General Introduction

"Innovation", "high performance" and "fast data transfer" are part of the expressions that can describe the progress of the semiconductor industry through the years. Indeed, in semiconductor-based technology big efforts are made to reduce cost, achieve greater embedded functionality and higher interconnect density, enable lower power consumption and better reliability [START_REF] Mounier | Silicon photonics 2018[END_REF]. In order to answer to these requirements, improvements can be done in terms of production process steps, device architecture and the materials that will constitute the devices. These kinds of developments are realized to push up the performance of electronic and opto-electronic devices. As a result, new transistor concepts such as tunnel field effect transistors (TFETs) as an alternative to the conventional metal oxide semiconductor field effect transistors (MOSFETs) are prominent alternatives to boost transistor efficiency. Meanwhile, optical interconnects could help with the issues of heat dissipation of integrated circuits (ICs). Indeed, the application of photons, for the data and information transfer, requires less energy and provides higher speed of transfer [START_REF] Wirths | Group IV Epitaxy for Advanced Nano-and Optoelectronic Applications[END_REF]. An efficient electronic-photonic integrated circuit (EPICs) would then be needed to implement a single material that could be used for both electronic and photonic components. Germanium-Tin (GeSn), group-IV alloys, are interesting materials to achieve these goals.

The number of research reports on GeSn materials steadily increased since 1960, as observed in Fig. 0.1 and abruptly augmented after 2010 thanks to the multiple applications to which this material can be used.

GeSn can be used as source and drain stressor in Ge MOSFETs, and to obtain high mobility channels in either pMOSFETs and pTFETs. The mobility enhancements with the Sn content have been proven, which indicates the effectiveness of using GeSn alloys for high-performance MOSFETs [START_REF] Zaima | Growth and applications of GeSn-related group-IV semiconductor materials[END_REF]. On the other hand, the addition of a sufficient amount of Sn in the Ge lattice (about 10 at.%) offers the potential to get a direct band-gap structure. This material could be then the final piece to achieve a monolithically integrated group-IV laser that is compatible with the Si-CMOS technology (complementary metaloxide-semiconductor). Further technological achievements in GeSn-technology can then enhance electronic and opto-electronic devices functionalities.

No matter the application, low-resistive, stable and reliable metallic or intermetallic contacts are key components to inject electric current in the devices. The integrated contacts are expected to fulfill certain requirements low contact or sheet resistance, low thermal processing budgets, high thermal stability and be compatible with Si CMOStechnology.

Ni-based contacts have been proposed and widely studied as they allow the formation of the Ni(GeSn) intermetallic at relative low temperatures (around 350 ) with low values of contact (R c ) and sheet (R sh ) resistances (around 10-30 Ω/sq). Nevertheless, Ni-based contacts exhibit a poor thermal stability linked to the agglomeration of the Ni(GeSn) layer and surface segregation of Sn atoms.

This PhD thesis was dedicated to the systematic and comprehensive development and characterization of Ni / GeSn intermetallics. The objective was to propose efficient and reliable contacts for GeSn-based devices. The thesis is divided in five chapters, as schematized in Fig. 0.2.

The first chapter will briefly describe the GeSn material properties, the interest of including this alloy in electronic and opto-electronic devices as well as the challenges that have to be overcome in order to actually succeed the integration. On the other hand, certain notions related to metal / semiconductors contact elaboration and the interest of obtaining them through a solid-state reaction will be discussed. Additionally, certain parameters that, in a general view, can characterize a good contact will be mentioned.

The second chapter will describe the characterization techniques used to assess the different properties of the developed contacts. The objective is to highlight the reasons why these techniques are properly suited to provide critical information about the contacts properties (physicochemical, structural or electrical properties) either on the as-deposited state and during the evolution of the metal / semiconductor solid-state reaction.

The third chapter will be introduced by an overview about the state-of-the-art of Ni / GeSn intermetallics. Later, the systematic and comprehensive analysis of this type of intermetallics, carried by us will be addressed. The Ni / GeSn general properties were analyzed in terms of the phase sequence, the morphological and the electrical evolution. In a more specific manner, different phenomena occurring during the solid-state reaction will be also assessed like the crystalline nature of the Ni-rich phase and the impact of Sn In order to go further, chapter four will present different alternatives to enhance the Ni / GeSn thermal stability from the use of pre-amorphization by implantation, the addition of an alloying element (Pt, Co) to Ni, to the use of laser pulse as the annealing technique. The objective will be to evaluate the impact of performing modification on different steps of the contact elaboration process. Indeed, pre-amorphization by ion implantation is expected to modify the GeSn substrate, meanwhile Pt or Co addition modify the metallization layer and laser annealing will be considered as an alternative thermal treatment for the final step of the process. The impact of the different alternatives will be systematically analyzed in terms of the phase sequence, the morphological and the electrical evolution during the solid-state reaction.

Finally, the fifth chapter will present a brief overview of an alternative to Ni-based metallization which consist in the use of a Ti layer. Equally as for previous cases, the impact of Ti metallization will be systematically analyzed in terms of the phase sequence, the morphological and the electrical evolution during the solid-state reaction.

To conclude, a review of the main significant results to the GeSn contact technology will be given. In addition, a process scheme to elaborate good quality, stable and reliable contacts compatible to GeSn-technology will be proposed together with the perspectives to improve further works.

GeSn material

This first chapter addresses an overview of the great interest of introducing GeSn material in electronic and opto-electronic devices and the necessity to elaborate high quality contacts for these devices. The division of this chapter is schematized in Fig. 1.1. A description of GeSn properties, historical research and perspectives is presented in section 1.1.1. Then, a brief summary of the main growth challenges and envisioned applications is addressed in sections 1.1.2 and 1.1.3, respectively. As stated in the introduction, no matter the application, low-resistive, stable and reliable metallic or intermetallic contacts are key components to inject electric current in the devices. A description of the elaboration of metal / semiconductor intermetallics for contact technology and the approaches used for GeSn layers are also described in section 1.2.

GeSn properties and historical background

Germanium (Ge) and tin (Sn) are both group-IV materials which can be considered for multiple applications for the semiconductor industry. From a brief description of the Ge-Sn equilibrium phase diagram in Fig. 1.2 (a) and (b), it can be mentioned that (i) the eutectic point is located at a temperature and composition very close to those of pure Sn, (ii) Ge and Sn can form two stable phases, the liquid solution and the solid solution of Sn in Ge, (iii) the maximum solubility, of Sn in Ge, is about 1.1 at.% Sn at about 400 and (iv) the solubility at the eutectic temperature is less than 1 at.% Sn [4].

Above those values, the phase diagram predicts a two component mixture (Ge + β-Sn). Likewise, above the melting point, i.e. 231 , a two phase mixture (solid Ge + liquid β-Sn) can occur. Below this temperature, two solid phases are created (body centered tetragonal β-Sn phase + diamond lattice Ge) [START_REF] Aubin | Low temperature epitaxy of Si, Ge, and Sn based alloys[END_REF]. Like germanium, α-Sn crystallizes in the diamond phase, but its lattice parameter (a α-Sn = 6.489 Å) is 14% larger than that of Ge (a Ge = 5.658 Å). Additionally, this allotropic element presents a phase transition from its semiconducting so-called α-Sn (gray tin) phase with diamond cubic lattice structure to the metallic phase (β-Sn or white tin) with a body-centered tetragonal structure.

This transition occurs at 13.2 and goes along with a volume change of approximately 26% [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF]. Owing to these features, the synthesis of GeSn layers is rather challenging. A brief revision of the GeSn growth characteristics is presented in next section 1.1.2.

The general properties of the binary Ge-Sn system have been studied since 1939 [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF]. Meanwhile, the possibility to obtain a fundamental direct band-gap from this material was first mentioned in 1982, as displayed in Fig. 1.3 [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF][START_REF] Doherty | Progress on Germanium-Tin nanoscale alloys[END_REF][START_REF] Goodman | Direct-gap group IV semiconductors based on tin[END_REF]. This predictions motivated investigations to epitaxially stabilized GeSn alloys with Sn concentrations exceeding the equilibrium solid solubility. The growth of Ge 1-x Sn x thin films has been reported using many fabrication techniques including molecular beam epitaxy (MBE), ion-assisted MBE, solid phase epitaxy (SPE), sputtering, pulse laser deposition (PLD) or chemical vapor deposition (CVD) [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF][START_REF] Doherty | Progress on Germanium-Tin nanoscale alloys[END_REF], as observed in Fig. 1.3. MBE and CVD techniques have however played the leading role in terms of practicability, layer quality and applicability [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF]. These techniques have also been the preferential techniques used for GeSn-based contacts research in later years, as Chapter I. GeSn material from growth to applications displayed in Table 1.1. The kinds of samples used for this PhD thesis consisted in thin GeSn layers that were grown by reduced pressure-chemical vapor deposition (RP-CVD), as it will be briefly mentioned in section 1.1.2.

Reference

Substrate Ge 1-x Sn x growth Sn% [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Han | High-mobility germanium-tin (GeSn) P-channel MOSFETs featuring metallic source/drain and sub-370 process modules[END_REF][START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF] Ge (0 0 1) MBE at 180 5.3% [START_REF] Li | Electrical characteristics of Ni ohmic contact on n-type GeSn[END_REF] n-type Si (0 0 1) MBE 4.9% [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF] Ge/Si RP-CVD 4% [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF] Ge (0 0 1) UHV-MBE 2, 4, 5 and 8% [START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF] Ge/Si RP-CVD 1-3% [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF] Ge (0 0 1) MBE 8.6% [START_REF] Zhang | Formation and characterization of Ni/Al ohmic contact on n + -type GeSn[END_REF] Si (0 0 1) MBE 3.2% [START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF] Si (0 0 1) RP-CVD 12% [START_REF] Tong | Ni(Ge 1-x Sn x ) ohmic contact formation on n-type Ge 1-x Sn x using selenium or sulfur implant and segregation[END_REF] Ge (0 0 1) MBE 4.2% [START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF] N-type Ge (0 0 1) MBE 2, 3.6, and 6.5%

GeSn growth challenges

In order to promote improvements on GeSn-based devices, great efforts have been made to implement growing conditions far from equilibrium to exceed the solid solubility limit of α-Sn in Ge. The idea is to ensure an amount of incorporated Sn in which the indirectdirect band-gap transition can be attain. Another aspect to take care, is to prevent GeSn phase separation via surface segregation or bulk precipitation [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF].

Chemical vapor deposition (CVD) techniques, at low temperatures, have been highly suited for those ends when combining high growth rates and really low growth temperatures (the lower the growth temperature is, the higher the Sn at.% that can be incorporated in the Ge lattice) [START_REF] Aubin | Low temperature epitaxy of Si, Ge, and Sn based alloys[END_REF]. Aubin et al. investigated low temperature epitaxy of high Sn content GeSn alloys in a retro-fitted industrial Reduced Pressure-Chemical Vapor Deposition (RP-CVD) tool. The main results can be summarized as follows -A linear GeSn growth rate increase together with a sub-linear Sn concentration increase occurred as the SnCl 4 precursor mass-flow increased, irrespective of the Ge 2 H 6 precursor mass flow (fixed or varying) and temperature. The Sn atoms indeed seemed to catalyze hydrogen desorption, resulting in higher GeSn growth rates for high SnCl 4 mass-flows.

-The GeSn growth rate exponentially increased with the temperature. Meanwhile, the Sn content decreased linearly as the temperature increased. at 100 Torr on Ge strain relaxed buffers (SRBs) [START_REF] Aubin | Low temperature epitaxy of Si, Ge, and Sn based alloys[END_REF].

From these studies, GeSn layers with Sn contents as high as 15 at.% were obtained at 300 of high quality, as observed in Fig. 1.4. This resulted in high Sn content (> 10 at.%) GeSn alloys grown using low temperature processes. Those layers were thin, fully compressively strained on the Ge strain relaxed buffer (SRB) underneath.

However, being able to grow partially / fully relaxed GeSn layers with a high Sn content (> 10 at.%) is mandatory to fabricate efficient opto-electronic devices. Strain relaxation, for a given Sn content, can modify the directness of the GeSn band-gap. Fully / partially relaxed GeSn layers with relatively high Sn contents are also interesting to increase tensile strain Ge layers and improve electron mobility. This necessary strain relaxation, can be achieved on thick GeSn layers in which, however, the thermal stability is low. The studies made by Aubin et al. concluded with the implementation of GeSn step-graded structures with high crystalline quality and surface morphology stability, as observed in Fig. 1.5.

The advantage of using this type of step-graded structure was the achievement of a gradual strain relaxation, minimizing Sn segregation / precipitation. Those results were thus very interesting to grow a direct band-gap, thick GeSn layer with 16 at.% of Sn, while keeping superior structural and optical properties. Such layers were also processed to fabricate efficient GeSn-based optical components [START_REF] Aubin | Low temperature epitaxy of Si, Ge, and Sn based alloys[END_REF]. The growth of proper GeSn layers , that can be integrated for electronic or opto-electronic applications, was thus achieved under these process parameters. For these reasons, this technique was used to grow the GeSn semiconductor substrates used for this PhD thesis.

GeSn perspectives for electronic and opto-electronic devices

The improvements on GeSn growth allowed numerous achievements for device integration, as observed in Fig. 1.3. There is a great interest to grow high quality GeSn, either for electronic or opto-electronic applications. Indeed, for electronic applications, GeSn material can be used for (i) source and drain stressors in Ge MOSFETs thanks to the control of the energy band structures and the effective mass for GeSn layers; (ii) high mobility channel to increase MOSFETs performance, in which the mobility increases with Sn content, and (iii) innovative architectures like fin field-effect transistor (FinFETs) or Tunnel Field-Effect Transistors (TFETs) [START_REF] Zaima | Growth and applications of GeSn-related group-IV semiconductor materials[END_REF][START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Ge is an indirect bandgap semiconductors with positive band-gap 0.7 eV, while semiconducting α-Sn with a diamond lattice structure exhibits a negative and direct band-gap [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF].

On the other hand, opto-electronic applications can, greatly profit from the indirectdirect band-gap transition that can be achieved with Sn introduction on the Ge lattice. α-Sn band structure has been described considering the material as a semi-metal with an inverted band structure, that is, the conduction and valence bands are bent downwards and upwards, respectively. This material presents then a negative band-gap of about -0.41 eV. Meanwhile, Ge is an indirect group-IV semiconductor, see Fig. 1.6 [START_REF] Wirths | Si-Ge-Sn alloys: From growth to applications[END_REF].

Nevertheless, the small energy separation of 140 meV between indirect (L) and direct (Γ) conduction band valleys in Ge can be overcome by inducing tensile strain. This has motivated numerous research towards Ge band-gap engineering. Consequently, certain solutions have been proposed to introduce tensile strain in Ge thin films such as (i) micromechanical strain and epitaxial growth of Ge on larger lattice constant buffer layers, (ii) substitution of Sn into Ge lattice to form a semiconducting GeSn material which results in a similar effect to that observed in Ge under tensile strain, i.e. lowering the energy of the direct valley below the indirect one, as observed in Fig. 1.7 [START_REF] Gupta | Achieving direct band gap in germanium through integration of Sn alloying and external strain[END_REF]. The introduction of a sufficiently high fraction of Sn in the Ge lattice can decrease the energy of the Γ-valley below the one of the L-valley. Recent calculations indicated that this indirect-direct transition for relaxed GeSn can occur with Sn concentrations in the range of 6.5-11.0 at.% as displayed in Fig. 1.7. This direct band-gap group-IV material can thus represent a pathway towards the monolithic integration of Si-photonic circuitry and complementary metal-oxide-semiconductor (CMOS) technology.

The first demonstration of GeSn lasing effect was achieved in 2015, at low temperatures and relatively high optical pumping [START_REF] Wirths | Lasing in direct-bandgap GeSn alloy grown on Si[END_REF]. This breakthrough result demonstrated that a Chapter I. GeSn material from growth to applications direct band-gap group-IV laser on Si (0 0 1) can be achieved and represents a promising proof of principle for a CMOS-compatible gain material platform for cost-effective integration of electronic and photonic circuits. The use of this direct band-gap material can then be envisioned for the elaboration of high performance p-i-n photodetector [START_REF] Mathews | Extended performance GeSn/Si (1 0 0) p-i-n photodetectors for full spectral range telecommunication applications[END_REF][START_REF] Zhang | High-responsivity GeSn short-wave infrared p-i-n photodetectors[END_REF], heterojunction / homojunction LEDs, lasers [START_REF] Wirths | Lasing in direct-bandgap GeSn alloy grown on Si[END_REF], emitters / detectors for mid-infrared applications spectroscopy or sensors [START_REF] Reboud | Optically pumped GeSn micro-disks with 16% Sn lasing at 3.1 µm up to 180 K[END_REF][START_REF] Lin | Midinfrared integrated photonics on silicon: a perspective[END_REF][START_REF] Jérémie | GeSn lasers covering a wide wavelength range thanks to uniaxial tensile strain[END_REF].

In summary, GeSn group-IV material presents properties that can be useful for source and drain stressors in Ge MOSFETs, high mobility channels in either pMOS-FETs and pTFETs and to achieve a monolithically integrated group-IV laser that is compatible with Si-CMOS technology. However, the low Sn thermodynamic solubility in Ge (< 1 at.%) and the large lattice parameter differences (> 14%) presents challenges for GeSn growth and stabilization due to Sn segregation and phase separation. Chemical vapor deposition techniques at low temperatures can however allow the growth of good quality GeSn layers with high Sn contents (> 10 at.%). The use of these kinds of layers can thus be envisioned for electronic and opto-electronic devices.

