line). RT CL spectra have been plot in linear-scale in h)-l). A gradient in the electron concentration along the NW is evidenced from the evolution of the fitted electron quasi-Fermi levels (E f c ).

Figure 4 . 26 :

 426 Figure 4.26: CL peak energy map where the nine rectangle areas from which the CL spectra have been integrated. Electron quasi-Fermi levels and electron concentration determined from the CL measurements are also summarized.

Figure 4 . 27 :

 427 Figure 4.27: Plot of the effective lifetime (τ eff ) as a function of the NW diameter for two surface recombination velocity S= 2.5×10 3 cm.s -1 (red line) and S=1×10 5 cm.s -1 (blue line). The evolution of the effective lifetime has been calculated from Equation4.6. We considered a τ bulk =20 ns.

  ) and i), and explained by the strain.

Figure 4 . 28 :

 428 Figure 4.28: CL maps of undoped GaAs NW with an a)b)c)d)e) InGaP shell (76851) and f)g)h)i)j) AlGaAs shell (76715) measured at low temperature (10 K). a) f) SEM image. b) g) CL intensity maps. c) h) Maps of the CL peak position energy. d)i) Maps of the CL FWHM. Both FWHM and CL peak position have been estimated using a 1-Gaussian function. j)e) Average CL spectra extracted from regions of theNWs displayed as white squares on CL maps.(white scale bars: 0.5 µm)

Figure 4 .

 4 Figure 4.29: a)b) streak images recorded with TRCL at 10K for a) the InGaP and the b) AlGaAs structures on the brightest part of the NW (described in Figure 4.28.) Low-temperature (10K) TRCL decays measured with the streak camera for GaAs-AlGaAs (blue line) and GaAs-InGaP (green line) coreshell nanowires. The decays have been estimated in NWs presented in Figure 4.28. The decays were determined by averaging streak images between: [810 nm;840 nm] for the AlGaAs and [830nm;850nm] for the InGaP passivated NW. Spectral ranges were informed on streak images with white lines.

Figure 4 . 30 :

 430 Figure 4.30: SEM images of a) GaAs/InGaP and b) GaAs/AlGaAs core-shell nanowires. Roomtemperature panchromatic CL maps for a) GaAs/InGaP and b) GaAs/AlGaAs. CL intensity scale has been set for each panchromatic CL map in order to optimize the contrast. The NW contour has been determined from the SEM contrast and plotted in the images.

Figure 4 . 31 :

 431 Figure 4.31: Room-temperature CL spectra recorded under pulsed conditions when the electron pulses are focused at a location of a NW, for three different GaAs/InGaP core-shell nanowires.

Figure 4 . 32 :

 432 Figure 4.32: Room-temperature TRCL streak images of three randomly picked GaAs/InGaP core-shell nanowires. The three NWs correspond to the ones for which we presented the CL spectra in Fiugre 4.31. streak image scale has been set for each image in order to optimize the contrast.

Figure 4 . 33 :

 433 Figure 4.33: Time decays measured with TRCL at RT on 3 single InGaP/GaAs NWs. Monoexponential fit and resulting effective lifetimes has been plotted (solid lines).

Figure 4 .

 4 34, τ ef f as a function of S and τ bulk for a d = 200 nm NW using Equation 4.6. In this Figure we can observe a transient behavior between vertical and horizontal lines. The vertical lines correspond to a regime where τ ef f is governed by the recombination in the bulk of the NW (low τ bulk values). On the other hand, an horizontal line corresponds to the

Figure 4 . 34 :

 434 Figure 4.34: Surface plot of the evolution of τ ef f as a function of S and τ bulk . Vertical (τ bulk = 0.35 ns) and horizontal (S = 10 4 cm.s -1 ) black dashed lines have been plotted corresponding to asymptotic values for τ ef f = 0.31 ns and d = 200 nm.

Figure 4 .

 4 Figure 4.35: a)b) Dark field TEM images of two representative InGaP-GaAs NWs heads. EDX maps have been recorded along the NW axis and are represented as green squares. Composition has been estimated for each EDX spectrum for the two NWs and the lattice mismatch has been calculated, and plotted in c), using Equation 4.11.