Contact technology for GeSn layers

Contact technology for GeSn layers

For the technological requirements of the different applications listed in section 1.1, a metallic contact would be necessary to integrate in the process flow for device elaboration to allow the pass of electric current. Indeed, thin-film metallizations allow to have a direct contact with the semiconductor material and to connect semiconductor devices from one to another on an integrated circuit. Due to its technological importance, metallic contacts have been an important research topic in the semiconductor domain. Fundamental research has focused on the elucidation of charge transfer mechanisms and to understand and characterize the junction properties on the atomic scale. Meanwhile for industry, from a device point of view, the control of barrier heights for a specific set of semiconductor and metal components is critical to properly design high performance micro-and opto-electronic devices. Even if great technological progess have been achieved, further metallurgical approaches had to be related to a deeper understanding of the metal / semiconductor properties, physics, chemistry and interactions. This way, current issues of device performance, reliability and reproducibility can be more easily solved [START_REF] Brillson | Contacts to Semiconductors: Fundamentals and Technology[END_REF]. For these reasons, it is important to identify the parameters that characterize a good contact, the ways to obtain them and the current considerations for GeSn layers.

Metal / Semiconductor (M / S) contact formation and properties

Basically, a contact refers to a metallization applied to a semiconductor to ensure good electrical and physical conduction towards the semiconductor. However, simply performing a metal deposition on a semiconductor surface is not enough to ensure high quality contacts. Indeed, a system formed by two materials, (metal / semiconductor for this case), M and S separated by a sharp M / S interface (as displayed in Fig. 1.8 (a)), will not be thermodynamically stable [START_REF] Gas | Formation of silicide thin films by solid state reaction[END_REF]. With enough time or temperature, the system will then evolve towards a more stable state in which the energy can be minimized. This state can be achieved by the formation of M x S y intermetallics as observed in Fig. 1.8 (b). This evolution, known as reactive diffusion, can occur in the solid state and will depend on the M / S equilibrium phase diagram. For totally miscible elements a solid solution of concentration ranging from pure M to pure S will occur and the M / S interface will be spread. For systems with a tendency to demixing, two solid solutions (M-rich and S-rich, respectively) will form and the M / S interface will change to a M(S) / S(M) interface. In most cases, intermediate phases exist Chapter I. GeSn material from growth to applications between M and S and they could be observed, at specific annealing temperatures, if they are stable [START_REF] Gas | Formation of silicide thin films by solid state reaction[END_REF]. However, even if the binary M / S phase diagram is useful to have a first hint about the phases that could develop during the solid-state reaction, the reaction of thin films could differ from bulk reaction.

It has been well established experimentally, by planar binary thin-film diffusion couples, that not all of the compound phases predicted by the equilibrium phase diagram are observed [START_REF] Gösele | Growth kinetics of planar binary diffusion couples: Thinfilm case versus bulk cases[END_REF]. Additionally, sequential appearance of phases, absence of certain stable phases or the nucleation of metastable phases coupled with rapid kinetics of formation was observed. These effects were related to (i) variations on the experimental conditions (quality of the M / S interface, nature of the substrate, presence of impurities) which had important consequences on the growth of the thin films, (ii) the use of low temperatures which might have favored the observation of phenomena related to nucleation and (iii) the limited thicknesses which can underline the competitive aspects of the growth of different phases.

For diffusion controlled reaction, the high rates of formation as well as the low values of the activation energies were related to the microstructure of the growing layer. The high density of grain boundaries and the high diffusivity along these paths can generate high rate growth kinetics. Other factors, such as chemical driving force and structural defects, were expected to have smaller influence [START_REF] Gas | Formation of silicide thin films by solid state reaction[END_REF]. For these reasons, relaying on bulk binary phase diagrams to foresee the system phase evolution will not be reliable enough. A full analysis of the phase sequence during the solid-state reaction is mandatory once a new couple metal / semiconductor has to be integrated in micro-electronic or opto-electronic devices. This analysis should be additionally coupled to a full assesment of the morphological and electrical properties and ranges of stability to supply Si CMOS-compatible contacts of high quality [START_REF] Rodriguez | Development of a CMOScompatible contact technology for IIIV materials and Si photonics[END_REF]. Indeed, certain general requisits should be fulfilled when an intermetallic compound is used to contact a semiconductor layer as follows [START_REF] Ghegin | Integration of innovative ohmic contacts for heterogeneous III/V-Silicon Photonic devices[END_REF].

-Low resistivity.

-Good adhesion to the semiconductor.

-Low contact resistance to the semiconductor.

-Appropriate Schottky barrier height or ohmic behavior with heavily doped semiconductors.

-Thermal stability in a wide range of temperatures.

-Appropriate / smooth morphology for subsequent lithography and etching processes.

-High corrosion resistance.

-Oxidation resistance.

-Good adhesion to and minimal reaction with the surrounding materials (other than the semiconductor).

-Low stress variations.

Contact technology for GeSn layers

-Compatibility with other processing steps such as lithography and etching.

-Minimizing metal diffusion.

-High electromigration resistance.

-Formation at low temperatures.

The verification of all these parameters is out of the scope of this thesis. However, the combination of multiple characterization techniques which will be further described in chapter II, can help to assess the intermetallic properties evolution during thermal annealing. Indeed, a great amount of optical (transmission spectroscopy, raman spectroscopy, reflectance modulation, etc.) physico-chemical (electron beam techniques, ion beam techniques, X-ray techniques, etc.) and electrical (resistivity, carrier concentration, mobility, contact resistance, barriers height, etc.) characterization techniques are available to fully understand the M / S interactions and behavior [START_REF] Mcguire | Characterization of semiconductor maerials: Principles and Methods[END_REF].

For our particular case the substrate corresponds to GeSn layers. For these kinds of semiconductor layers, the use of Ni stanogermanides, i.e. Ni(GeSn), has been proposed in order to start the phase sequence at low temperatures, reach low specific contact (R c ) and sheet (R sh ) resistances, while they are CMOS-compatible [START_REF] Han | High-mobility germanium-tin (GeSn) P-channel MOSFETs featuring metallic source/drain and sub-370 process modules[END_REF][START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. For these reasons, different complementary characterizations techniques (described in chapter II) will be used to analyze either Ni / GeSn intermetallics (presented in chapter III) or enhance Ni-based metallization (presented in chapter IV). The objective is to identify the most suited metallization or an alternative process scheme that ensures high quality, stable and reliable GeSn-based contacts. In general lines, the study will be focused on (i) the full assessment of the phase sequence during the solid-state reaction, (ii) the determination of the ranges of morphological and electrical stability, (iii) finding strategies that could either reduce resisitivity values or enhance thermal stability. The results obtained will be described in the next chapters.

Metallization for GeSn-based layers

Different studies have addressed the formation of Ni-based intermetallics on GeSn since 2010. Takeuchi et al. investigated the growth of Ni(GeSn) layers and Sn incorporation by X-ray diffraction (XRD) and transmission electron microscopy (TEM) techniques. Meanwhile, scaning electron microscopy images (SEM) of the Ni(GeSn) layers top surface evidenced important surface roughening and low thermal stability [START_REF] Takeuchi1 | Assessment of Ge 1-x Sn x alloys for strained Ge CMOS devices[END_REF]. Later, Vincent et al. were one of the first groups to use complementary techniques to assess different properties of Ni / GeSn intermetallics. They studied the phase sequence of 10 nm Ni on 50 nm Ge 0.95 Sn 0.05 on Ge layers by real-time XRD measurements, the evolution of the electrical properties by sheet resistance (R sh ) measurements at 400 after several minutes of annealing and the agglomeration of the layers was confirmed by scanning electron microscopy (SEM) images of the intermetallic surface [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF]. After these results, many other groups have analyzed the Ni / GeSn solid-state reaction. However, the metallization parameters varied from one group to the other, from the surface pre-clean to the annealing parameters, as summarized in Table 1.2. The multiple combination of process parameters for Ni / GeSn contact technology can impact the different properties of the intermetallics. The phase sequence, the evolution of the morphological or the electrical properties are expected to be impacted during the solid-state reaction when more than one process parameter is changed. The state of the art concerning Ni / GeSn system evolution is further adressed in chapter III, section 3.1 to understand the different properties of the Ni / GeSn intermetallics. This thesis focalized in a systematic and comprehensive analysis of the Ni / GeSn system evolution with annealing temperatures. The objective, in the first instance, is to fully comprehend the Ni / GeSn solid-state reaction and the properties evolution. After this, an analysis of the impact of changing different process conditions can be made. From this study we expect to propose a proper metallization or an adequate process scheme to elaborate contacts for GeSn-based technology.

To summarize this section, an intermetallic layer is needed to contact GeSn-based devices. Indeed, the formation of M x S y intermetallics corresponds to a more stable configuration than a sharp M / S interface.

In order to ensure high quality, stable, reliable and low resistive contacts a full understanding of the phase sequence during the solid-state reaction is necessary. Ni / GeSn intermetallics are considered as well suited to fulfill the requirements of GeSn-based technology. However, the multiple combination of process parameters that can be chosen to treat the Ni / GeSn system (from surface preparation to annealing treatments) is large. These variations can impact the different properties of the intermetallics and their evolution when annealed. For these reasons, a systematic and comprehensive analysis of the Ni / GeSn system evolution with annealing temperatures (with specific process conditions) was carried out and is presented in the next chapters.

Chapter I. GeSn material from growth to applications

Conclusion of Chapter I

For this chapter, a brief description of GeSn layers properties, historical background, envisioned applications and a GeSn-based contact technology overview was realized. In general, this alloy corresponds to a group-IV material which present a direct band-gap. This properties are advantageous for different applications for the semiconductor industry. The main issues of this material are the low solubility of Sn in the Ge lattice (< 1 at.%) and the important lattice parameter differences (> 14%). These aspect can difficult the growth of stable GeSn layers, phase separation and Sn segregation can easily occur. Even though, high quality GeSn layers with high Sn content (> 10 at.%) can be achieved with low temperature CVD growth conditions. One important aspect to have in mind then, for GeSn-devices, corresponds to the develop and elaboration of high quality, stable and reliable contacts.

For the case of thin-films M / S couples, the minimization of the energy will allow the growth of M x S y intermetallic phases. These phases will serve as an intermetallic contact. For the case of GeSn-based layer, Ni metallization has been considered as a well suited solution. However, the metallization parameters can vary importantly. Combined characterization techniques will be then necessary to asses certain critical properties of the layers, as will be presented in chapter II. The phase sequence during solid-state reaction, the ranges of morphological and electrical stability when annealing, together with other properties can be acquired from these techniques. The final goal is to be able to proposed a suited metallization or process scheme that can ensure reliable GeSn-base contacts that can be fully integrate in diverse electronic or opto-electronic devices. The elaboration process of a semiconductor device can demand many steps, in which the quality of the integrated materials should be high. In order to ensure that all the components fulfill the standard demands, the properties and performance should be probed and measured. The only way to do it, is with the consecutively use of different characterization techniques, which allow to understand and estimate the material properties.

Chapter II

Characterization techniques for intermetallic contacts

A metal / semiconductor (M / S) contact is a fundamental component of numerous electronic and opto-electronic devices. Indeed, this M / S junction will allow the exchange of electronic charges between the semiconductor device and an external circuit. For these reasons, the characterization and improvement of the contact components have been the topic of active research for decades [START_REF] Brillson | Contacts to Semiconductors: Fundamentals and Technology[END_REF]. Fundamental studies have focused on elucidate and understand certain physical properties that are key parameters for conceiving a device as the charge transfer mechanisms, the junction properties in an atomic scale, the ways to control effective barrier heights, etc., together with the impact of reducing the size of the components to the properties evolution. As described in chapter I, section 1.2, from a device point of view the reliability, stability and high electrical performance of the metallic contact are critical for micro/nano and opto-electronics.

A comprehensive and complete understanding of the interface physics, chemistry and its evolution when changing certain process parameters (temperature, alloy element, surface treatment, etc.) was thus considered as necessary. For these aims, different characterization techniques were chosen thanks to their abilities to generate detailed and accurate information about the contacts properties either on the as-deposited state and during its evolution when the metal / semiconductor solid-state reaction occurred. In the following sections, different techniques well adapted to characterize our kinds of samples will be described, as schematized in Fig. X-Ray diffraction (XRD) analysis was applied to determine the phase sequence, composition and texture of the layers. Atomic force microscopy (AFM) and transmission electronic microscopy (TEM) coupled to energy-dispersive X-ray spectroscopy (EDS) and Structural characterization by X-ray diffraction electron energy loss spectroscopy (EELS) analyses helped to understand the morphology of the layers, the composition and atom distribution. The electrical characterization was enabled thanks to the four-point probe method to follow sheet resistance (R sh ) evolution.

Structural characterization by X-ray diffraction

X-rays are high-energy electromagnetic radiation of exactly the same nature as light but of very much shorter wavelength [START_REF] Suryanarayana | X-Ray Diffraction: A Practical Approach[END_REF]. In 1912, X-ray diffraction by crystals was discovered and this discovery provided a new non-destructive method for investigating the fine structure of matter [START_REF] Cullity | Elements of X-Ray Diffraction[END_REF]. When an X-ray beam falls on an atom the beam may be scattered, due to process of absorption and re-emission of electromagnetic radiations [START_REF] Warren | X-Ray Diffraction[END_REF]. Then, the incident radiation can be re-emitted in different directions. This random emission can generate interference, which are the superposition of two or more of these scattered waves, producing a resultant wave that is the sum of the overlapping wave contributions. Further, when constructive interference of more than one scattered wave happen, we can talk about diffraction [START_REF] Suryanarayana | X-Ray Diffraction: A Practical Approach[END_REF]. Diffraction is the phenomenon that explains the modification of the behavior of an X-ray beam by its interaction with an object.

Bragg-Brentano θ-2θ diffraction configuration

The information acquired, when using X-ray radiation, can be used to determine the crystal structure following the Bragg's Law [START_REF] Bragg | The diffraction of short electromagnetic waves by a crystal[END_REF].

n.λ = 2.d hkl . sin(θ) (2.1)
Where λ is the wavelength of the incident beam wavelength, n is an integral number of wavelength or order of diffraction, θ is the incident angle and d hkl is the atomic interreticular distance, i.e. the distance between two atomic planes that belong to the same (h k l). Fig. 2.2 shows the schematic representation of diffraction on atomic planes. The use of X-ray diffraction enables to determine and verify then the crystal structure, the chemical composition and the texture of the layers, as well as the evolution of these properties when changing certain process parameters. The θ-2θ Bragg-Brentano configuration is a type of geometry in which the X-ray source (that irradiates the sample) is fixed, and the detector moves through a range of angles. The radius of the focusing circle is not constant but increases as the angle 2θ decreases, as it can be seen in Fig. 2.3. The 2θ measurement range is typically from 0°to about 170°. In an experiment scan the whole range of the detector angles is yet not necessary [START_REF] Suryanarayana | X-Ray Diffraction: A Practical Approach[END_REF][START_REF] Guinebretiere | X-ray Diffraction by Polycrystalline Materials[END_REF]. The angle between the plane of the specimen and the X-ray source is θ, the Bragg angle, while the angle between the projection of the X-ray source and the detector is 2θ. For this reason the patterns produced with this geometry are often known as θ-2θ scans and are considered as one of the most common configuration in order to obtain crystallographic information of a sample from a XRD source [START_REF] Suryanarayana | X-Ray Diffraction: A Practical Approach[END_REF][START_REF] Guinebretiere | X-ray Diffraction by Polycrystalline Materials[END_REF]. Indeed, this technique allows to acquire diffraction patterns which gather information from the lattice planes that are parallel to the surface samples, as schematized in Fig. 2.4 (a). These patterns are unique and characteristic of the crystallographic structure of the phase that is present and thus allows their identification. This technique has been largely used for the phase identification of many metal / semiconductor systems, for contact Structural characterization by X-ray diffraction elaboration as III-V materials [START_REF] Ghegin | Integration of innovative ohmic contacts for heterogeneous III/V-Silicon Photonic devices[END_REF], silicides [START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Mangelinck | Kinetics of a transient silicide during the reaction of Ni thin film with (1 0 0) Si[END_REF], germanides [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Nemouchi | Simultaneous growth of Ni 5 Ge 3 and NiGe by reaction of Ni film with Ge[END_REF][START_REF] Nemouchi | A comparative study of nickel silicides and nickel germanides: Phase formation and kinetics[END_REF] and stanogermanides [START_REF] Han | High-mobility germanium-tin (GeSn) P-channel MOSFETs featuring metallic source/drain and sub-370 process modules[END_REF][START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Schulte-Braucks | Process modules for GeSn nanoelectronics with high Sn-contents[END_REF] contacts. θ-2θ scans over a large range of 2θ position (around 20-90°), using a cooper (Cu Kα) source were considered adequate to characterize our kinds of samples. Analyses were performed in a X'Pert Pro PANalytical X-ray diffractometer. An example of the results that can be obtained with this technique is displayed in Fig. 2.5. It will be seen in the following chapters, that this technique was well-aimed in order to follow the phase sequences during the solid-state reaction, the preferential orientation of the intermetallics, the changes in the out-of-plane texture and to estimate the out-of-plane crystallite size. For phase identification, of Ni / GeSn intermetallics, the obtained patterns were compared to the experimental patterns extracted from the wide database of the International Centre for Diffraction Data (ICDD) PDF-4 database. For the kinds of analyses explained before, samples were treated by rapid thermal annealing (RTA) at different temperatures and analyzed when cooled down to room temperature. This "cooking and looking" method has the risk of overlooking transient but crucial stages during the reaction. In order to reduce this risk, in-situ θ-2θ analyses scans over a large range of diffraction angles (20-90°) were performed with a 2°offset in order to attenuate the symmetric (0 0 4) reflection of the Si substrate. Indeed, Demeleumester et al. and Vincent et al. used these kinds of analyses in order to properly follow the phase sequence evolution with temperature, for Ni / GeSn systems [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. In our study, an Empyrean PANalytical X-ray diffractometer equipped with a copper (Cu Kα) source, a linear detector (PIXcel1D) and an HTK 1200 Anton Paar furnace under secondary vacuum (the pressure was maintained at about 10 -6 bar, equivalent to 0.1 Pa, to avoid atmospheric contamination) was used. The kinds of patterns that were obtained are displayed in Fig. 2.6. For these measurements, samples were placed in the HTK 1200 Anton Paar furnace support and in the diffractometer. Then, samples were annealed from 50 up to 600-800 (depending on the sample), with a 5 step and scanning time of 7 min per step as displayed in Fig. 2.6 (a). A pattern was acquired each 5 , and was plotted with the diffraction angle on the x axis, the temperature on the y axis and the intensity variation on z axis as a 3D pattern, as shown in Fig. 2.6 (b). Then, these patterns were assembled to obtain a 2D contour map with the diffraction angle on the x axis and the temperature on the y axis, with the respectively intensity scale bar, as displayed in Fig. 2.6 (c). The pattern analyses, for either classical θ-2θ and in-situ XRD, were realized with the High-Score Plus software from PANalytical and the ICDD PDF-4 database.