Figure 4 . 36 :

 436 Figure 4.36: RT CL maps for three GaAs-InGaP NWs. It has been recorded on the same set of NW as the one considered since the beginning: NW1 a)d)g), NW2 b)e)h) NW3c)f)i). a)b)c) live SEM images recorded at the same time than the CL maps. d)e)f) panchromatic CL maps. g)h)i) CL energy peak estimated by a 1-Gaussian picking function. White scale bars: 5 µm.

Figure 4 .

 4 Figure 4.38: a) SEM image of GaAs NW with an InGaP shell (76851). b) RT CL peak position map. c) RT CL spectra extracted from the bottom to the top of the NW (red solid lines). Each spectra has been normalized and shifted vertically for clarity. Spectra have been fitted using a two-Gaussian function (blue dashed lines) where the first Gaussian function is defined by (E 1 = 1.395 eV and σ 1 = 17 meV ), the second Gaussian function is defined by (E 2 = 1.435 eV and σ 2 = 22 meV ). From the results of the fit, the evolution of the low energy peak has been normalized by the maximum CL intensity and plotted in d). (white scale bar: 0.5 µm)

Figure 4 .

 4 Figure 4.39: a)-e) Cartographies panchromatiques haute résolution (128 × 128 pixels) de CL à température ambiante mesurées sur des couches minces de CdTe a) sans recuit et avec recuit sous atmosphère CdCl 2 à b) 400 °C c) 420 °C d) 440 °C et e) 460 °C.(Barre d'échelle: 5µm). Des histogrammes des intensités de CL à température ambiante sont présentés en f)-j), ils ont été déterminés à partir des intensités de CL des cartographies panchromatiques: f) sans recuit g) 400 °C h) 420 °C i) 440 °C j) 460 °C. Deux populations de points ont été mises en évidence : joints de grains (bleu) et intérieurs de grains (vert). k) Évolution de la longueur caractéristique en fonction de la température de recuit. l) Évolution des intensités intégrées de CL à basse température des pics de défauts (rouge) et excitoniques (vert). Ces intensités ont été estimées à partir de spectres moyens déterminés sur des cartographies (128×128 pixels)hyperspectrales de CL à basse température.

Figure 4 .

 4 Figure 4.40: a) Profil d'évolution de l'intensité radiative de CL et de b) la position du pic de CL en énergie en fonction de la profondeur pour des échantillons de CdSe x Te 1-x avec différentes compositions de Se (x Se =0, 0.1, 0.2, 0.3, 0.4). c) Evolution de l'intensité de CL en fonction de la position du pic estimée pour les échantillons de CdSe x Te 1-x . d) Evolution de la position du pic de CL du spectre moyen, estimé dans la partie du film riche en Se pour les 4 échantillons, en fonction de la quantité de Se initialement introduite dans les films.

Figure 4 .

 4 Figure 4.41: a) Image Live-SEM d'un NF de GaAs et b) l'image de CL panchromatic de celui-ci acquise à basse température.c) Spectres de CL déterminés en moyennant sur les zones représentées par des carrés sur la cartographie b). d) Evolution de l'intensité d'émission de la phase WZ (bleu) et ZB (rouge) en fonction de l'orientation du polariseur. e) Décroissance temporelle mesurée sur les phases WZ et ZB d'un même NF à basse température (10K).

L

  'utilisation de la technique de croissance de NFs par VLS conduit, dans certaines conditions de croissance, à la coexistence de phases Zinc-blende (ZB) et Wurtzite (WZ). La méconnaissance des propriétés optiques et de transport de la phase WZ rend difficile la prédiction de l'impact d'une alternance non-contrôlée des phases ZB et WZ au sein d'un même NF. Dans le but d'étudier ces deux phases, des NFs en GaAs avec des segments de phases pures WZ et ZB ont été synthétisés au C2N. Nous avons réalisé par la suite des cartographies de CL à basse température de ces NFs (Figure4.41 a)b)). A partir de ces cartographies, couplées avec des mesures résolues en injection, on a mis en évidence des gap en énergie similaires (E g (10K) = 1.513 eV ) pour les phases WZ et ZB (Figure4.41 c)). De plus, la polarisation de l'émission de ces deux segments a été mesurée par polarimétrie de CL. On a mis en évidence une très forte polarisation de l'émission de la WZ, perpendiculaire à l'axe du NF (i.e. axe c du réseau hexagonal de la WZ), et nous avons mesuré un degré de polarisation de -74% (Figure4.41 d)). Ce résultat a été ensuite utilisé comme un outil pour distinguer sans ambiguïté les deux segments du NF in-situ.Ce résultat a permis la mesure des temps de vie à différents endroits d'un même NF à basse température (Figure4.41 e)).