In-plane reciprocal space map (RSM)

The classical θ-2θ configuration is well adapted in order to characterize polycrystalline films. However, when thin films are used, the intensities can be low compared to the ones related to the substrate. In addition, the θ-2θ configuration only allows to recover information about grains which crystallographic planes have a parallel orientation with respect to the sample surface. Information about grains oriented without fulfilling the Bragg's condition could then be lost.

In-plane configuration with a grazing incident X-ray beam can thus be a solution in order to increase the signal obtained from the thin films [START_REF] Kobatashi | X-ray thin film measurement techniques: In-plane XRD measurements[END_REF]. Contrary to the θ-2θ config-Structural characterization by X-ray diffraction uration, an in-plane set up allows to acquire crystallographic information concerning the planes that are normal to the surface of the sample as observed in Fig. 2.4 (b). Information is obtained, thanks to the control of the sample rotation angle (φ) and the position of the detector (2θ χ ) [START_REF] Kobatashi | X-ray thin film measurement techniques: In-plane XRD measurements[END_REF], as shown in Fig. 2.7. For these reasons, complementary in-plane reciprocal space map (RSM) characterizations were carried out using a Rigaku SmartLab X-ray diffractometer. The grazingincidence angle was fixed at 0.5°and the exit angle for detection at 1°. The in-plane RSM were obtained using an in-plane 2θ range of 20-90°, for a φ range of 95°explored with 1°s teps.

The kinds of patterns that can be obtained with this method are exemplified in Fig. 2.8 (a). After integration over the full φ range (see Fig. 2.8 (b)), an equivalent phase diagram to θ-2θ in-plane patterns can also be obtained, see Fig. 2.8 (c). The interpretation of these kinds of patterns can be very helpful to additionally identify peaks shift, broadening or overlap. Indeed, this technique was used to determine the in-plane orientation of the Ni / GeSn intermetallic phases, verify the nature of certain peaks displacements in 2θ position and determine the texture of the intermetallic phases, as will be observed in further chapters. As mentioned before, one of the advantages of using this technique is to increase the signal gathered from the thin intermetallic film compared to the substrate. Indeed, in the Bragg-Brentano configuration the incident angle is large which allows a large penetration inside the samples. Then, diffraction intensities from the substrate can be quite large and can otherwise reduce the level of detection for the thin film signal. The opposite case occurs when a grazing incident X-ray beam is used in which the incident angle is rather small. The beam can thus travel along the thin film recovering larger information about it without penetrating in the substrate. However, if the sample surface is extremely flat and the incident angle extremely small, the incident X-ray beam might be totally reflected. This occurs because the incident angle is that small that the X-ray beam cannot properly penetrate in the sample, instead it travels parallel to the sample surface. This generates patterns with diffraction intensities very weak, which can make the indexation difficult [START_REF] Kobatashi | X-ray thin film measurement techniques: In-plane XRD measurements[END_REF].

In-plane Pole Figures (IP-PF)

When thin films are used, the crystal orientation and the degree of preferred orientation can have a great influence on their properties [START_REF] Nagao | X-ray thin film measurement techniques: Pole figure measurement[END_REF]. Then, in order to have a better understanding of the microstructure of the films later used for devices, a full assessment of the Ni / GeSn texture was realized thanks to in-plane pole figures (IP-PF). The advantage of PF measurements by XRD corresponds to the possibility to acquire the texture information of samples averaged over a large area / volume ratio [START_REF] Nagao | X-ray thin film measurement techniques: Pole figure measurement[END_REF].

A pole figure is a stereographic projection, with a specific orientation relative to the specimen. This projection shows the variation of pole density with pole orientation for a selected set of crystal planes [START_REF] Cullity | Elements of X-Ray Diffraction[END_REF], in which a scattering vector of fixed length is rocked to trace out a sphere, as shown in Fig. 2.9. For these measurements, the combination of the out-of-plane and in-plane configuration geometries were used as observed in Fig. 2.10. This kind of set-up uses the in-plane arm like for the in-plane RSM analysis, while the movement of the four axes ensures the gathering of the whole signal. The control of the χ axis (which corresponds to the Structural characterization by X-ray diffraction tilt angle of the sample) is realized with the combination of 2θ and 2θ χ . Meanwhile the detector is fixed at the position of interest, φ is scanned at each χ angle.

When analyzing a sample, the diffraction angle 2θ is fixed and the diffracted intensity is collected by varying two geometrical parameters the χ angle (tilt angle from the sample surface normal direction) and the φ angle (rotation angle around the sample surface normal direction). The obtained diffracted intensity data are plotted as a function of χ and φ on the plane of the stereographic projection with polar coordinates, as shown in Fig. 2.9 [START_REF] Nagao | X-ray thin film measurement techniques: Pole figure measurement[END_REF]. Pole figure analyses have enable to (i) identify a new type of complex texture called "axiotaxy", in which an alignment of the thin film and substrate plane is possible thanks to similar d-spacing values, while the fibre texture lies off-normal to the surface [START_REF] Detavernier | An off-normal fibre-like texture in thin films on single-crystal substrates[END_REF]. (ii) Evaluate and understand the texture structure of Ni / Ge intermetallics on different Ge substrate and its evolution with annealing temperatures [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Balakrisnan | Texture of NiGe on Ge (0 0 1) and its evolution with formation temperature[END_REF][START_REF] Van Stiphout | Formation of ultrathin Ni germanides: solid-phase reaction, morphology and texture[END_REF][START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF][START_REF] Huang | Texture evolution and grain competition in NiGe film on Ge(0 0 1)[END_REF].

This characterization technique was thus considered as a need to understand the Ni / GeSn intermetallic texture. Complete in-plane pole figures were then recorded using a Smartlab set-up, with equivalent steps in φ and χ of 3°. The pole figures were then built using the 3D Explore software (Rigaku Corporation, Tokyo, Japan). The pole figure analysis utilized simulations with the CaRIne 3.1 software, which is a software that deals with different geometrical representation of crystals and crystalline structures. An example of the kinds of results that were acquired with this technique is presented in Fig. 2.11, together with the analysis procedure performed with the CaRine 3.1 software. Fig. 2.11 presents the type of circular stereographic patterns that can be acquired. Indeed, the data are obtained in the way that bright spots at specific χ and φ positions can be observed for this Ni / GeSn sample. In order to identify the poles associated to the germanium (Ge) substrate, a simulation with the CaRine 3.1 software was done. This program allows to (i) construct and simulate the phase lattice (for this case the Ge lattice simulation), (ii) give the lattice a specific orientation (Ge (1 0 0) oriented for this case) and (iii) choose the poles that have to be arranged in the polar stereographic projection (for this case Ge (3 3 1) poles).

The chosen orientation of the lattice corresponds to the real orientation of the substrate or to the Ni / GeSn intermetallics in the out-of-plane direction (which can be gathered from θ-2θ XRD analyses). Meanwhile, the projected (h k l) poles correspond to the poles that can diffract at the 2θ angle chosen to record the pole figure (2θ = 76°for this case). As mentioned before, the combination of the out-of-plane θ-2θ and the inplane configurations allows to acquire complete pole figures. Basically, when χ = 0°the geometry corresponds to Bragg-Brentano configuration, and when χ = 90°, the in-plane configuration is achieved. Then, like for the case of in-plane RSM analyses, a geometrical constrain can jeopardize the diffracted information. Indeed, from χ = 80°up to χ = 90°, the relative intensity received can be importantly decreased. Then, acquiring the crystalline information from the poles at these coordinates can be difficult.

Morphological characterization

Morphological characterization

Metal / semiconductor contact surfaces generally show a flat interface after deposition. However, one technological requirement for a good metal contact is the preservation of interface flatness upon subsequent annealing during device processing and in a wide range of temperatures [START_REF] Brillson | Contacts to Semiconductors: Fundamentals and Technology[END_REF]. Silicides, germanides and stanogermanides all suffer from degradation after a certain threshold temperature. This degradation can compromise the electrical properties, generate leakage currents and unpredictable electrical performance. For these reasons atomic force microscopy (AFM) and transmission electron microscopy (TEM) images were systematically realized to verify the morphological evolution of the layers.

Atomic Force Microscopy (AFM) Tapping mode

Atomic force microscopy (AFM) is a method used to observe and monitor the shape of a surface with three dimensional (3D) details down to the nanometer scale [START_REF] Haugstad | Atomic Force Microscopy: Understanding Basic Modes and Advanced Applications[END_REF]. An AFM physically "feels" the sample surface with a sharp probe, building up a map of the height of the sample surface or topography as it goes along [START_REF] Eaton | Atomic Force Microscopy[END_REF]. The surface morphology is not perceived in the usual way, that is, by line-of-sight, reflections or shadows. Rather, at each point or pixel, within a 2D array over the surface, a measurement of surface height is made using a sharp solid force probe. It essentially uses the touch to be able to image a surface, unlike light or electron microscopes. The force probe may move over a stationary sample or remain stationary as the sample is moved under the probe. After this process, an image of the surface topography is obtained [START_REF] Haugstad | Atomic Force Microscopy: Understanding Basic Modes and Advanced Applications[END_REF]. Fig. 2.12 presents a schematic representation of the AFM principle. In the intermittent, or tapping, mode of imaging the cantilever is allowed to oscillate at a value close to its resonant frequency. When the oscillations occurs close to a sample surface, the probe is repeatedly engaged and disengaged with the surface, restricting the amplitude of oscillation, as observed in Fig. 2.12. As the surface is scanned, the oscillatory amplitude of the cantilever changes as it encounters different topography details. By using a feedback mechanism to alter the z-height of the piezo-crystal and maintain a constant amplitude, an image of the surface topography is obtained [START_REF] Eaton | Atomic Force Microscopy[END_REF].

The advantage of this technique consists in its ability to image a wide variety of samples, defined on the nanoscale range with an atomic resolution, as observed in Fig. 2.13, while the operation is rather simple. Indeed, a proper idea of the surface features can be acquired together with the quantitative information about the sample roughness. Fundamental applications in physical sciences include imaging of the fine structure of metals and semiconductors. Specifically for the semiconductor industry, the AFM use started around the mid-1990s, for wich one of the main application was in metrology [START_REF] Eaton | Atomic Force Microscopy[END_REF]. One important parameter to evaluate would then be the surface roughness, to properly characterize the surface of the deposited metallization. This parameter is easy to measure with the AFM, because it produces high contrast on relatively flat surfaces, and produces three-dimensional, digital data by default, which gives accurate values for surfaces with roughness down to the level of atomic flatness (as observed in Fig. 2.13). On the other hand, measuring roughness is a simple and quantitative way to compare surfaces [START_REF] Eaton | Atomic Force Microscopy[END_REF].

Nishimura et al. used AFM measurements to verify (i) the roughness evolution with temperature, (ii) the surface degradation and agglomeration, (iii) the grain size of the agglomerates and (iv) the impact of Sn content variation on roughness variation, when analyzing Ni / GeSn samples [START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF]. Scanning Electron Microscopy (SEM) technique was also used for these aims [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF]. However, this technique can give the impression that surfaces with roughness values less than 10 nanometers are extremely smooth and featureless, due to the fact that SEM detects electrons scattered from a few nanometres into the surface. In addition, it cannot supply data suitable to measure roughness parameters [START_REF] Eaton | Atomic Force Microscopy[END_REF].

It will be seen in the following chapters, that the capacity to easily compare Ni / GeSn surfaces, with AFM images allowed us to properly evaluate the surface roughness evolu-Morphological characterization tion with temperature and when varying other process conditions as the alloying element, the surface preparation, the GeSn amorphization prior to the metallization step and follow the intermetallics agglomeration and the Sn segregation mechanisms.

On the other hand, one aspect to have in mind when using this technique, is to ensure the same parameters when scanning, processing and analyzing the images to properly compare the results. Two of the tools to process AFM raw images corresponds to "Plan-eFit" and "Flatten" commands from the NanoScope Analysis software. These commands level the image by operations realized in the entire image or following a line by line basis, respectively. This way, thanks to a polynomial fitting, the noise can be subtracted. The use of these commands can change the contrast of the images, which varies with the polynomial order that can be chosen. If different parameters are chosen, one image can look smother than another one, but this can only be an artifact due to image processing parameters. Other artifacts, that can modify the images obtained and can lead to a wrong analysis of the results can come from the AFM tool or the sample itself. Indeed, if the probe is either damaged, broken or contaminated, it can be responsible of the following facts (i) the features on a surface can appear too large or too small, (ii) the appearance of strangely shaped objects or repeating patterns and (iii) the generation of non homogeneous images which appear normal (compared to what is expected) on the top or bottom of the image, but not on their sides [START_REF] Eaton | Atomic Force Microscopy[END_REF]. Moreover, an external contamination present in the sample prior to the analysis can avoid a proper inspection of the sample surface and rather allow the probe to break or get damaged.

Transmission Electron Microscopy (TEM)

Transmission electron microscopy (TEM), corresponds to a microscopy technique in which a specimen is irradiated with an electron beam. Thanks to the electron interactions with the sample, an image with high resolution can be obtained [START_REF] Reimer | Transmission Electron Microscopy: Physics of Image Formation and Microanalysis[END_REF]. Thanks to its high capability to image a sample, this technique was used to evaluate the interface state and morphological evolution when the process conditions changed. In addition, X-ray energydispersive spectrometry (EDS) and electron energy-loss spectrometry (EELS) analyses were coupled in order to map atoms distributions.

An X-ray energy-dispersive spectrometer (EDS) is a sophisticated instrument in which a detector generates voltage pulses that are proportional to the X-ray energy. This technique uses the fundamental principle that the elements have different and unique atomic structures, to gather a specific set of peaks due to its electromagnetic emission. Electronic processing of the pulses translates the X-ray energy into a signal that enables elemental or chemical characterization. The input can be transformed and displayed as a spectrum or a quantitative compositional profile or image or "map" of the species that are present in the sample [START_REF] Williams | Transmission Electron Microscopy: A Textbook for Materials Science[END_REF]. An example of the kinds of images that can be acquired is presented in Fig. a specific temperature [START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF].

On the other hand, electron energy-loss spectrometry (EELS) corresponds to the analysis of the energy distribution of electrons that come through the specimen. This technique is considered as an excellent complement to EDS analysis, since it offers substantially more information than mere elemental identification and is well suited to the detection of light elements. A similar image, like for the EDS case, can be generated, which provides information on elemental composition, based upon the atomic transition during electron-electron interactions [START_REF] Williams | Transmission Electron Microscopy: A Textbook for Materials Science[END_REF].

One limitation to the use of TEM techniques, to obtain images from the samples, is that it present 2D images of 3D specimens, viewed in transmission. This drawback can be called the projection-limitation and occurs because generally all the TEM information is averaged through the thickness of the specimen. Then, there is basically no depth sensibility. Meanwhile, EDS technique is limited by its poor energy resolution which makes a proper quantitative analysis difficult. Artifacts in the spectra and overlapping of X-ray emissions from different species can also prevent a proper image analysis, additionally it is difficult to detect light elements. EELS technique, on the other hand, can quantify all the elements in the periodic table and is especially good for analyzing the light elements, offering better spatial resolution and analytical sensitivity than the EDS technique. However, it can be a challenging experimental technique which requires very thin specimens while understanding and processing the spectra and images requires somewhat more of a physics background than EDS. For these reasons, both are complementary and can be quite useful when analyzing thin intermetallic films [START_REF] Williams | Transmission Electron Microscopy: A Textbook for Materials Science[END_REF].

Electrical characterization

Electrical characterization

In order to meet the increasing requirements for the semiconductor industry, different methods and techniques can be used to acquire precise and relevant electrical information for electronic and opto-electronic devices as resistivity, carrier concentration, mobility, contact resistance and so on [START_REF] Mcguire | Characterization of semiconductor maerials: Principles and Methods[END_REF]. Particularly, the metallic contacts for devices need to ensure low values of sheet (R sh ) and contact (R c ) resistance. These properties should additionally remain stable during the device processing. Indeed, when thin films are used, the sheet resistance (or surface resistivity) is an important property to quantify. It allows to determine the ability for charges to travel along uniform thin films. Furthermore, resistivity (and hence conductivity) can be determined from the measurement of the sheet resistance, enabling electrical characterization of a material [START_REF] Smits | Measurement of sheet resistivities with the four-point probe[END_REF]. We then decided to used a common technique in order to obtain the sheet resistance values by means of the four-point probe method.

Four-point probe electrical measurements

A very convenient method for measuring the R sh is the four-point probe technique, which is a well established technique used for semiconductors [START_REF] Mcguire | Characterization of semiconductor maerials: Principles and Methods[END_REF][START_REF] Smits | Measurement of sheet resistivities with the four-point probe[END_REF]. It is performed using four electrical probes (in line) with equal spacing between each set of probes, as shown in Fig. 2.15. Then, a current is applied between the outer two probes and a voltage drop is measured between the inner two probes. The sheet resistance is related to the current and voltage from the following equation

R sh ∝ V / I (2.2)
Here, R sh is the sheet resistance, V is the change in voltage measured between the inner probes, and I is the current applied between the outer probes. The sheet resistance is expressed with the units Ω/sq or ohms per square [START_REF] Smits | Measurement of sheet resistivities with the four-point probe[END_REF]. It will be seen in next chapters, that Ni / GeSn intermetallics have a low thermal stability, which promotes Chapter II. Characterization techniques for intermetallic contacts agglomeration and Sn segregation when temperatures rise. Then, the annealing temperature become one of the most important process parameters when elaborating the contacts. R sh evolution was consecutively measured with the annealing temperatures in order to evaluate the electrical performance of the layers and evaluate the impact of changing another process parameters (i.e. the alloying element, the surface treatment, etc.). For Ni-stanogermanides, R sh evolution with temperature enabled to (i) link electrical variation with the phase transformation, (ii) determine the temperature range from which minimum and stable values can be obtained and (iii) verify the layers degradation and stability variation depending on the Sn content of the layers [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Lasing in direct-bandgap GeSn alloy grown on Si[END_REF][START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF].