Figure 4 .

 4 Figure 4.42: a) Spectres de CL mesurés à température ambiante (TA) et b) basse température sur des films fins de GaAs:Te. Les spectres ont été normalisés et décalés par soucis de clarté. Les fits du spectre de CL à TA basés sur la loi de Planck généralisée ont été représentés en rouge.Les cercles et carrés ajoutés sur les courbes marquent les paramètres E g et E f c obtenus par modélisation, respectivement. c) Évolution du quasi-niveau de Fermi en fonction de la concentration en électrons. Les courbes correspondent aux évolutions théoriques dans le cas du modèle d'absorption parabolique (ligne pointillée) et non-parabolique (ligne continue). Les points correspondent aux quasi-niveaux de Fermi mesurés sur les échantillons. d)Évolution de la largeur à mi-hauteur (FWHM) du pic de CL à basse température (10K) en fonction de la concentration. Des évolutions théoriques empiriques ont été reportées sur ce graphique. Les points correspondent aux mesures sur nos échantillons.

Figure 4 .

 4 Figure 4.43: a) Evolution du temps de vie effectif (τ ef f ) en fonction du diamètre d'un NF de GaAs (τ bulk = 20 ns) pour deux vitesses de recombinaison de surface S = 10 5 cm.s -1 (bleu) et S = 2.5 × 10 3 cm.s -1 (rouge). b)d) Images live-SEM acquises et les cartographies de CL panchromatiques d)e) correspondantes mesurées sur des NFs de structure b)c) GaAs-AlGaAs et d)e) GaAs-InGaP. f) Décroissances temporelles pour trois NFs GaAs-InGaP mesurées par TRCL à température ambiante. Une fenêtre montrant une image TEM d'un NF du même échantillon a été ajoutée. f) Graphique en 3D de l'évolution du temps de vie effectif (τ ef f ) pour un NF de 200nm de diamètre. Une asymptote pour la valeur de vitesse de recombinaison de surface a été représentée (ligne blanche pointillée).

Figure 4 .

 4 43 b)c)) et l'autre avec une couche d'InGaP (Figure 4.43 d)e)). La comparaison quantitative de la luminescence émise par les deux structures à température ambiante et à basse température a montré que la structure avec une coquille en InGaP possédait une plus grande efficacité radiative à ces deux températures (Figure 4.43 c)e))

  Highest efficiency cells have just reached 23.35i % with a MaPbI 3 absorber. Despite good efficiencies reached by this technology, its low stability remains as a huge drawback.

TiO 2 is first deposited. Then classically perovskite is sandwiched between a mesoporous TiO 2 (n-type selective layer) and a Spiro-OMeTAD (p-type selective layer). Then gold or silver back contacts are deposited.

Chapter 2 Cathodoluminescence charac- terization setup Chapter content 2.1 The CL setup: main features . . . . . . . . . . . . . . . . . . . . . . . 2

  .1.1 Overview of the cathodoluminescence setup . . . . . . . . . . . . . . . . 2.1.2 The electron source . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.3 Light collection system . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1.4 Nanopositioning system and control of the temperature . . . . . . . . . 2.1.5 Dispersion and detection system . . . . . . . . . . . . . . . . . . . . . . 2.1.6 CL mapping and hyperspectral cube . . . . . . . . . . . . . . . . . . . .

2.2 Time-resolved cathodoluminescence . . . . . . . . . . . . . . . . . . . 2.3 Electron-matter interactions . . . . . . . . . . . . . . . . . . . . . . . . 2

  .3.1 Interaction volume: CASINO simulations . . . . . . . . . . . . . . . . . 2.3.2 Estimation of the injection level . . . . . . . . . . . . . . . . . . . . . .

2.4 CL emission and calibration . . . . . . . . . . . . . . . . . . . . . . . .

  Calibration of the CL setup . . . . . . . . . . . . . . . . . . . .

	2.4.1 Luminescence of semiconductors . . . . . . . . . . . . . . . . . . . . . .
	2.4.1.1 The generalized Planck law . . . . . . . . . . . . . . . . . . . .
	2.4.1.2 Absorption model . . . . . . . . . . . . . . . . . . . . . . . . .
	2.4.2 Calibrated CL Setup: protocol and preliminary results . . . . . . . . . .
	2.4.2.1 Calibration of the reference diode irradiance . . . . . . . . . .
	2.4.2.2

2.5 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Table 2 . 1 :

 21 Estimation of the generation rate of electron-hole pairs in GaAs and CdTe for standard operating conditions: 6keV and I = 2nA. Values used in order to estimate for the generation rate have also been summarized.

Table 2 .

 2

2: Estimation of the carriers density per second in GaAs and CdTe for standard operating conditions: 6 keV and I=2 nA. Carriers density generated from CL have been compared to operating conditions with the 1-sun injection, which has been determined from Equation

2

.8.

  .3.

		Photon flux
		(photons.s -1 )
	Center	2.37×10 12
	Edges	1.45×10 13

Table 2 . 3

 23

: Photon flux of the GaAs diode (30µm diameter) determined with the calibrated system. Flux from the sides and from the center have been separated.

Table 2 . 4 :

 24 Quantum efficiency of the dispersion and detection system at 870nm and theoretical collection efficiency of the C2N CL Setup. Results of the calibration Ratio, considering λ = 870 nm, from theoretical consideration and the different experimental methods (diode and LASER) are summarized.

	syst.resp.(870nm)	0.31
	f coll. eff.	0.4
	I diode	1.16 10 17 cts.s -1
	I diode+edges	1.25 10 18 cts.s -1
	Ratio diode	2.04×10 5 photons.cts -1
	Ratio diode+edges	1.3 ×10 5 photons.cts -1
	Ratio LASER	1.5×10 5 photons.cts -1
	Ratio theoretical	3.6×10 4 photons.cts -1

CdTe thin-films Chapter content 3.1 Background and state of the art of CdTe

  State of the art for CdTe polycrystalline thin films . . . . . . . . . . . . 3.1.2 V oc deficit in polycrystalline CdTe . . . . . . . . . . . . . . . . . . . . . 3.1.3 Recombinations at interfaces in polycrystalline materials . . . . . . . . .

• Luminescence of semiconductors is characterized by the generalized Planck law (See Subsection 2.4.1). Fitting the luminescence spectrum collected by CL allows to determine fundamental materials properties (bandgap, Urbach's tail, radiative efficiency...). An absolute calibration is needed in order to further investigate the light emitted by the SC. Different protocols for the absolute calibration of the C2N CL setup were described in Subsection 2.4.2. Results were presented, showing consistency of the calibration factors from the different protocols. Polycrystalline . . . . . . . . . . . . . . . . 3.1.1 3.1.3.1 Role of interfaces in polycrystalline materials . . . . . . . . . . 3.1.3.2 Post deposition CdCl 2 annealing . . . . . . . . . . . . . . . . . 3.1.4 CL on polycrystalline CdTe thin films . . . . . . . . . . . . . . . . . . . 3.

2 Impact of the CdCl 2 annealing temperature on the defects passi- vation

and formation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1 Description of the series of CdTe thin films . . . . . . . . . . . . . . . .

3.2.2 Analysis of RT CL maps . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.2.1 Evolution of the radiative efficiency with the CdCl 2 annealing temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.2.2 Estimation of the average grain size . . . . . . . . . . . . . . . 3.2.2.3 Statistical analysis of the luminescence intensity . . . . . . . . 3.2.2.4 Local analysis: characteristic length . . . . . . . . . . . . . . . 3.2.3 Defects nature and distribution in grains . . . . . . . . . . . . . . . . . . 3.2.3.1 Defects in as-deposited CdTe thin-film . . . . . . . . . . . . . . 3.2.3.2 Impact of the CdCl 2 annealing temperature on defects nature and distribution . . . . . . . . . . . . . . . . . . . . . . . . . . 3.