Sheet resistance measurements with the four point-probe method was thus considered as a well adapted technique for our purposes. However, certain attention should be taken concerning the probes used for the analysis. If the probes are extremely sharp, they can scratch the surface sample which can make difficult to obtain R sh measurements and can generate artifacts when analyzing the surface morphology.

Conclusions of Chapter II

The first motivation of our work was to elaborate well adapted Ni-based contacts for GeSn-based devices. However, this motivation cannot be fulfilled if the contacts are not properly tested, probed and characterized. A summary of the main characterizations techniques applied to (i) acquire valuable information about the properties of the Ni / GeSn contacts and (ii) follow the evolution of these properties when certain process conditions were changed, was then addressed in this chapter.

In order to properly learn and understand the Ni / Ge(Sn) phase sequence, crystalline orientation and degree of texture, complementary XRD techniques (θ-2θ patterns, inplane RSM and in-plane pole figures) were used. Meanwhile, morphological properties concerning surface and cross-section were assessed with AFM and TEM analyses, respectively. TEM images were additionally coupled to EDS and/or EELS techniques in order to follow atoms accumulation or segregation. Finally, electrical characterization was mainly possible thanks to the four-point probe method. This quick and easy technique allowed to properly extract the R sh values of the samples.

Chapter III

Study of the Ni / GeSn system during the solid-state reaction

State of the art: Ni / GeSn system

Previously in chapter I, it was highlighted the need to form efficient and reliable contacts for GeSn-based electronic and opto-electronic devices. Additionally, it was discussed how an intermetallic layer can be formed (see chapter I, section 1.2.1) and the well suited characterization techniques that are used to fully assess its properties in terms of phase sequence, crystallinity, morphological and electrical properties (see chapter II). This chapter is divided in seven section as schematized in Fig. The state of the art concerning Ni / GeSn intermetallics and the experimental procedure used for this precise study will be introduced first. The contribution of our results to the current state of the art will be presented then. The results are described in terms of the phase sequence, the morphological and the electrical properties evolution during the solid-state reaction. In addition, comprehensive analyses of the impact of the surface preparation, the nature of the Ni-rich phase, the impact of Sn content variation and the texture of the Ni(GeSn) phase will be addressed.

State of the art: Ni / GeSn properties and evolution with annealing temperatures

As expressed in chapter II, the technological importance of the stanogermanides requires to have a full understanding of their properties. Indeed, these properties can have a direct impact on (i) the fabrication steps to follow when conceiving an electronic or an optoelectronic device and (ii) the final performance of the device [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. In order to predict, understand and be able to adjust or control certain properties and process parameters of Ni / GeSn intermetallics, a proper analysis of the phase sequence during and the properties evolution during the solid-state reaction, is then fundamental. This section summarized the the different findings concerning the Ni / GeSn system during the recent years in terms of the phase sequence, the morphological and electrical evolution during thermal annealing.

Ni / GeSn phase sequence

No data about the Ni-Ge-Sn ternary system is available in literature. However, the Ni / Ge system can be taken as reference. Indeed, the GeSn compound is obtained thanks to a small incorporation of Sn, (for this thesis the highest amount was 15 at.%), that can substitute certain Ge atomic sites. Then, the phase that could grow during the solid-state reaction are expected to follow the Ni / Ge phase diagram. The occurrence of Sn segregation or Ni x Sn y compounds formation can, nevertheless, not be dismissed, which supports the fundamental value of knowing and understanding the phase sequence. The Ni-Ge binary phase diagram is thus shown in Fig. 3.2. The binary diagram is rather complex, due to the existence of several equilibrium and metastable phases that are summarized in Table 3.1 [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF][START_REF] Peter | Phase analysis and thermal stability of thin films synthesized via solid state reaction of Ni with Si 1x Ge x substrate[END_REF]. These kinds of diagrams for bulk reactions represent a good base to foresee the phases that can be obtained when analyzing the reactions of thin metal films with a semiconductor. However, precautions should be taken in account. In some cases, the phase sequence might not agree with what is expected from bulk reactions. Indeed, sequential appearance of phases, absence of certain phases or very rapid kinetics of formation can occur [START_REF] Gas | Formation of silicide thin films by solid state reaction[END_REF]. In addition, the experimental process conditions (temperature, atmosphere, etc.) can impact and modify these reactions, as expressed in chapter I, section 1.2.

The Ni / Ge phase sequence during the solid-state reaction was studied by various groups on different types of Ge substrates (amorphous, (0 0 1) oriented, (1 1 1) oriented, etc.) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF][START_REF] Peter | Phase analysis and thermal stability of thin films synthesized via solid state reaction of Ni with Si 1x Ge x substrate[END_REF][START_REF] Nemouchi | Réactivité de films nanométriques de Nickel sur substrats Silicium-Germanium[END_REF][START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF][START_REF] Nemouchi | Dopant effect on NiGe texture during nickel germanide growth[END_REF][START_REF] Jablonka | Formation of nickel germanides from Ni layers with thickness below 10 nm[END_REF]. In general, the reaction started with the formation of a Ni-rich germanide, i.e. Ni 5 Ge 3 . Finally, when the Ni-rich phase was consumed, the growth of the NiGe phase (which was the last phase appearing) occurred. However, the identification of the Ni-rich phase has been controversial. First studies indexed the Ni-rich phase as the Ni 2 Ge or Ni 3 Ge 2 phase [START_REF] Hsieh | Partial epitaxial growth of Ni 2 Ge and NiGe on Ge (1 1 1)[END_REF][START_REF] Jin | The interfacial reaction of Ni with (1 1 1) Ge, (1 0 0) Si 0.75 Ge 0.25 and (1 0 0) Si at 400[END_REF], while certain analyses suggested the presence of the monoclinic ε'-Ni 5 Ge 3 [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Nemouchi | Simultaneous growth of Ni 5 Ge 3 and NiGe by reaction of Ni film with Ge[END_REF]. Finally, De Schutter et al. were able to identify the nature and the texture of the Ni-rich phase using X-ray pole figures and transmission electron microscopy (TEM). These techniques were used to unambiguously identify the transient phase that grew first. This phase corresponded to the ε-Ni 5 Ge 3 phase, a nonstoichiometric Ni-rich germanide with an hexagonal crystal structure that can exist for Ge concentrations between 34-48% [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. The results obtained by this group indicated that the incorrect identification of this phase, from past studies, could be attributed to the limited information that can be acquired from the standard θ-2θ geometry in case of textured phases. Indeed, as explained in chapter II, section 2.1.1, only planes that are oriented parallel to the surface can be measured. Hence, texture measurements combined with compositional depth profiles (to obtain stoichiometry information) can provide valuable information to properly identify the crystalline phases. This is the perfect example of the importance of both the in-plane pole figure characterization technique and the proper identification of the phase appearing during the solid-state reaction.

For the case of the Ni / GeSn system, a similar phase sequence was expected. Indeed, it is widely accepted that the phase sequence initiates with the appearance of a Ni-rich phase. The reaction occurs mainly thanks to Ni diffusion [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Jin | The interfacial reaction of Ni with (1 1 1) Ge, (1 0 0) Si 0.75 Ge 0.25 and (1 0 0) Si at 400[END_REF][START_REF] Patterson | Kinetics of Ni/a-Ge bilayer reactions[END_REF][START_REF] Comrie | Determination of the dominant diffusing species during nickel and palladium germanide formation[END_REF]. Then, the consumption of the Ni-rich phase allows the growth of the less resistive Ni(Ge 1-x Sn x ) orthorhombic phase. Demeleumester et al. and Vincent et al. followed the Ni / GeSn phase sequence by real-time XRD. Demeulemester et al. identified the first phase appearing as the Ni 5 Ge 3 phase, which was soon followed by the growth of the NiGe phase. Sn was observed to be soluble in both phases. The two Ni / Ge intermetallic phases grew simultaneously.

Ni, Ni 5 Ge 3 , and NiGe diffraction peaks were present since 290 up to 310 . At these temperatures, the Ni peak disappeared [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. Vincent et al. highlighted the formation of Ni(GeSn) starting from 300 , while the Ni-rich phase was not indexed [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF].

Concerning the rapid thermal annealing (RTA) technique, different studies identified the Ni-rich phase as the Ni 5 (GeSn) 3 phase as well. Meanwhile, the Ni(GeSn) phase was latter obtained [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Schulte-Braucks | Schottky barrier tuning via dopant segregation in NiGeSn-GeSn contacts[END_REF]. The composition and the crystallographic nature of the last phase appearing, the orthorhombic Ni-monostanogermanide phase, has been unambiguosly identified. However concerning the Ni-rich phase, different crystalline structures could match the crystalline peaks observed by XRD in the studies mentioned previously. When observing the Ni / Ge binary system (see Fig. 3.2 and Table 3.1), three Ni-rich phases were reported to be thermodynamically stable at room temperature β-Ni 3 Ge (cubic), Ni 2 Ge (orthorhombic), ε'-Ni 5 Ge 3 (monoclinic). Meanwhile, five other phases were only stable at elevated temperatures γ-Ni 3 Ge (crystal structure unknown), Ni 5 Ge 2 (hexagonal), ε-Ni 5 Ge 3 (hexagonal), Ni 19 Ge 12 (othorhombic) and Ni 3 Ge 2 (hexagonal) [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF][START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF][START_REF] Larsson | An electron diffraction study of modulated Ni 1+x Ge B8 type phases[END_REF]. Wirths et al. indexed the Ni / GeSn rich phase as the monoclinic Ni 5 (GeSn) 3 phase. However pole figures analyses made on Ni / (SiGeSn) system showed that the Ni 5 (SiGeSn) 3 phase might be hexagonal with in-plane domains rotated by 90° [START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Schulte-Braucks et al. also indexed the Ni-rich phase as the monoclinic Ni 5 (GeSn) 3 , however, the reflection, in XRD patterns, could also be attributed to the following Ni-rich phases the Ni 2 (GeSn) or Ni 3 (GeSn) 2 . Aditionally, fast Fourier transform (FFT) analysis in the cross-section TEM (XTEM) showed an hexagonal pattern similar to that reported for the Ni 5 (SiGe) 3 phase [START_REF] Schulte-Braucks | Schottky barrier tuning via dopant segregation in NiGeSn-GeSn contacts[END_REF][START_REF] Liu | Ultrathin highly uniform Ni(Al) germanosilicide layer with modulated B8 type Ni 5 (SiGe) 3 phase formed on strained Si 1x Ge x layers[END_REF]. Due to the diversity of the possible phases, a full crystallographic analysis of the Ni-rich nature was mandatory. This analysis was assessed by our own and is discussed hereafter in section 3.5.

Ni / GeSn systems are highly metastable due to the low solid solubility of Sn in the Ge lattice, i.e. less than 1% at the thermodynamic equilibrium, and the large difference in lattice parameters (around 14%). Sn segregation can easily occur if the annealing temperatures exceed the Ge-Sn eutectic point (231 ), the Sn melting point (232 ) or the GeSn epitaxy temperature (which varies with Sn content). Currently, it is not entirely known how Sn incorporates in the Ni / GeSn phases. Sn might (i) be soluble in the phase, (ii) segregate, and/or (iii) form Ni-Sn alloys [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. Demeulemeester et al. and Nishimura et al. showed that substitutional Sn was incorporated in the Ni(GeSn) phase while segregation can occur with the increase of the annealing temperatures [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF]. Meanwhile, Liu et al. related the incidence of certain XRD peaks to the metastable Ni-monostannide (NiSn) phase, when analyzing the Ni / GeSn solid-state reaction [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF]. No matter how Sn incorporates in the intermetallics, we believe that the addition of this element could have an impact on various aspects of the phase sequence during the solid-state reaction and the State of the art: Ni / GeSn system Ni-stanogermanides crystalline, morphological and electrical properties. In order to verify these aspects, a comprehensive analysis of the behavior of Sn in Ni / GeSn intermetallics with different Sn concentrations was fully assessed and is presented in section 3.6.

Ni / GeSn morphological properties and stability

Ni / GeSn intermetallics, formed by a solid-state reaction, consist of polycrystalline layers [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Schulte-Braucks | Schottky barrier tuning via dopant segregation in NiGeSn-GeSn contacts[END_REF], in which the necessary amount of Ni to produce a certain quantity of the Ni(GeSn) phase can be known. The solid-state reaction approximately uses a proportion of Ni : GeSn of 1 : 2.25 for the GeSn consumption and a proportion of Ni : Ni(GeSn) of 1 : 2.5 for the Ni(GeSn) formation [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Han | High-mobility germanium-tin (GeSn) P-channel MOSFETs featuring metallic source/drain and sub-370 process modules[END_REF][START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF]. That means that for 1 nm of Ni, 2.5 nm of Ni(GeSn) can be formed, while consuming 2.25 nm of GeSn.

In addition, the integrity of these layers can be analyzed by means of scanning electron microscopy (SEM), transmission electron microscopy (TEM) or atomic force microscopy (AFM), to verify either surface or cross-section evolution when the process conditions were changed, as observed in Fig. In general, the layers can remain smooth and continuous up to a threshold temperature which has been commonly related to the start of Sn segregation, agglomeration or degradation processes [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Li | Electrical characteristics of Ni ohmic contact on n-type GeSn[END_REF][START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Zhang | Formation and characterization of Ni/Al ohmic contact on n + -type GeSn[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF]. At these temperatures, degradation was identified thanks to a lost of continuity in the layers, compounds agglomeration, together with island in the surface that were related to Sn accumulation. Indeed, Nishimura et al. and Wirths et al. both observed the influence of Sn content and thermal annealing on the surface morphological evolution and degradation (see Fig. 3.3) [START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. The evolution of the morphological properties might depend on different factors annealing conditions, Sn content, crystallites size, crystalline orientation and so on. A systematic evaluation of Chapter III. Study of the Ni / GeSn system morphological integrity was then performed and is addressed in section 3.3. The advantage of this analysis is that a correlation on the impact of the morphological evolution and the electrical properties can be realized.

Ni / GeSn electrical properties and stability

Electrical measurements on semiconductor devices are routinely performed for the analysis of a wide range of properties. Some of the routine methods of electrical characterization are related to the measurements of the electrical resistivity (conductivity), carrier mobility and carrier types. For very thin semiconductor layers, sheet resistance (R sh ) measurements can be used to assess electrical properties as discussed in chapter II, section 2.3. Commonly, for the case of Ni / GeSn layers, sheet resistance was found to decrease down to a minimum value, with increasing annealing temperatures, followed by an increase of resistance values as the annealing temperatures further increased, as observed in Fig. 3.4 [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Li | Electrical characteristics of Ni ohmic contact on n-type GeSn[END_REF][START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Zhang | Formation and characterization of Ni/Al ohmic contact on n + -type GeSn[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF]. The threshold temperatures, for the reduction in sheet resistance values, was related to the growth of the less resistive Ni(GeSn) phase. Additionally, the range of temperatures in which these values remained constant were related to the thermal stability of the layers. Meanwhile, the further increase of R sh values at higher annealing temperatures was related to the agglomeration or degradation of the film or the interfaces, increase of defects or Sn segregation. On the other hand, the minimum R sh values obtained and the range of thermal stability was observed to vary with different process conditions, as illustrated in Fig. 3.5.

Wirths et al. observed that the Ni / GeSn electrical stability varied with the Sn content in the GeSn layer. The higher the Sn addition, the smaller the thermal stability range, while the minimum R sh values remained comparable [START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Liu et al. and Wang et al. both studied Ni / GeSn samples with 5.3 at.% Sn content, while the range of electrical thermal stabilities were substantially different (see Fig. 3.5) [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF]. Moreover, the strain variation due to GeSn growth conditions also changed the thermal range of stability [START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF]. Ensuring high electrical performance and stability is of high importance for the metallic contacts used in electronic or opto-electronic devices. The effects of annealing temperature during the solid-state reaction, the Sn content, the morphological evolution and other process conditions, on the electrical properties were assessed and presented in section 3.3.

In summary, the Ni / GeSn phase sequence during the solid-state reaction starts with the growth of a Ni-rich phase. The identification of the crystalline nature of this phase is still ambiguous. A proper identification is thus mandatory. The Ni-rich phase is consumed to the benefit of the orthorhombic Ni(GeSn) phase, which is the less resistive phase of the system. These kinds of layers present a low thermal stability. The annealing temperature corresponds then to an important process condition to adjust. The goal is to ensure proper morphological and electrical properties of the Ni / GeSn intermetallics.

Different characterization techniques were deployed to verify various aspects concerning the morphological evolution with temperature (see Fig. 3.3). The annealing temperatures will have an impact on the electrical properties as well (see Fig. 3.4), while other process conditions (Sn content, layers strain, etc.) might also modify the electric transport.

For our studies, specific process conditions in order to obtain Ni / GeSn samples were realized and are described in the next section 3.2. Due to the importance of a full characterization of the metallic contacts and the impact that the specific process conditions can have on the properties evolution, the phase sequence was assessed to relate the layer microstructure to their properties. In addition, the evaluation of the morphological and the electrical properties with annealing temperatures and the impact of changing other process conditions was also necessary to determine. Knowing the range of thermal stability, the minimum R sh values that can be obtained, the temperature thresholds for layer agglomeration and Sn segregation, was necessary.

A comprehensive analysis was then realized with combined characterization techniques and is presented hereinafter through different sections.

As stated in chapter I, section 1.2, different studies have addressed the formation of Ni-based intermetallics on GeSn, in which the metallization parameters varied from one group to the other, from the surface pre-clean to the annealing parameters, as summarized in Table 1.2. This section gives an overview of the parameters used to elaborate the samples analyzed in this PhD thesis, from the GeSn growth to the metallization procedure.

Growth and characterization of GeSn layers

In order to realize a systematic and comprehensive analysis of the Ni / GeSn system, samples were prepared as follows. Pure Ge layers and GeSn layers with 6, 10 and 15 at.% of Sn were epitaxially grown on Ge-buffered Si (1 0 0) substrates in a 200 mm reduced pressure chemical vapor deposition (RP-CVD) tool [START_REF] Aubin | Very low temperature epitaxy of heavily in situ phosphorous doped Ge layers and high Sn content GeSn layers[END_REF]. Although the lattice mismatch between Sn and Ge is large (14%) and the solubility of Sn in Ge is low at thermodynamic equilibrium (less than 1%), high quality stacks were obtained, as explained in chapter I, section 1.1.2. The Sn content was measured from ω-2θ XRD scans around the (4 0 0) Si reflection, while surface morphology was evaluated by AFM images (see Fig. The intense, slightly asymmetric peak positioned around 33.0-33.5°in each ω-2θ spectrum in Fig. 3.6 corresponded to the 2.5 µm thick Ge strain relaxed buffer (SRB), which was slightly shifted towards the Si substrate peak [START_REF] Aubin | Very low temperature epitaxy of heavily in situ phosphorous doped Ge layers and high Sn content GeSn layers[END_REF][START_REF] Aubin | Very low temperature epitaxy of Ge and Ge-rich SiGe alloys with Ge 2 H 6 in a reduced pressurechemical vapour deposition tool[END_REF]. The peaks related to the Ge 1-x Sn x layer were shifted to lower angles depending on the Sn content. The higher Sn composition corresponded to the lower angle peak. The larger atomic radii (a r ) of substitutional Sn (a r = 140 pm) in the Ge lattice (a r = 122 pm) generated different levels of strain. A higher Sn content induced a higher strain, consequently shifting the peaks to lower ω values. Surface morphology observed in AFM images in Fig. 3.6 exhibited a cross-hatch pattern along the 1 1 0 crystallographic direction. This pattern resulted from the propagation of the 60°threading arms of misfit dislocations on (1 1 1) planes [START_REF] Aubin | Very low temperature epitaxy of heavily in situ phosphorous doped Ge layers and high Sn content GeSn layers[END_REF][START_REF] Aubin | Very low temperature epitaxy of Ge and Ge-rich SiGe alloys with Ge 2 H 6 in a reduced pressurechemical vapour deposition tool[END_REF][START_REF] Gaspare | Epi-cleaning of Ge/GeSn heterostructures[END_REF]. Thus, this surface morphology directly originated from the epitaxy of GeSn layers on the Ge (SRB). No island formation due to Sn segregation was observed while root-mean-square (RMS) roughness values were low (see Fig. 3.6), which demonstrated the high quality and uniformity of the layers. For the Ge 0.85 Sn 0.15 alloy, local defects were observed on the surface (see Fig. 3.6). The latter corresponded to inverted pyramids with a square base. Such pits might be the emerging points of threading dislocations, coming from the Ge SRB underneath that propagated inside the GeSn layer. On the other hand, the trace of the dislocations movement was observed in the different GeSn samples, verified by marked tranches or lines traversing the surface.

Ni / GeSn samples preparation

The thickness of the GeSn layers (60 nm) was low enough for them to be fully compressively strained on the Ge SRB underneath. In order to remove particles and native oxides from the surface, samples were dipped for 1 min in diluted HF (1% in water) followed by a rinse in deionized water and dried with a N 2 gun. Then, 10 nm of Ni were deposited at room temperature followed by 7 nm of TiN capping layer deposited at 100 by physical vapor deposition (PVD-magnetron sputtering), in a 200 mm equipment. Giving that no in-situ pre-clean was used, the air break between the surface preparation and the sample transfer under vacuum for metallization was reduced to the minimum (less than 10 min). On the other hand, no air break occurred between the Ni and TiN deposition samples were maintained at a pressure as low as 10 -7 Torr during the transfer between the two deposition chambers. The TiN capping layer protected the layers of interest from atmospheric contamination later on. This way the solid-state reaction was not impacted by oxygen or carbon contamination. A summary of the process step is schematized in Fig. 3.7. Next sections will present our results concerning the Ni / GeSn phase sequence during the solid-state reaction, the texture of the intermetallics, the impact of Sn content on the evolution of morphological and electrical properties with annealing temperatures.

3.3 General characterization of the Ni / Ge 0.90 Sn 0.10 system

As stated in section 3.2, pure Ge layers and GeSn layers with 6, 10 and 15 at.% of Sn were epitaxially grown on Ge-buffered Si (1 0 0) and metallized with Ni and TiN capping layers under the same conditions. This section, however, focuses on the phase sequence, the morphological and the electrical evolution of the Ni / Ge 0.90 Sn 0.10 system only, as a reference example. The results given in this section will be further compared to samples with different Sn contents in section 3.6. At low temperatures, characteristic peaks of TiN (2 0 0) at 42.4°, Ni (2 0 0) at 52.1°, Si (4 0 0) and Ge (4 0 0) around 67.0°were observed. In addition, an unexpected peak was visible at 31.6°and another one was distinguishable at around 66.5°. These latter were more visible in Fig. 3.9 depicting a focus on the low temperature range (i.e. 50 to 180 ). These peaks have not been reported at such temperature in the previous in-situ XRD studies performed on Ni / Ge 1-x Sn x systems under similar conditions [START_REF] Vincent | Characterization of GeSn materials for future Ge pMOSFETs source/drain stressors[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. The presence of these peaks was attributed to the process conditions involved in this study. Indeed, TiN was deposited at 100 . This temperature was high enough to allow the reaction between Ni and Ge 0.90 Sn 0.10 to start and generate the growth of a Ni-rich phase since such phase was reported to form on as-deposited samples during the deposition of Ni thin films on Ge [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. Looking at previous investigation on Ni / Ge(Sn) solid-state reaction, this Ni-rich phase could be attributed to Ni 2 (GeSn), Ni 5 (GeSn) 3 , Ni 3 (GeSn) 2 or Ni 19 (GeSn) 12 phases [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Schulte-Braucks | Process modules for GeSn nanoelectronics with high Sn-contents[END_REF][START_REF] Nemouchi | Dopant effect on NiGe texture during nickel germanide growth[END_REF]. The exact nature of the Ni-rich phase will be discussed hereinafter, (see section 3.5). At 120 , the Ni layer was fully consumed. Between 125 and 150 , the peaks corresponding to the Ni-rich phase underwent a slight shift towards higher 2θ angles. The peak position importantly evolved once the Ni layer was consumed. After this shift, the peaks were located at 32.2°and around 67.0°. The intensity of both peaks increased with temperature. The origin of the shift towards higher 2θ angles and the increase of the Ni-General characterization of the Ni / Ge 0.90 Sn 0.10 system rich phase peak's intensity, while Ni was totally consumed, will be discussed hereinafter in section 3.5.

Around 245 , diffraction peaks related to the Ni-rich phase vanished and new diffraction peaks appeared at 34.2°, 45.0°, 53.9°and 87.7°. These new peaks were related to the growth of the mono-stanogermanide Ni(Ge 0.90 Sn 0.10 ) phase and related to the following out-of-plane orientations (1 1 1)/(1 0 2), (2 1 1)/(2 0 2), (0 2 0)/(0 1 3) and (0 1 5). A double indexation has to be done for certain peaks due to the close 2θ values of the poles. The preferred orientation, i.e. the one for which diffraction peaks were the most intense, was the (0 2 0)/(0 1 3). These privileged out-of-plane orientations were not obvious taking into account the relative peak intensities from powder diffraction files from ICDD data base (the PDF chart can be consulted in Appendix A). The mono-stanogermanide phase was stable up to 600 .

In order to complement the in-situ XRD results, Ni / GeSn samples, obtained under the same process conditions, were ex-situ annealed for 30 s at temperatures ranging from 150 up to 550 in a rapid thermal annealing (RTA) furnace under N 2 environment.

Then, θ-2θ scans over a large range (20-90°) using a Cu source were performed in a X'Pert Pro PANalytical X-ray diffractometer. Results are presented in Fig. 3.10. From the as-deposited state and up to 200 annealing, TiN (2 0 0) and Ni (2 0 0) peaks were observed (at 2θ = 42.0°and 2θ = 52.0°respectively). Two peaks around 2θ = 31.4°and 66.4°were also visible. They were attributed to the Ni-rich phase as well as for the in-situ XRD case. At 250 , the Ni layer was fully consumed and the Ni-rich peak underwent a slight shift towards higher 2θ angles. After this shift, the peaks were located

Chapter III. Study of the Ni / GeSn system at around 32.0°and 68.0°. This Ni-rich phase remained stable up to 300 . At 350 , the Ni(GeSn) phase was then obtained, with various out-of-plane orientations. Then, at the early stage of the reaction (350-400 ), other peaks at 2θ = 30.6°and 32.0°were observed. These were attributed to the β-Sn phase which is due to Sn segregation. Peaks related to Sn segregation where not observed in the in-situ XRD results, the reasons of this behavior will be discussed in section 3.6.

The summary of the phase sequence between a Ni thin film and Ge 0.90 Sn 0.10 layer obtained by in-situ and ex-situ XRD is presented in Fig. 3.11. In general, the phase sequence was comparable for in-situ and ex-situ XRD samples a Ni-rich phase was formed first. Then, the Ni-rich phase disappeared for the growth of the mono-stanogermanide phase Ni(GeSn). This phase appeared to be stable up to 600 . Differences were however observed concerning the threshold temperatures for phase transition. Indeed, for rapid thermal annealed samples, the phase sequence seemed to be delayed up to higher temperatures, as illustrated in Fig. 3.11. Thermal budgets chosen during the solid-state reaction might have affected the sequence of the phase formation [START_REF] D'heurle | The kinetics of reactive phase formation: silicides[END_REF]. In solids, diffusion is an important parameter to achieve mass transport, which may be defined as the migration of an atomic or molecular species within a given matrix under the influence of a concentration gradient [START_REF] Ohring | Materials Science of Thin Films: Deposition and Structure[END_REF]. Additionally, thin-films silicides and germanides reaction kinetics are expected to be limited by the diffusion step. That means that the diffusion step is the slowest step of the reaction and the one that can limit the phase formation [START_REF] Nemouchi | Réactivité de films nanométriques de Nickel sur substrats Silicium-Germanium[END_REF]. Diffusion is a time-dependent process, in a macroscopic sense, the quantity of an element that is transported within another as a function of time. Ad-ditionally, temperature has a great influence on the coefficients and diffusion rates [START_REF] Callister | Materials Science and Engineering: An Introduction[END_REF]. Between in-situ and ex-situ analysis the difference in thermal budgets is obvious. For insitu XRD analysis, one sample was heated from 50 up to 600 following temperature steps of 5 . Once the targeted temperature was reached, the sample remained at this temperature for 7 min while the XRD pattern was acquired. Meanwhile, for RTA treated samples, each annealing temperature corresponds to a sample that was heated up to the targeted temperature for 30 s and immediately cooled down to acquired the ex-situ XRD pattern. The 7 min thermalisation promoted during the in-situ measurement might allow mass-diffusion to occur which was beneficial for the phase growth at lower temperatures. Our observation thus highlighted the importance of the kinetics and thermal budget (time vs. temperature) for the Ni / Ge 1-x Sn x systems.

3.3.2

The Ni / Ge 0.90 Sn 0.10 morphological evolution Ni / GeSn intermetallics are well suited to contact GeSn-based devices, in terms of phase formation and electrical performance [START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Nevertheless, they exhibit a poor thermal stability linked to the precipitation of Sn atoms and the agglomeration of the Ni(GeSn) layer, as presented in section 3.1.2. These phenomena were observed to be Sn content dependent and strong for GeSn layers with high Sn concentrations [START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Atomic force microscopy images, in tapping mode, with a Bruker FastScan equipment were thus acquired in order to verify surface stability when increasing the annealing temperatures. Fig. 3.12 shows the AFM images obtained for samples annealed at various temperatures by RTA for 30 s in N 2 atmosphere for the Ni / Ge 0.90 Sn 0.10 system. As mentioned in chapter II, section 2.2.1, this technique is well adapted to characterize the surface morphology of the samples. Additionally, understand and rely the results obtained with other complementary techniques (XRD and R sh ) is quite easy. Then, relating morphological changes of the surface to the phase sequence and the electrical evolution can be possible. At low temperature (i.e. for T < 300 ), samples exhibited a cross-hatch pattern along the 1 1 0 crystallographic direction, as observed in section 3.2.1 when analyzing the epitaxial GeSn layers. This surface morphology directly originated from the epitaxy of Ge/GeSn layers. This means that the Ni and TiN layers were extremely smooth and their deposition did not modify the original surface morphology. The cross-hatch pattern was preserved up to 300 . Above this temperature, the Ni-rich phase growth was observed by XRD analysis (see Fig. 3.10). This means that the formation of the Ni-rich phase did not modify the original morphology either.

At 350 , the surface morphology changed. These changes were linked to the formation of the Ni-monostanogermanide phase (verify by XRD analysis, see Fig. 3.10) and the subsequent layers degradation. The degradation of the layer continued with temperature and finally, at 550 , dots were observed in the surface. These dots can be related to Sn segregation (verified by XRD analysis in Fig. 3.10).

For the case of germanides/stanogermanides, degradation can start due to grain boundary grooving up to the full agglomeration of the films. As illustrated in Fig 3 .13, the thin film layer tries to minimize the interface energies and therefore reduce the angles of contact between the film and the substrate. A groove can be formed at the grain boundaries then. When this etching reaches the substrate, the continuity of the film is broken and a hole is formed, which is define as agglomeration. During the grain boundary grooving, the film can remain continuous while the grain boundaries are curved. This step is defined then as the prior step to agglomeration in which a discontinuous film is obtained [START_REF] Bourjot | Germanosiliciuration a base de Ni et d'alliage Ni 1-x Pt x pour le p-MOS 14 nm FDSOI[END_REF].

After 350 , a strong increase of the surface roughness occurred which gradually increased with temperature, see Fig. 3.14. In order to verify the layers continuity, TEM analysis were realized. The results are presented in Fig. 3.15.

For this sample and in the as-deposited state, Ge, GeSn, Ni-rich, Ni and TiN layers were well identified. The intermetallic layer of about 12 ± 1 nm was related to the Ni-rich phase and was observed to be continuous and uniform. At 350 , a polycrystalline layer of 30 ± 1 nm was related to the Ni(GeSn) phase. The presence of both the Ni-rich phase and the Ni(Ge 0.90 Sn 0.10 ) at the temperatures mentioned was verified by XRD analysis, see section 3.3.1. The start of the grain boundary grooving (previous step to the agglomeration) was evidenced for the sample annealed at 350 , which confirmed the start of the surface degradation evidenced in Fig. 3.12. At 400 , the intermetallic layer completely lost its continuity due to the agglomeration process.

The changes in the surface morphology in Fig. 3.12 and the increase of the RMS roughness values in Fig. 3.14 were then linked to the layers agglomeration. A complete lost of continuity of the layers was verified for the Ni / Ge 0.90 Sn 0.10 system by TEM analysis in Fig. 3.15. The degradation of the layers was then due to the amplification of both phenomena the layer agglomeration and the Sn segregation (dots observed in Fig. 3.12 at 500 ). The attractiveness of Ni / GeSn intermetallics for contacts relies on the low values of contact (R c ) and sheet (R sh ) resistances that can be achieved [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. However, as explained in sections 3.1.2 and 3.1.3, and verified previously in section 3.3.2, Ni / GeSn intermetallics exhibit a poor thermal stability. Agglomeration and Sn segregation can degrade the electrical properties compromising the device performances. The R sh evolution with annealing temperature was acquired with a four-point probe meter. The objective was to verify the impact of the phase sequence and the morphological evolution with temperature on the electrical properties, verify the process window at which values can remain low, stable and in an acceptable limit for integration. Fig. 3.16 shows the evolution of the sheet resistance, for the Ni / Ge 0.90 Sn 0.10 system after various RTA from 150 up to 550 . Beyond 350 , a strong increase of the R sh values was observed. This behavior was directly linked to the agglomeration of the Ni(GeSn) layer and to the subsequent Sn segregation. The lost of continuity in the surface layer and the increase in root-meansquare roughness values was verified by AFM (Fig. 3.12 and Fig. 3.14). TEM analyses, in Fig. 3.15, confirmed the agglomeration process and the lost of the layer continuity at 400 . Then, the morphological degradation had a direct impact on the electrical properties stability. The lost of continuity of the layers had a detrimental impact on R sh values.

To summarize this section, an overview of the Ni / Ge 0.90 Sn 0.10 system evolution was presented as a reference example of the Ni / GeSn systems evolution. The phase sequence during the solid state-reaction of a thin Ni film (10 nm) and Ge 0.90 Sn 0.10 layers was analyzed. These results support previous studies from the literature mentioned in section 3.1.1. We evidenced a sequential growth of the layers, in which a Ni-rich phase was formed first. Later, the Ni-rich phase was consumed to the benefit of the Ni(GeSn), orthorhombic mono-stanogermanide phase, which is the less resistive phase (verified by the correlation of XRD and R sh measurements in Fig. 3.10 and 3.16). The nature of the Ni-rich phase is still ambiguous. The proper identification of this phase is then necessary and will be addressed in section 3.5. A variation of the threshold temperatures for phase formation was observed when comparing in-situ and ex-situ XRD patterns. These results highlighted the importance of the kinetics and thermal budget (time vs. temperature) for Ni / Ge 1-x Sn x systems.

During the solid-state reaction, the morphological and electrical properties evolved as well. R sh values were observed to decrease to a minimum value (see Fig. 3.16), in which the growth of the Ni(GeSn) phase was verified by XRD analysis. Coupled to this, the surface morphology changed. The common cross-hatch pattern that was the result of the Ge / GeSn epitaxial growth dissipated. In return, evidence of layers agglomeration was encountered by AFM images (see Fig. 3.12) and verified by TEM analysis (see Fig. 3.15). When annealing temperatures further increased, degradation issues aggravated due to Sn segregation. This degradation on morphological properties was directly related to the increase of R sh values, which demonstrated the additional degradation of the electrical properties.