3 Investigation of CdSe x Te 1-x alloys properties

  Quantitative estimation of τ bulk and S GB by effective lifetimes measurements in CdTe grains . . . . . . . . . . . . . . . . . . . . . . . . . . . .

	3.4.1

. . . . . . . . . . . . . . 3.3.1 Notable results on CdSeTe/CdTe layers . . . . . . . . . . . . . . . . . . 3.3.2 Impact of the Se composition on the bandgap and the defects passivation within CdSeTe/CdTe layers . . . . . . . . . . . . . . . . . . . . . . . . . 3.3.2.1 Description of the series of CdSe x Te 1-x beveled samples . . . . 3.3.2.2 In depth analysis of CdSeTe thin films . . . . . . . . . . . . . . 3.3.2.3 Defects passivation with Se in CdSeTe layers . . . . . . . . . . 3.4 Lifetime assessment in CdTe grains . . . . . . . . . . . . . . . . . . . 3.4.2 Proof of feasibility: low temperature lifetime measurement in a CdTe grain . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

Table 3 . 1 :

 31 CdTe samples. The post deposition CdCl 2 annealing temperature is adressed Five CdTe samples were grown by CSS with CdCl 2 and annealed under different CdCl 2

Table 3 . 2 :

 32 An average grain area is determined for each sample, and expressed using the grain diameter of a disc of same area (seeFigure 3.11). Standard deviation are also specified in brackets.

Table 3 .3: Fitting

 3

parameters of each family of pixels (grain boundary and grain interior) using two Gaussian functions, see Figure

3

.9. The values are normalized to the mean value of I GB obtained for the sample annealed at 400°C.

Table 3 .4: An

 3

average grain area is determined for each sample, and expressed using the grain diameter of a disc of same area (see

Figure 3.11)

. The average characteristic length is determined by fitting the luminescence decay close to grain boundaries with Equation 3.2, see

Figure 3.15

Table 3 . 5

 35

	. The 20%

: CL Peak positions, FWHM and the maximum CL intensity values determined from CL profiles for the 5 CdSeTe samples.

Table 4 . 1 :

 41 19 cm -3 . Carrier concentration and mobility were first estimated with Hall effect measurements (see Subsection 1.4) and summarized in Table4.1. Electron concentrations ranging from 4.9 × 10 17 cm -3 to 1.3 × 10 19 cm -3 have been measured. Description of n-GaAs:Te thin films grown by MBE at the University of Tampere. The doping level was assessed using Hall effect measurements. Mobilities which have been measured have also been provided.

	Sample name Carrier concentration	Mobility
		(cm -3 )	(cm 2 .V -1 .s -1 )
	D1	1.3×10 19	1260
	C3	6.8×10 18	1660
	C1	4.2×10 18	1545
	B1	9.8×10 17	2629
	A3	7.5×10 17	2960
	A1	4.9×10 17	2970

Table 4 . 2 :

 42 Peak position and optimal fitting parameters (E g , E fc , Urbach's) tail γ ) for RT CL spectra measured on planar n-GaAs (Te dopant). Electron concentration determined from Equation4.4 are also indicated.

		Carrier density Peak	Eg	E fc -E CB	γ	Exp. electron conc. (RT)
		cm -3	(eV) (eV)	(eV)	(eV)	cm -3
	D1	1.3×10 19	1.554 1.394	0.219	0.066	9.1×10 18
	C3	6.8×10 18	1.473 1.376	0.139	0.021	4.5×10 18
	C1	4.2×10 18	1.440 1.407	0.024	0.020	(6.4×10 17 *)
	B1	9.8×10 17	1.430 1.409	0.00	0.013	<1×10 18
	A3	7.5×10 17	1.433 1.409	0.006	0.011	<1×10 18
	A1	4.9×10 17	1.428 1.408	0.00	0.008	<1×10 18

  Carrier density Peak FWHM Exp. electron conc. (LT)

		cm -3	(eV)	(meV)	cm -3
	D1	1.3×10 19	1.635	0.185	1.1×10 19
	C3	6.8×10 18	1.572	0.115	5.4×10 18
	C1	4.2×10 18	1.534	0.049	1.5×10 18
	B1	9.8×10 17	1.527	0.038	1.0×1.0 18
	A3	7.5×10 17	1.52	0.031	7.5×10 17

Table 4 . 3 :

 43 CL Peak positions and FWHM values determined from LT CL spectra (see Figure 4.22 a)).Experimental electron concentration calculated with Equation4.5 are also reported.