This comprehensive analysis of Ni / GeSn intermetallics allowed us to lead our research towards the understanding of different main subjects that are discussed through different sections the effect of the surface pre-clean (section 3.4), the proper identification of the Ni-rich phase (section 3.5), the impact of the Sn content on the Ni / GeSn intermetallics (section 3.6 and 3.7) and the impact of changing different process conditions the amorphization of GeSn layer (section 4.1), addition of an alloy element to the Ni layer (section 4.2) and the change in annealing conditions (section 4.3), which will be addressed in the following sections and chapter.

Effect of surface preparation on the Ni / GeSn properties evolution during the solid-state reaction

After the GeSn epitaxial growth and before the metal deposition, a surface preparation corresponds to a necessary step. Whatever the targeted application, wet cleaning is indeed mandatory to remove particles, organic materials, metallic impurities and native oxides from the surface [START_REF] Raynal | GeSn surface preparation by wet cleaning and in-situ plasma treatments prior to metallization[END_REF][START_REF] Aspnes | Chemical etching and cleaning procedures for Si, Ge, and some III-V compound semiconductors[END_REF]. Literature data on GeSn surface preparation prior to the metallization step is currently scarce [START_REF] Gupta | GeSn Technology: Extending the Ge Electronics Roadmap[END_REF][START_REF] Gupta | Atomic layer deposition of Al 2 O 3 on germanium-tin (GeSn) and impact of wet chemical surface pre-treatment[END_REF]. Based on the existing knowledge on pure germanium, the surfaces are unstable due to the high reactivity upon air and oxygen exposure which can demand the use of high dissolution in oxidant solutions (GeO 2 dissolution in water). Ex-situ wet cleanings will likely suffer from a fast native oxide re-growth on the GeSn surface. In-situ plasma treatments, in a chamber connected to the metal deposition chamber itself, i.e. without any air break, can prevent that issue. From a previous analysis concerning the impact of wet and plasma surface treatmenst realized on GeSn layers it was evidenced that [START_REF] Raynal | GeSn surface preparation by wet cleaning and in-situ plasma treatments prior to metallization[END_REF]:

-The oxidation kinetics of GeSn surface is extremely fast upon air exposure, Sn was more sensitive to oxidation than Ge.

-HF surface treatment was more efficient for removing Sn oxides than Ge oxides.

-The rinse in deionized water partially re-oxidized the GeSn layer.

-Helium (He) or argon (Ar) plasma treatments efficiently removed a fraction of the GeSn native oxide without degrading the surface morphology.

We decided then, to verify the impact of the surface preparation treatment on the phase sequence, the morphological and the electrical evolution during the Ni / GeSn solid-state reaction. For these aims, compressively-strained, 60 nm thick layers of Ge 0.90 Sn 0.10 , obtained under the same conditions explained in previous sections (section 3.2.1) were used. One part of the wafer was kept without surface preparation as reference. Wet treatments were performed in a cleanroom bench. Diluted HF immersion (i.e. a 1 min dip in HF with a 1% molar concentration in water) followed by a deionized water (DIW) rinse were realized. Then, samples were dried with a N 2 gun. Meanwhile, He or Ar plasma treatments were otherwise performed in a pre-clean chamber connected to a 300 mm automatic Endura platform from Applied Materials. Then Ni and TiN layers deposition were performed successively by RF-PVD, both at room temperature in the same 300 mm tool.

It has to be highlighted that certain changes, compared to the samples analyzed in section 3.3, were yet applied. First, the deposition at room temperature of the TiN layer for this case and the fact that the deposition (both of Ni and TiN layers) was realized in a 300 mm equipment.

Effect of surface preparation on the Ni / GeSn phase sequence

In order to verify the impact on the phase sequence, in-situ XRD patterns were acquired respecting the same parameters explained in chapter II, section 2.1.1. The Ni / Ge 0.90 Sn 0.10 , with beforehand HF surface treatment, phase sequence was followed by in-situ XRD and is shown in Fig. 3.17 as a contour map. The y-axis shows the temperature and the x-axis, the diffraction angles. Characteristic peaks of TiN (1 1 1) at 36.4°, Ni (1 1 1) at 44.9°, Si (4 0 0) and Ge/GeSn (4 0 0) around 67°were observed. Around 85 , two diffraction peaks appeared at 46.0°and 85.6°. These peaks were due to the growth of the Ni-rich phase. These peaks slightly shifted towards higher 2θ angles above 120 (after total Ni consumption) as observed for the Ni / Ge 0.90 Sn 0.10 system in Fig. 3.8. This shift will be described further in section 3.5.

At 210 (marked by a dashed line in Fig. 3.17), the diffraction peaks linked to the Ni-rich phase dissipated and new diffraction peaks appeared at 42.3°, 45.2°and 53.4°. They were due to the formation of the Ni mono-stanogermanide phase (Ni(GeSn)). This phase remained stable up to 670 . At this temperature the dissipation of the diffraction peaks was related to a phase destabilization due to reaching temperatures close to the NiGe-Ge eutectic point [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]. Even if the eutectic point between NiGe and Ge was reported to be at around 760 [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]. The calibration of the furnace used for the in-situ XRD analyses showed that, in this range of temperature, a gap of 50 up to 100 between the temperature displayed by the furnace and the actual temperature was encountered. Indeed, Ni / Ge samples were analyzed under the same conditions, (the full in-situ XRD pattern can be observed in Fig. 4.13), in which destabilization of the NiGe phase was encountered at the same range of temperature. Then, Sn addition might not importantly modify the Ni-Ge eutectic point.

In order to study the potential impact of the surface treatment, similar in-situ XRD analyses were made on Ni / GeSn systems without any treatment and with He or Ar plasma treatments. The results obtained are displayed in Fig. 3.18. In general, the same phase sequence was observed from a straightforward comparison of in-situ XRD patterns. The phase sequence for the samples went as follows. (i) A Ni-rich phase was formed first, (ii) then, the Ni-rich phase disappeared to the benefit of the mono-stanogermanide phase (i.e. Ni(GeSn)). This phase was stable until reaching temperatures close to 660-670 . The dissipation of the diffraction peaks was related to a phase destabilization linked to the reach of the NiGe-Ge eutectic point. Differences were however observed concerning transition temperatures for the phase growth as summarized in Table 3.2.

The Ni-rich phase was obtained at around 115-120 for the reference sample. Meanwhile for Ar-treated and HF-treated samples, the Ni-rich phase was observed since the as-deposited state and at 85 , respectively. Finally, for the He-treated samples the threshold temperature did not vary much [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF]. For the non-treated sample, native oxides are expected to be present in the surface. For Ar and He pre-clean treatment, a proper elimination of oxides is expected, while for HF-treated samples the oxides regrowth is quite possible due to the unavoidable air-break prior to the metallization. In order to verify the efficiency of Ar and He plasma treatments, X-ray photoelectron spectroscopy (XPS) measurements were performed on a Thermo Electron Theta 300 spectrometer on Ge 0.94 Sn 0.06 samples. The results are expected to be indicative and comparable to what can be expected for Ge 0.90 Sn 0.10 samples. Quasi in-situ analyses were ensured with the use of a vacuum wafer carrier and transfer interface is directly plugged to the XPS via a vacuum transfer chamber (10 -7 mBar range). A N 2 pressure of 10 -2 mbars range was generate in the carrier during the transfer. XPS spectra was used for various collection angles of 23.7 ± 3.7°and 76.2 ± 3.7°, corresponding to photoelectrons escaping from the bulk surface and the near surface, approximately 10 and 1 nm respectively of depth depending in the energy regions observed. All the analyses were performed with a high resolution monochromatic Al Kα X-ray source (1488.6 eV photons) for a pass energy of 60 eV. Ge bulk , Ge GeO , Ge GeO 2 , Sn bulk and Sn Ox concentrations were extracted from Ge 2p 3/2 and Sn 3d 5/2 core level energy regions. The Ge 2p 3/2 level (1220 eV binding energy) was used due to the high surface sensitivity obtained at this low level of photoelectron kinetic energy. Spectral deconvolution was performed to extract the peak contributions in the acquired energy regions using a numerical fitting procedure. Line shapes were simulated with a combination of Lorentzian and Gaussian functions. Ge GeO peak was fit at a binding energy (BE) shift of +1.45 eV from Ge bulk , and Ge GeO 2 at a BE shift of +3.05 eV. BE shift between Sn bulk and Sn Ox was 1.45 eV after the plasma treatments and 1.83 eV for the substrate as-received. Fig. 3.19 presents the XPS spectra obtained for the reference as-received Ge 0.94 Sn 0.06 sample while Table 3.3 gives a summary of the oxides proportions measured for 23.7 ± 3.7°and 76.2 ± 3.7°angles collection corresponding to the bulk and the near surface responses from the photoelectrons of the different samples analyzed. Fig. 3.19 shows the energy area fit deconvolution for a Ge 0.94 Sn 0.06 substrate asreceived focusing on the Ge 2p and Sn 3d energy region. The GeSn sample as-received presented mainly the GeO 2 specie and the Ge metal signal was low. A broad SnO 2 signal was additionally observed. From the results displayed in Table 3.3, it can be observed that both Ar and He plasma significatively reduce the amount of Ge and Sn oxides. However, Ar plasma was more efficient for oxides elimination. For the reference (without pre-clean) and He treated samples native oxides are expected to be present prior to the metallization. For the HF treated sample, oxide regrowth is expected to occur due to the air-break, meanwhile the Ar treated samples had just a small fraction of oxidation. Surface preparation to remove oxides and contamination is indeed critical to obtain proper interfaces for contacting the metal layer [START_REF] Park | A novel approach to clean surface for high mobility channel materials with in-situ atomic hydrogen clean[END_REF][START_REF] Rodriguez | In situ cleaning/passivation of surfaces for contact technology on III-V materials[END_REF]. Additionally, the quality of the surface might not only alter devices performance but modify microstructural properties [START_REF] Aspnes | Chemical etching and cleaning procedures for Si, Ge, and some III-V compound semiconductors[END_REF]. The remaining native oxides (mostly for the reference and He plasma treated samples) or regrowth of the oxides (mostly for HF treated samples) might act as a diffusion barrier for atom diffusion. The Ni-rich phase growth can then be delayed. On the contrary, an oxide-free surface enhanced the Ni 5 (GeSn) 3 phase growth (since the as-deposited state) for the Ar treated system. The temperatures for the Ni(GeSn) phase growth remained comparable, values of 210 , 215 , 220 and 225 were observed for reference, HF treatment, Ar-treatment and He treatment samples, respectively. These variations can be due to a variation in the interface state and interface energy after the treatments, then the interfacial reaction rates might vary. However, the variation in threshold temperatures for the Ni(GeSn) phase growth are not as large as observed for the Ni-rich phase growth (variation of maximum 15 for the Ni(GeSn) growth compared to the maximum difference of around 65 for the Ni-rich growth). Then, once the native oxide barrier was overcome (and the Ni-rich phase was obtained), the diffusion mechanisms to growth the Ni(GeSn) phase appeared to be similar. In this way, only the growth of Ni-rich phase was importantly impacted.

On the other hand, one aspect to consider corresponds to the texture of the intermetallics, i.e. the preferential crystallographic orientations of the layer. Indeed, an influ-ence of texture on phase formation was evidenced when comparing the phase sequences for a thin Ni film on Si (0 0 1) and Si (1 1 1) substrates and when comparing the growth of Ni films on Ge (0 0 1) and Ge (1 1 1) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Gaudet | The thermally-induced reaction of thin Ni films with Si: Effect of the substrate orientation[END_REF]. For the Ni / Ge case, Ni 5 Ge 3 and NiGe appeared consecutively on Ge (1 1 1) while they grew simultaneously on α-Ge and Ge (0 0 1). Phase formation temperatures changed strongly with the nature of the substrate being the lowest on α-Ge and the highest on Ge (1 1 1) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF]. Furthermore, the phase stability, i.e. the maximum temperature up to which an intermediate phase is stable, before it reacts with the substrate to form a phase that is more thermodynamically stable, can vary with texture [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. The surface states prior to the metallization can then modify the preferential texture of the thin film layers which might modify the phase sequence.

The in-situ XRD patterns acquired for the reference sample and the Ar treated sample are compared in Fig. 3.20 to verify the impact on the variation of the Ni / GeSn intermetallic phases texture. From a straightforward comparison of both patterns a difference on the out-of-plane orientations, for both the Ni-rich and the Ni(GeSn) phases, was evidenced. For the reference sample (without pre-clean), four peaks related to the Ni-rich phase were observed at 31.6°, 46.5°, 67.9°and 85.2°. Later, three peaks of the Ni(GeSn) appeared at 42.3°, 44.9°and 53.6°and were related to the (2 1 0), (2 1 1)/(2 0 2) and (0 1 3)/(0 2 0) orientations, respectively. For the Ar treated sample, two peaks were also related to the Ni-rich phase and were observed around 46.5°and 85.6°. Meanwhile, three peaks were detected for the Ni(GeSn) phase at 42.4°, 45.1°and 53.6°, related to (2 1 0), (2 1 1)/(2 0 2) and (0 1 3)/(0 2 0) orientations, respectively. Then, a variation of the out-of-plane crystalline orientation for both Ni / GeSn intermetallics was confirmed.

The impact changes on the crystalline orientation, on both phases, could then impact Chapter III. Study of the Ni / GeSn system the stability and the threshold temperatures for phase growth. Nevertheless, in order to correlate the variations of the phase thermal stability to the crystalline variations with total certainty, complementary XRD characterization techniques (i.e. in-plane RSM, pole figures) should be performed. Indeed, as expressed in chapter II, section 2.1.1 , the limitations of θ-2θ configuration prevent a proper identification of the layer texture.

Effect of surface preparation on the morphological evolution

In order to continue the systematic analysis of the Ni / GeSn systems, the impact of the different preparation steps on the morphological evolution was analyzed. Atomic force microscopy images, in tapping mode, with a Bruker FastScan equipment were acquired in order to verify surface stability, when increasing annealing temperatures. Fig. 3.21 shows the AFM images obtained for Ni / Ge 0.90 Sn 0.10 samples annealed at various temperatures by RTA for 30 s in N 2 atmosphere for the different treatments. Whatever the system studied and at low temperature (i.e. for T < 300 ), samples exhibited the characteristic cross-hatch pattern evidenced in section 3.3.2. At 350 , the surface morphology changed for all the samples. A drastic change was observed for the HF treated sample, suggesting that important agglomeration already took place. For the samples without treatment and treated by Ar plasma, small cracks at the sub-micron-scale on the surface were observed, which might be related to grain boundary grooving (prior Effect of surface preparation on the Ni / GeSn system step to agglomeration), explained in section 3.3.2. Meanwhile, for He treated samples no bigger disruption was observed. At higher temperatures (i.e. for T > 350 ), the agglomeration of the surface was evidenced for all the samples. The root-mean-square roughness values further increased with the annealing temperature which can be related to the successive agglomeration and Sn segregation, as observed in Fig. 3.22. The differences observed for the morphological evolution could be related to chemical variations due to the nature of the oxide layers. Indeed, for the reference and He treated sample native oxides are expected to be present in the surface, while a re-growth of the oxides can occur for the HF treated sample and the Ar treated sample is rather oxide-free. However, this hypothesis needs a more detailed analysis.

Effect of surface preparation on the electrical evolution

In order to follow the layers electrical properties, R sh evolution was acquired with a four-point probe meter when the annealing temperatures changed. Fig. 3.23 shows the evolution of the sheet resistance after various RTA from 150 up to 550 .

Whatever the system observed, from the as-deposited state and up to 300 , R sh values remained relatively constant at around 30-35 Ω/sq. Around 200 and up to 250 , a small increase in the values was observed. This was related to the growth of the Ni-rich phase. An important reduction of the R sh values was observed after 300 . Minimum values of R sh were obtained at 350 for all the samples. This reduction occurred concomitantly with the variations in the surface morphology evidenced by AFM, see Fig. 3.21, and can be related to the growth of the Ni-monostanogermanide phase.

Beyond 350 and for all the samples, a strong increase of the R sh values was observed. This behavior was directly linked to the agglomeration of the Ni(GeSn) layer and to the subsequent Sn segregation. A drastic change on the surface morphology and the increase in root-mean-square roughness values was verified in Fig. 3.21 and 3.22. Higher degradation was observed for the HF treated samples. This variations can be related to the nature of the oxides remaining in the surface as mentioned earlier. However, this hypothesis needs a more detailed analysis. On the other hand, the impact of phase sequence and the morphology degradation on the electrical properties stability was verified.

In summary, the different surface treatments realized were verified to not strongly modify the general phase sequence. The same phases grew in a sequential manner. First, the Ni-rich phase was obtained. The Ni(GeSn) followed after. This latter was stable until reaching the system eutectic point. Concerning morphological degradation, a similar sequence was found for all the samples. However, degradation was more critical for the HF treated samples. This phenomenon could be related to the nature of the remaining or re-growth oxides in the surface. The evolution of the electrical properties was also comparable. Minimum R sh values were obtained at 350 for all the samples, when the mono-stanogermanide phase was acquired.

After agglomeration and successive Sn segregation, the R sh values increased. The degradation observed for the HF treated samples was more important. Coupling the WET-HF treatment to a plasma in-situ treatment seems then, a good alternative to ensure clean and smooth surface to which the phase sequence and the morphological or electrical properties evolution would not be extremely impacted.

Ni-rich phase crystallographic structure determination

So far we have shown, using in-situ and ex-situ XRD, that 10 nm thick Ni film deposited on a Ge 0.90 Sn 0.10 layer resulted in the following phase sequence a Ni-rich phase was obtained first followed by the Ni(GeSn) phase, in a sequential manner. There is, however, some controversy regarding the phases that are present at low temperatures (i.e. temperatures below 250 ). Common phases were indexed either on RSM or on the integrated patterns for both samples. First, diffraction peaks located at 2θ = 42.8°and 2θ = 62.2°were respectively linked to TiN (2 0 0) and (2 2 0) in-plane orientations. Then, peaks localized at 2θ = 45.4°a nd 2θ = 66.4°were attributed to Ge (2 2 0) and (4 0 0) in-plane orientations, respectively. For the as-deposited sample (Fig. 3.24 (a)), diffraction spots of Ge (2 2 0), Ni (2 0 0) and TiN (2 0 0) around φ = 45°and spots of Ge (4 0 0), Ni (2 2 0) and TiN (2 2 0) around φ = 0°and 90°were well aligned suggesting a strong epitaxial relationship between these three phases.