Table 4 . 4 :

 44 GaAs:Te NW grown by VLS in a MBE. The substrate temperature (T substrat ), the As specie (As 2 or As 4 ) and the shell presence have been summarized. NWs lengths and diameters have been measured using SEM images

1.4. Materials characterization and aim of the thesis crystalline thin-films solar cells and GaAs nanowire-based solar cells.

CL maps have been measured at low temperature over the very same area as previous RT maps and are compared in Figure 3.17. This allows a direct comparison between the radiative properties and characteristic lengths investigated at room temperature with the defects revealed by low-temperature measurements. The LT CL spectra averaged over the whole maps are shown in Figure 3.17a) for the different annealing temperatures, T. Contrary to the as-deposited thin film, they exhibit a single defect band in the 1.35-1.5eV spectral range. The disappearance of the Z-band can be attributed to the passivation effect of the V Cd by the Cl introduced during the CdCl 2 treatment. A shoulder at 1.475eV, overlapped with the A-band, attributed to Y-band corresponding to exciton bounded to a dislocation [START_REF] Halliday | A study of the depth dependence of photoluminescence from thin film CdS/CdTe solar cells using bevel etched samples[END_REF][START_REF] Kraft | Comprehensive photoluminescence study of chlorine activated polycrystalline cadmium telluride layers[END_REF][START_REF] Okamoto | Characterization of Highly Efficient CdTe Thin Film Solar Cells by Low-Temperature Photoluminescence[END_REF], appears for annealing temperatures over 420°C.

The intensity of each emission band varies with the annealing temperature. The luminescence a strain due to the lattice mismatch is expected.

In Figure 4.35 we plotted RT CL maps of the GaAs-InGaP core-shell NWs. The CL peak positions have been picked for each spectrum in order to follow the evolution of the peak position along the NW. Very inhomogeneous CL peak positions can be observed on the different maps for the three NWs. Besides, we can notice the same trend along the NW, for each NW. Indeed, a shift of about 20∼30 meV can be observed from one end of the NW to the other. The AlGaAs shell sample has been chosen as a reference as no strain on the core is expected. In Figure 4.37 is plotted the average spectra recorded on the GaAs-AlGaAs NW and on the three investigated GaAs-InGaP NWs. On the AlGaAs shell spectrum, we observe only one CL peak at 1.42 eV which corresponds to the band to band transition (See Section 4.2). However, we can observe that the CL spectra seems to be composed of two distinct peak for the InGaP shell NWs. A second lower energy peak at 1.395 eV is observed. Signorello et al. [START_REF] Signorello | Tuning the Light Emission from GaAs Nanowires over 290 meV with Uniaxial Strain[END_REF] has quantified the shift of the luminescence peak induced by a tensile strain of 85 meV/%. In our case, if we consider a 30 meV shift of the CL peak, a strain of 0.4 % is found.

A red-shift of the peak has been previously observed in the GaAs-InGaP core-shell NWs compared the GaAs-AlGaAs structure. A lattice mismatch between the core and the shell has been measured with TEM/EDX maps to be of ≈1 %. Using the shift of the CL peak position we estimated a strain of about 0.4 %.

Estimation of a diffusion length in a nanowire

An inhomogeneous peak position has been observed previously. Indeed the shift in the CL peak position seems to depend on the location on the NW. In Figure 4.28 c), we can notice a red-shift near the top of the NW. The top of the NW is supposed to suffer from the highest strain as it is surrounded by the InGaP shell [START_REF] Glas | Why Does Wurtzite Form in Nanowires of III-V Zinc Blende Semiconductors?[END_REF].