The in-plane RSM integrated patterns shown in Fig. 3.25 exhibited differences between the as-deposited sample and the one annealed at 180 . The most obvious difference was the lost of the diffraction peaks related to the Ni peak and located at 2θ = 51.3°and 75.8°(respectively, (2 0 0) and (2 2 0) in-plane orientations) on the pattern corresponding to the sample annealed at 180 . This correlated well with the in-situ XRD presented in Fig. 3.8 where the Ni diffraction peak was lost at this temperature. The missing spots were symbolized by crosses in Fig. 3.24 (b).

For the as-deposited sample (Fig. 3.24 (a) and Fig. 3.25), diffraction spots located By comparing the RSM obtained for the as-deposited sample and the sample annealed at 180 , differences were discernible. First, a slight shift towards higher 2θ angles (from 31.8°to 32.0°) was observed for the first diffraction peak attributed to the Ni-rich phase.

Then, at 180 , diffuse diffraction spots appear at 2θ = 46.3°and at around φ = 95°and 155°(these latter were highlighted by rectangular dotted symbols on Fig. 3.24 (b)). The presence of these diffraction spots was comforted by the fact that the peak attributed to the Ge (2 2 0) at 2θ = 45.0-46.0°in Fig. 3.25 showed an asymmetry and an enhanced intensity compared to the one obtained for the as-deposited sample. This observation was attributed to the fact that there is a dual peak one at 2θ = 45.4°and related to Ge (2 2 0) orientation and one at 2θ = 46.3°related to the Ni-rich phase. For the Ni-rich phase, a shift towards higher 2θ angles was highlighted by out-of-plane in-situ XRD measurements. Fitting of the XRD peak located around 32.0°in Fig. 3.8 was performed and was shown in Fig. 3.26 and 3.27. Two distinct areas were clearly evidenced by these data. Before the Ni consumption, i.e. for temperature lower than 120-130 , the peak intensity was nearly constant without any clear evolution with respect to the temperature increase (Fig. 3.26), the FWHM was constant at about 0.75°and the peak was located around 31.6°(Fig. 3.27). These observations suggested that the Ni-rich phase formed during the TiN deposition realized at 100 for about 20 s and related to a very thin intermixed Ni / GeSn layer did not evolve in this range of temperatures.

From 120-130 , when the Ni layer was totally consumed, the peak intensity strongly raised before monotonously increasing with temperature (see Fig. 3.26). The intensity of the peak abruptly dropped down when the Ni-rich phase was consumed for the benefit of the growth of the mono-stanogermanide phase. The Ni consumption was also related to a clear-cut 2θ angle shift since the peak position evolved from 31.6°to 32.2°and to a FWHM decrease which reached a plateau at 0.47°(see Fig. 3.27). The fact that the intensity monotonously increased while the FWHM remained stable indicated that this phenomenon could not be linked to an increase of crystallinity. On the other hand, thanks to in-plane RSM analyses, this increase of intensity might be explained by in-plane texture variation (appearance of new diffraction spots at 2θ = 46.3°for the sample annealed at 180 ) which promoted preferential out-of-plane orientations. This kind of shift towards higher 2θ angle was reported in literature for the Ni / Ge system and might be linked to stoichiometry variations (changes in Ge concentration) in the ε-Ni 5 Ge 3 phase or stress and thermal expansion phenomena as suggested by De Schutter et al. [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. In this study, this shift occured when the Ni consumption was total and was also evidenced by in-plane RSM analyses. The fact that the shift towards higher 2θ angles happened in out-of-plane and in-plane XRD measurements was a sign that it might not be related to mechanical deformation but to a stoichiometry variation, i.e. lattice constant fluctuation of the ε-Ni 5 (GeSn) 3 phase when the GeSn concentration varied as reported for the Ni / Ge system [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF][START_REF] Ellner | Zur struktur der mischung Nickel-Germanium[END_REF]. In order to comfort this hypothesis, the exact nature of the Ni-rich phase must be determined.

Nature of the Ni-rich phase

Concerning the Ni / Ge system in Fig. 3.2, 4 phases were reported to be thermodynamically stable at room temperature β-Ni 3 Ge (cubic), Ni 2 Ge (orthorhombic), '-Ni 5 Ge 3 (monoclinic) and NiGe (orthorhombic) whereas 5 others phases are only stable at elevated temperatures γ-Ni 3 Ge (?), Ni 5 Ge 2 (hexagonal), ε-Ni 5 Ge 3 (hexagonal), Ni 19 Ge 12 (othorhombic) and Ni 3 Ge 2 (hexagonal) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF][START_REF] Larsson | An electron diffraction study of modulated Ni 1+x Ge B8 type phases[END_REF]. The Ni-Ge crystal structure data was summarized in Table 3.1 in section 3.1.1. For Ni / Ge 1-x Sn x systems, most of the studies stated that the Ni-rich phase corresponded to the Ni 5 (GeSn) 3 phase [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. When indexed, the latter was related to the monoclinic structure stable at room temperature [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. In our case, for the peak located at 31.6°, the Ni 5 (Ge 0.90 Sn 0.10 ) 3 phase represents the most plausible option. Nevertheless, its out-of-plane orientation might be related to (3 1 2) in the monoclinic system ε'-Ni 5 (Ge 0.90 Sn 0.10 ) 3 or (1 0 1) in the hexagonal system ε-Ni 5 (Ge 0.90 Sn 0.10 ) 3 .

To clearly identify the crystal structure of the Ni-rich phase, in-plane pole figures were recorded. Additionally, the location of the low-index substrate (Ge / GeSn) poles with respect to the substrate normal are presented in Fig. 3.28 for (0 0 1) oriented samples, in order to properly identify the epitaxial relationship with respect to the substrate. First, both pole figures obtained at 31.9°were identical. They exhibited mutual diffraction spots at the same χ and φ angles. The spot intensities appeared to be higher for the sample annealed at 180 , this could be link to the quantity of matter which should be more important at this temperatures. Taking as hypothesis a (3 1 2) orientation in the monoclinic system or a (1 0 1) orientation in the hexagonal system for the out-of-plane orientation of the peaks located at 31.6-32.2°, stereographic projection simulations were realized thanks to the CaRIne 3.1 software. For the simulation obtained in the monoclinic system, see Fig. 3.31, results cannot be linked to the obtained pole figures and thus the Ni-rich phase cannot be linked to this kind of crystal structure. On the other hand, the simulation obtained for the hexagonal crystal structure shown in Fig. 3.29 perfectly fitted the experimental pole figures after the application of successive rotations of 90°. 3.4. It was interesting to notice that a similar alignment was reported for a very different system, i.e. Ni / In 0.53 Ga 0.47 As [START_REF] Zhiou | Reaction of Ni film with In 0.53 Ga 0.47 As: Phase formation and texture[END_REF]. Indeed, Ni/InGaAs/InP intermetallics exhibited three kinds of orientations, simplified as group A and group B orientations. Group A corresponded to the projection of (1 0 1 0) plane of the intermetallic phase on the (1 0 0) plane, while group B corresponded to the projection of (1 0 1 1) on the (1 0 0) plane along two in-plane directions of the hexagonal structure. We observed the same alignments as group B orientation for the Ni/InGaAs/InP system [START_REF] Zhiou | Reaction of Ni film with In 0.53 Ga 0.47 As: Phase formation and texture[END_REF]. The poles at χ = 55°might be related to (i) the projection of the (1 1 1) planes of Ge on the (0 0 1) plane, (ii) the projection of the Table 3.4: Summary of the orientations observed for the Ni 5 (Ge 0.9 Sn 0.1 ) 3 phase. Each symbol corresponds to an orientation in Fig. 3.32

Symbol Number

Orientation 1 Ni 5 (GeSn) 3 (1 1 0 1) GeSn (0 0 1) Ni 5 (GeSn) 3 (1 0 1 2) GeSn (0 1 1) 2 Ni 5 (GeSn) 3 (1 1 0 1) GeSn (0 0 1) Ni 5 (GeSn) 3 (1 2 1 0) GeSn (0 1 1)
Thanks to pole figure measurements and the simulation of the stereographic projections, the nature of the Ni-rich phase was clearly identified. This phase is then related to the hexagonal -Ni 5 (Ge 0.9 Sn 0.1 ) 3 metastable phase (P6 3 /mmc space group). Thus, as demonstrated for the Ni / Ge [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF][START_REF] Schutter | Phase formation in intermixed Ni-Ge thin films: Influence of Ge content and low-temperature nucleation of hexagonal nickel germanides[END_REF] and suspected for the Ni / SiGeSn [START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF] systems, a metastable hexagonal stanogermanide phase was obtained at low temperature for the Ni / Ge 0.90 Sn 0.10 system. This identification comforts the fact that the shift observed for the Ni-rich peak at low temperature and concomitant to the total consumption of Ni was due to a stoichiometry variation. Based on the clarification of the Ni-rich phase nature, the Ni-rich peaks observed around 31.6°and 66.5°, in Fig. 3.8 and Fig. 3.10, can be indexed as the Ni 5 (GeSn) 3 phase with (1 0 1) and (2 0 2) out-of-plane orientation, respectively. This comprehensive analysis was based on analysis realized for the Ni / Ge 0.90 Sn 0.10 system. In spite of that, these results can be extended to the Ni / Ge, Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.85 Sn 0.15 systems. Then, the nature of the Ni-rich phase is the same independant on the Sn content in the GeSn layer.

A comprehensive analysis of the crystalline nature of the Ni-rich phase was assessed, thanks to complementary XRD diffraction techniques (i.e. in-situ XRD, in-plane RSM and in-plane PF). These analyses allowed to unambiguously identify the Ni-rich phase as the non-stoichiometric hexagonal ε-Ni 5 (GeSn) 3 phase. This phase allows stoichiometric changes, i.e Ge/GeSn concentration variations in a range between 35 and 44 at.% [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. The stoichiometric variations explained the shift in the Ni-rich phase observed in Fig. 3.8 and 3.9. Additionally, the epitaxial alignment of this phase to the Ge/GeSn substrate was determined. The texture components can be observed in Table 3.4. These results allowed to proper index the Ni-rich phase during the Ni / GeSn solid-state reaction.

However, one can expect that the reaction pathways and rate limiting mechanisms, substrate crystallinity/orientation can have a major impact on the resulting film texture. Other factors such as the possible crystallographic relationships between the film and the substrate, the relative surface and interface energies, the addition of doping or alloying elements to the film, the annealing conditions or surface treatments before metallization could also impact the resulting texture. Due to this range of possibilities, attention concerning crystallographic indexation, should then be paid when process conditions are significantly different. This section 3.5 issued to a publication in the Journal of Applied Crystallography 51, 1133, (2018).

Impact of Sn content on the Ni / GeSn system evolution

Controlling the Sn content and understanding its behavior in the Ni / GeSn intermetallic layers, are both important for device performances. Indeed, higher Sn content in GeSn layers can increase the energy difference between the (L) and the (Γ) valleys and shifts the lasing threshold towards the mid-infrarred (MIR) wavelength region, which is interesting for sensing applications. The challenge of this approach is to preserve the crystallographic quality of the layers [START_REF] Reboud | Optically pumped GeSn micro-disks with 16% Sn lasing at 3.1 µm up to 180 K[END_REF]. In the near future, high quality stano-germanide contacts will be needed (refer to chapter I, section 1.2.1).

Currently, it is not entirely known how Sn incorporates in Ni / GeSn phases. Sn might (i) be soluble in the phase, (ii) segregate, and/or (iii) form Ni-Sn alloys [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. Demeulemeester et al. showed that Sn can homogenously incorporate in the Ni / GeSn intermetallics with 8.6 at.% of Sn addition. However their measurements could not reveal how Sn incorporated in the germanides. After anneal, Sn segregated at the interface between the germanide and the GeSn and at the surface. Additionally, they considered that the effects on the phase sequence and growth kinetics should be limited and that Sn addition is not expected to form an impediment for atoms diffusion [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF]. On the other hand, Sn segregation was evidenced on rapid thermal annealed samples either by XRD or by AFM or SEM images [START_REF] Noroozi | Effect of strain on Ni-(GeSn) x contact formation to GeSn nanowires[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Additionally, Nishimura et al. evidenced that surface degradation can happen at lower temperatures with increasing Sn at.% content.

No matter how Sn incorporates in the intermetallics, we believe that the addition of this element and further segregation when annealing, could have an impact on different aspects of the phase sequence during the solid-state reaction and the Ni-stanogermanides properties. A comprehensive analysis of the behavior of Sn in Ni / Ge 1-x Sn x intermetallics with different Sn concentrations was thus realized. The impact of varying Sn content from 0 up to 15 at.% in the system is presented hereinafter.

The impact on the phase sequence during the SSR and the elements distribution were studied by in-situ X-ray diffraction (XRD) with the parameters explained in chapter II, section 2.1.1. Additionally, transmission electron microscopy (TEM) coupled to energydispersive X-ray spectroscopy (EDS) and electron energy loss spectroscopy (EELS), were realized. EDS and EELS mappings were acquired in a probe-corrected FEI Titan microscope operated at 200 kV. EDS spectra were acquired using the four windowless silicon drift detectors (SDD) integrated into the pole piece, which enabled X-ray detection over a 0.8 srad solid angle. EELS spectra were acquired in STEM nanoprobe mode in dual-EELS mode using a Gatan Imaging Filter (GIF) Quantum with a dispersion of 1 eV/channel and convergence and collection angles of 21 mrad and 40 mrad, respectively. EELS and EDS spectra were acquired thanks to a Gatan Microscopy Suite 3 (GMS3) software with a dwell time of 0.05 s. The pixel size was approximately 1 nm for all maps. The focus was made on the impact of changing the Sn content and the Sn segregation effect on the Ni / GeSn phase sequence during the solid-state reaction.

Sn content variation on the phase sequence, the morphological and the electrical properties evolution

As stated in section 3.2.1, layers of pure Ge and GeSn with 6, 10 and 15 at.% of Sn were epitaxially grown on Ge-buffered Si (1 0 0). Ni and TiN deposition was also realized in these layers under the same conditions described in section 3.2.1. The objective was to verify the effect of Sn content variation on the Ni / GeSn properties evolution during the solid-state reaction.

The Ni / Ge, Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.85 Sn 0.15 phase sequences were comparable to the one observed for the Ni / Ge 0.90 Sn 0.10 system in Fig. 3.8, from a straightforward comparison of in-situ XRD patterns. Differences were however observed concerning the threshold temperatures for phase growth. These differences are summarized in Fig. 3.33 and Table . 3.5 For all the samples, the Ni layer was consumed at 120 . The Ni-rich phase was obtained first and later it was consumed to the benefit of the Ni-monogermanide or Nimonostanogermanide phase. For the Ni / Ge system, the NiGe phase was observed above 180 annealing. Meanwhile the threshold temperature to obtain the Ni(GeSn) phase Impact of Sn content on the Ni / GeSn system evolution increased with Sn content. Indeed, for Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.90 Sn 0.10 systems, the threshold temperatures increased up to 235 and 245 , respectively. Meanwhile for the Ni / Ge 0.85 Sn 0.15 system a decrease down to 230 occurred. A more depth analysis of this behavior is addressed in sections 3.6.6 and 3.6.7.

The variations of the threshold temperatures for phase formation are expected to have an impact on the morphological and electrical properties evolution. For these reasons AFM images, RMS roughness and R sh evolution were acquired for the different systems. Whatever the system studied and at low temperature (i.e. for T < 300 ), samples exhibited the cross-hatch pattern which was preserved up to 300 . For the Ni / Ge system, some variation at the sub-micron scale of the surface was evidenced at this temperature. This behavior was related to the growth of the Ni-monogermanide phase. Indeed, as presented in Fig. 3.33, the phase growth was delayed with the Sn content variation. The NiGe monogermanide phase was then expected to grow at lower temperatures compared to the Ni / GeSn counterparts. On the other hand, the grow of this phase would thus modify / impact the surface morphology. At 350 , the surface morphology changed for all the samples. These changes were linked to the formation of the Ni-monostanogermanide phase and the subsequent degradation of the layers. However, the changes were more As explained in section 3.3.2, the Ni / Ge 0.90 Sn 0.10 sample annealed at 350 showed the presence of a polycrystalline layer of 30 ± 1 nm related to the Ni(GeSn) phase. At 400 , the intermetallic layer completely lost its continuity due to the agglomeration process.

For the Ni / Ge 0.85 Sn 0.15 sample, a similar behavior was evidenced. At 250 , uniform and continuous layers were observed, in which the Ni-rich layer of 17 ± 1 nm was well identified. At 350 , the lost of the Ni(Ge 0.85 Sn 0.15 ) continuity was verified, due to the agglomeration process. These results verify the more narrow thermal stability when increasing the Sn content in the layers and will be further discussed in depth in section 3.6.7.

The layers degradation was due to the amplification of two phenomena the layers agglomeration and the Sn segregation (dots were observed in Fig. 3.34 at 500 ). Indeed, The RMS roughness values increased gradually for all the samples. However higher values were obtained for higher Sn content stacks. This means that, the first phenomenon which degraded the surface morphology corresponded to the agglomeration process which was aggravated by the Sn segregation.

R sh measurements were realized for all the samples. The objective was to verify the Sn content impact on the electrical properties evolution and verify the process window at which values remained low, stable and in an acceptable limit for integration. Fig. 3.37 shows the evolution of the sheet resistance, for different systems and after various RTA from 150 up to 550 .

Whatever the system observed, from the as-deposited state and up to 250 (for the Ni / Ge system) and 300 (for Ni / GeSn systems), R sh values remained constant at around 28-30 Ω/sq. This temperature range corresponded to the Ni-rich phase domain.

An important reduction of R sh values was observed at 300 for the Ni / Ge system. This reduction occurred concomitantly with the variation in the surface morphology evidenced by AFM, see Fig. 3.34, and can be related to the growth of the Ni-monogermanide phase. This phase was verified to grow at lower temperatures compared to Ni / GeSn samples, see Fig. 3.33.