In Figure 4.38, we provide a more in-depth analysis of RT CL maps on GaAs-InGaP NWs. We can observe in Figure 4.38 b) that the map presents two distinct area with an intermediate region. In Figure 4.38 c) a linescan of the CL spectra at RT has been plotted. On these spectra

Summary

In this Chapter we focused on the characterization of GaAs NWs grown by MBE using the VLS method with a prepatterned substrate. Because of their low dimension, characterization of the optical properties of single NWs is challenging but necessary for efficiency improvements of NW-based solar cells:

• In section 4.2, the impact of the crystalline phase on optical and transport properties has been studied. Indeed, NWs GaAs can grow under ZB or WZ depending the growth conditions. NWs with distinct ZB and WZ segments were prepared and analyzed using CL maps at LT (10K) ( • In section 4.3, we performed contactless quantitative characterization of carrier densities with a nanometer scale. First, we detailed the method developed by Chen et al. [34] at C2N in order to determine quantitatively the electron concentration in Si-doped GaAs NWs.

Then we demonstrated that this method can also be used to determine quantitatively the electron concentration in the case of a Te doping. We applied this technique to demonstrate that we can achieve high doping (n=3.3 10 18 cm -3 ) in GaAs NW doped with Te.

Besides we showed that the Te incorporation results in a axial dopant gradients, with a spatial resolution of∼ 300 nm

• In section 4.4, surface passivation of the NWs has been investigated using high bandgap III-V shells. We first discussed theoretically the impact of surface passivation on the minority carriers lifetimes. A gain in voltage of 150 meV can be achieved by reducing the surface recombination velocity from 10 5 cm.s -1 to 2.5 10 3 cm.s -1 for a GaAs NW (d=200 nm), assuming τ bulk =20 ns. Then two core-shell structures were compared: GaAs-AlGaAs and GaAs-InGaP. Quantitative comparison of LT and RT CL hyperspectral maps show a higher radiative efficiency for the GaAs-InGaP structure compared to the GaAs-AlGaAs one. Lifetimes measurements have been performed at RT on three different GaAs-InGaP NWs. RT Average lifetimes of 0.31 ns for d=200 nm NWs were measured by TRCL. By combining these results with a simple theoretical model, an upper bound of 1.7 10 4 cm.s -1 for the SRV was estimated. Besides, the impact of the lattice mismatch strain on the core properties was examined. Because of ≈1.0% lattice mistmatch, a stress of 0.4% was determined in the GaAs core of the NW from the shift of the CL peak. From RT CL linescan along the NW axis, we observe that the strain is localized at the top of the NW. A diffusion length of L n = 370 nm was measured on the very same NW on which we measured a lifetime. From these results we determined a diffusion coefficient D=3.34 m 2 .s -1 and an electron mobility of µ=130 cm 2 .V.s -1 .

Articles and communications

Articles

Title : Characterization of nanomaterials by cathodoluminescence for photovoltaics applications

Keywords : cathodoluminescence, nanowires, CdTe, GaAs Abstract : The recent development of photovoltaic devices makes use of materials with properties that may vary at the sub-micrometer scale, for instance nanowire-based III-V solar cells based or CdTe polycrystalline thin films. This thesis aims at developing new methods and techniques based on cathodoluminescence (CL) and time-resolved CL (TRCL) in order to probe material properties at the nanometer scale. A first part is devoted to CdTe polycrystalline thin films. The role of CdCl 2 annealing on the formation and passivation of defects in the grain interior and at the grain boundaries is analyzed by combining high-resolution CL maps at roomand low-temperature. The incorporation of Se at different concentrations and its effect on defect passivation and radiative efficiency of CdSe x Te 1-x thin films is investigated and correlated to macroscopic measurements (V oc , lifetimes).

A second part is devoted to GaAs nanowires. The optical and transport properties of wurtzite (WZ) and zincblende (ZB) phases are compared in single nanowires, and radiative decay measurements are presented (𝜏(𝑊𝑍) = 0.55 𝑛𝑠 et 𝜏(𝑍𝐵) = 0.62 𝑛𝑠). N-doping obtained by incorporation of Te is determined quantitatively, and its homogeneity is studied in single nanowires by CL maps. We evidence doping levels up to n=3.3x10 18 cm -3 , and a doping gradient along nanowires. Finally, the passivation of GaAs nanowires by an AlGaAs or InGaP shell is characterized by CL measurement of the radiative intensity, diffusion length and lifetime of charge carriers. The techniques and methodologies presented in this thesis may be applied to a wide variety of semiconductor materials for photovoltaic or optoelectronic applications.