At 350 , minimum values of R sh were reached for all the systems 5.3, 12.9, 9.1 and 12.2 Ω/sq for Ni / Ge, Ni / Ge 0.94 Sn 0.06 , Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 samples, respectively. At this temperature, the growth of the Ni-monogermanide or Nimonostanogermanide phases was verified. R sh values for the Ni / Ge system were lower, therefore GeSn layers were slightly more resistive. These differences might be linked to the intermetallic thickness variation, intrinsic properties of the intermetallic layers induced Then, the morphological degradation had a direct impact on the electrical properties stability. The lost of continuity of the layers had a detrimental impact on R sh values, while Sn segregation aggravated the issues. For Ni / Ge samples, values remained low and constant up to 500 , and slight increased at 550 . For these samples, even if the features of the surface morphology changed since 300 due to phase change, RMS values remained low and stable. Higher increase was observed at 550 . Then, the effect of surface degradation on the electrical properties was verified. The small difference observed in the 2θ position for Ni / Ge with respect to Ni / GeSn systems could be related to Sn introduction in the Ni 5 Ge 3 phase. Sn incorporation was expected to increase the lattice parameter, resulting in a lower 2θ position. When the Ni layer was totally consumed (at 120 ), the Ni-rich peak shifted towards higher 2θ angles. This shift was observed for all samples (Ni / Ge system included). It thus cannot be due to Sn rejection. On the other hand, it was entirely related to stoichiometry variations in the ε-Ni 5 (Ge(Sn)) 3 phase as discussed in section 3.5. After the shift, the Ni-rich phase remained stable until it was consumed for the formation of the mono-germanide or mono-stanogermanide phase. The threshold temperature above which the Ni-rich phase disappeared increased with the Sn content from 180 up to 235 and finally 245 , for 0, 6 and 10 at.% of Sn, respectively. For the Ni / Ge 0.85 Sn 0.15 system (see Fig. 3.40), the trend observed was not entirely fulfilled. The Ni-rich phase remained stable up to 230 . Additionally, two peaks at 30.7°and 44.9°around 220 (dashed area in Fig. 3.40) were also observed. They coexisted with the Ni-rich phase in a small temperature window (220-230 ), until the latter disappeared. After that, the peak intensity increased with the growth of the Ni-monostanogermanide phase. Peaks associated with (0 1 1)/(0 0 2), (1 1 1)/(1 0 2) and (2 1 1)/(2 0 2) crystalline orientation of the Ni(Ge 0.85 Sn 0.15 ) phase at 30.1°, 34.6°and 44.0°, were observed at 230 and remained stable up to 600 . The peaks at 30.7°and at 44.9°vanished at 270 . They could thus be related to the diffraction lines of the (2 0 0) and (2 1 1) β-Sn phase and were due to Sn segregation. This latter will be further discussed in section 3.6.7.

Focus on the Ni-rich phase

Depth profile analysis cross-sectional TEM together with EELS and EDS mappings

In order to complement the in-situ XRD characterizations, samples were annealed in a rapid thermal annealing (RTA) furnace for 30 s in N 2 from 150 up to 550 . The same phase sequence than the one observed by in-situ XRD was evidenced, see Fig. At 350 , a 30 ± 1 nm thick layer was observed due to the Ni(GeSn) phase growth.

The EDS profiles in Fig. 3.42 indicated that the Ge / Sn ratio in the GeSn layer and the intermetallics (both for the Ni-rich and the Ni-monostanogermanide phase) remained the same. The solubility of Sn in the Ni / GeSn intermetallics was thus evidenced. Additionally, a slight increase of Sn content 1-2 at.% around the GeSn / intermetallic vicinity was also observed. The Ni(GeSn) layer was mostly uniform, with a local film thickness decrease due to the grain boundary grooving in the vicinity of grain boundaries (GB). Sn preferentially piled-up along GB and at the interface between the intermetallic and the TiN capping layer. Additionally, a small Sn depletion zone 6 ± 1 nm thick was observed close to the Ni(GeSn) / GeSn interface, with mainly Ni and Ge atoms in it.

At 400 , the Ni(GeSn) layer lost its continuity due to agglomeration (verified by atomic force microscopy (AFM) analyses 3.34). Grains were mainly made of Ni and Ge atoms. Meanwhile, Sn (coming from the intermetallic and the GeSn layer) was predominantly localized at GB and close to the surface. Indeed, the Sn atoms from the GeSn layer that did not react with the Ni atoms segregated on the surface. At 550 , Sn was no longer observed at the grain boundaries but completely distributed on the surface (Sn Impact of Sn content on the Ni / GeSn system evolution islands were observed by AFM 3.34).

Sn behavior during the Ni / GeSn solid-state reaction

As a result of XRD measurements and TEM coupled to EDS / EELS depth profiles, the Sn behavior during the Ni / GeSn solid-state reaction is schematically shown in Fig. 3.43 and summarized as follows (i) From the as-deposited state and up to 200 , the Ni-rich phase was present, with a homogeneous distribution of Sn in it. The 1-2 at.% of Sn increase in the GeSn vicinity was expected to generate a Sn-rich zone around the GeSn / intermetallic interface.

(ii) At 350 , the temperature was considered high enough to allow the atoms to move grain boundary grooving was evidenced by TEM images in Fig. 3.41 together with Sn segregation. Due to the metastability of Sn in this system, segregation was expected to happen first from the Sn-rich zone in the intermetallic / GeSn vicinity. Consequently, a Sn-depletion zone formed around this interface. Additionally, Sn preferentially accumulated at GB and segregated later at the surface. The TiN acted as a diffusion barrier and Sn accumulated below this capping layer.

(iii) At 400 , Sn segregation continued, with higher concentrations near GBs and at the surface. At this temperature, the Ni(GeSn) layer agglomerated. Well spread and separate grains mainly made of Ni and Ge were observed, suggesting that Sn had completely escaped the Ni(GeSn) lattice. The separation in between one grain and another enabled the GeSn semiconductor layer to be in direct contact with the TiN capping layer, namely closer to the surface. This likely generated faster paths for Sn segregation towards the surface. Indeed, Sn atoms in the non-reacted GeSn layer segregated, without further reincorporating in the NiGe grains. Sn atoms were expected to migrate directly to the surface once agglomeration occurred. If some Sn was still present in the intermetallic or in the Ge(Sn) semiconductor layer the expected percentage should be lower than 1 at.% (due to the low solubility of Sn in the Ge lattice at thermodynamic stability). Quantification thanks to EDS of the remaining Sn was difficult, as concentrations would be close to the lower limit of detection of the technique.

(iv) Finally, at 550 , the segregated Sn migrated towards the surface, leaving NiGe grains on a Ge semiconductor layer underneath. If some Sn was still incorporated the percentage was expected to be lower than 1 at.%.

Incorporation of Sn in the intermetallics

During the solid-state reaction, Ni is expected to diffuse towards the Ge(Sn) to form intermetallic phases (i.e. Ni 5 (GeSn) 3 and Ni(GeSn)) in which the Sn concentration should remain steady. However, the systems are highly metastable due to the low solid solubility of Sn in the Ge lattice, i.e. less than 1 at.%. Sn segregation can easily occur if the annealing temperatures exceed the Ge-Sn eutectic point (231 ), the Sn melting point (232 ) or the GeSn epitaxy temperature.

During the phase sequence, the difference in the Ni-rich phase peak position from the Ni / Ge system compared to Ni / GeSn systems indicated the possibility of Sn incorporation into the layer. Additionally, Sn was observed to be regularly distributed in the Ni-rich phase in the as-deposited state and after annealing at 200 . The Ge / Sn ratio in the GeSn layer and in the Ni-rich phase was quantified by EDS in Fig. 3.42 and remained the same after annealing. Sn was verified to be soluble in the Ni-rich phase then.

Although no data about the Ni-Ge-Sn ternary system or the Sn solid solubility in Ni-Ge intermetallics is available in literature, an hypothesis can be drawn from the evidenced behavior. The Ni-rich phase was verified to correspond to the ε-hexagonal phase in the Ni / GeSn system, as for the Ni / Ge system [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF], see section 3.5.2 and Fig. 3.44 (a). This latter is non-stoichiometric and can accept a wide range of Ge concentration between 35-44%, a similar behavior is expected for GeSn. This capability to allow a wide stoichiometry range might favor Sn incorporation in the Ni-rich phase. When the temperature increased, the phase sequence continued with the consumption of the Ni-rich phase to form the orthorhombic mono-stanogermanide phase (see Fig. 3.44 (b)). From the EDS results in Fig. 3.42 Sn was observed to be homogeneously distributed in the Ni(GeSn) layer after annealing at 350 . Thus, even if the phase changed, from an hexagonal towards an orthorhombic structure (as observed in Fig. 3.44), and Sn partially started to segregate, Sn remained soluble in the mono-stanogermanide phase. Sn solubility was thus evidenced to be possible in the Ni / GeSn intermetallics.

However, when the annealing temperatures increased (400-550 ), a complete Sn segregation was promoted, both from the Ni(GeSn) intermetallic and the GeSn semiconductor.

Impact of Sn on the Ni / GeSn phase sequence during solid-state reaction

The phase sequence was modified when Sn content varied. The threshold temperature for the Ni(GeSn) formation increased with the Sn at.% in the GeSn layer. Sn behavior during the solid-state reaction might thus have an impact on the diffusion mechanisms during the phase formation.

For Ni-germanides, Ni is considered to be the main diffusing species to grow the Nirich and NiGe phase [START_REF] Jin | The interfacial reaction of Ni with (1 1 1) Ge, (1 0 0) Si 0.75 Ge 0.25 and (1 0 0) Si at 400[END_REF][START_REF] Patterson | Kinetics of Ni/a-Ge bilayer reactions[END_REF][START_REF] Comrie | Determination of the dominant diffusing species during nickel and palladium germanide formation[END_REF]. Ge diffusion was also observed during the early stages of the NiGe phase formation, however Ni dominated the diffusion in the later stages of the reaction [START_REF] Comrie | Determination of the dominant diffusing species during nickel and palladium germanide formation[END_REF]. Ni is thus expected to be the most mobile species in the GeSn system as well [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF], in which the growth of the Ni-rich and the Ni-monostanogermanide phases are likely controlled by the diffusion step [START_REF] Jin | The interfacial reaction of Ni with (1 1 1) Ge, (1 0 0) Si 0.75 Ge 0.25 and (1 0 0) Si at 400[END_REF][START_REF] Patterson | Kinetics of Ni/a-Ge bilayer reactions[END_REF]. It is however likely that Sn distribution and segregation along grain boundaries and the surface might modify the Ni / Ge(Sn) solid-state reaction.

Due to the role that GB have on atomic transport, Sn segregation around this high energetic area might reduce atoms mobility. However, information about diffusion mechanisms (both for silicides or germanides) is difficult to gather because the analysis of diffusion profiles (radioactive, thermal and so on), is usually modelled with diffusion equations such as Fick's or Nernst-Einstein's which are not sensitive to diffusion mechanisms [START_REF] Gas | Diffusion in Silicides: Basic Approach and Practical Applications[END_REF]. Yet, some indirect evidence can be used in order to distinguish between lattice and grain boundary diffusion. The comparison between the formation kinetics of Ni 2 Si thin films by solid-state reaction (between a Ni layer and a Si substrate) and the mechanisms of self-diffusion in bulk Ni 2 Si suggested that the growth of thin Ni 2 Si films was limited by Ni GB diffusion through the growth of the Ni 2 Si layer [START_REF] Gas | Diffusion mechanism in bulk silicides: Relation with thin film behaviour (case of Ni 2 Si formation)[END_REF]. For Ni-silicides, the relatively low values of the activation energies with respect to the melting points might suggest that the mechanism involved was mostly grain boundary diffusion (both for Ni-rich silicide and Ni-monosilicide) [START_REF] Gas | Diffusion in Silicides: Basic Approach and Practical Applications[END_REF].

Low activation energy (1.3 eV) and a low melting point (756 ) were also associated to NiGe formation [START_REF] Jin | The interfacial reaction of Ni with (1 1 1) Ge, (1 0 0) Si 0.75 Ge 0.25 and (1 0 0) Si at 400[END_REF][START_REF] Patterson | Kinetics of Ni/a-Ge bilayer reactions[END_REF], making it likely that Ni GB diffusion is favored with pure Ge. Futhermore, when thin films are used, the impact of grain boundary on transport is maximized [START_REF] Gas | Formation of silicide thin films by solid state reaction[END_REF]. They offer preferential paths which can enhance atomic mobility, dominating atomic "traffic" [START_REF] Nicolet | Diffusion barriers in thin films[END_REF]. Then, if GBs are modified / saturated by other species, the diffusion rates can be impacted. Such defects will definitely modify Ni diffusion, which is likely the main diffusing species for the Ni / GeSn systems.

Based on (i) in-situ XRD trends and (ii) the Sn distribution and segregation evidenced by TEM images, we believe that Sn acted as a diffusion barrier, more precisely as a "stuffed barrier", as schematized in These types of barriers rely on the segregation of impurities along rapid diffusion paths such as grain boundaries, hampering two-way atomic traffic [START_REF] Nicolet | Diffusion barriers in thin films[END_REF]. At low temperatures, the system was reactive enough to start the solid-state reaction for the formation of the Nirich phase. At this stage, Sn atoms remained in the intermetallic layer, GB were then not crowded. Thus, no impact of the Sn content was evidenced on the Ni-rich phase growth temperature. Then, as the temperature increased, Sn segregation occurred in a range of temperatures close to the formation temperature of the Ni(GeSn) phase. Sn atoms accumulated at grain boundaries (as schematized in Fig. 3.45) which became stuffed barriers that reduced the paths for atoms diffusion, delaying the mono-stanogermanide phase growth. Sn segregation and distribution can also explain the specific behavior of the Ni / Ge 0.85 Sn 0.15 system. This specific case will be discussed in the following section 3.6.7.

3.6.7 Specific case: the Ni / Ge 0.85 Sn 0.15 system For the Ni / Ge 0.85 Sn 0.15 system, there were some disagreements with the phase sequence trends described in previous sections (see section 3.6.1 and 3.6.2). The monostanogermanide phase appeared at 230 , a temperature lower than for the Ni / Ge 0.90 Sn 0.10 system. Additionally, peaks at 30.7°and 44.9°related to Sn segregation were evidenced. Due to the large lattice mismatch between Sn and Ge (14%) and the low solubility of Sn in Ge at thermodynamic equilibrium (less than 1 at.%), all GeSn stacks under study were in a metastable state. Sn segregation was thus expected to occur in all cases (at sufficiently high temperatures). However, in-situ XRD did not show peaks due to Sn segregation for the Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.90 Sn 0.10 samples (see Fig. Takase et al. showed that most of Sn atoms can remain in the Ge lattice after thermal annealing, even when the annealing temperatures were much higher than the eutectic point of the Ge / Sn binary system (231 ) and the Sn concentrations were beyond the solubility limit of the binary system [START_REF] Takase | Behavior of Sn atoms in GeSn thin films during thermal annealing: Ex-situ and in-situ observations[END_REF]. For these two systems (6 and 10 at.%) Sn remained in the intermetallic phases (Ni 5 (GeSn) 3 and Ni(GeSn)) even when the anneal-ing temperature exceeded the eutectic point. As far as the Ni / Ge 0.85 Sn 0.15 system is concerned, it seems that 15 at.% of Sn introduction was too far from the equilibrium solubility, resulting in highly unstable layers in which Sn segregation was promoted. To verify the stability of the phases a heating and cooling cycle, with the same temperature steps and time per step used for the in-situ XRD analysis was performed on Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 samples. Fig. 3.46 shows that, during the heating up, the same phase sequence as the one described in section 3.6.1, occurred for both systems. However, new diffraction lines were observed when the temperature decreased again. For the Ni / Ge 0.85 Sn 0.15 cooling cycle, a peak around 30.7°was observed at 231 while for the Ni / Ge 0.90 Sn 0.10 cooling cycle, peaks around 30.7°and 32.0°appeared at 145 . These peaks were linked to the diffraction lines of β-Sn phase and were due to the recrystallization of Sn that previously segregated.

The sequence for the Ni / Ge 0.85 Sn 0.15 system was thus as follows when heating and reaching temperatures close to the Ge / Sn eutectic temperature (231 ) or Sn melting point (232 ), the ejection of Sn atoms from the intermetallic metastable layer occurred and XRD peaks of the β-Sn phase were observed. When the system reached 270 the peaks vanished, Sn atoms were likely in a non crystalline phase (melted) at the annealing conditions, hence the fading of the diffraction lines. Later on, when the temperature decreased and reached once again values close to the eutectic temperature, Sn solidified in a crystalline phase and the peak at 30.7°coming from β-Sn (2 0 0) was observed again. A similar phenomenon was observed for the Ni / Ge 0.90 Sn 0.10 sample. However, Sn was expected to stay incorporated in the intermetallics at temperatures beyond the eutectic point. Then, no peaks corresponding to Sn diffraction lines were observed at low temperatures during the heating up. Sn might segregate beyond this critical point, at a temperature in which Sn will be melted thus no peak could be detected. Once the system cooled down and temperatures reached values close to the eutectic temperature, the segregated Sn recrystallized and the peaks corresponding to the β-Sn phase were observed at 30.7°and 32.0°. The same effect can explain the fact that Sn diffraction lines were detected for the ex-situ annealed samples in Fig. 3.10. Indeed, for these samples the annealing was performed and XRD patterns were obtained once the samples reached room temperature, which allowed the segregated Sn to crystallize. Sn distribution and segregation evidenced before could explain the phase sequence behavior of the Ni / Ge 0.85 Sn 0.15 system. Due to the high meta-stability, segregation happened at lower temperatures for this system. In a first stage Sn accumulated around GB and in a second stage it migrated towards the surface. This migration was expected to happen at lower temperatures compared to the Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.90 Sn 0.10 systems, "freeing" the grain boundaries. Once the GB were released from Sn accumulation, Ni mobility was enhanced and the Ni-monostanogermanide formation temperature was reduced up to 230 . This behavior enabled Ni / Ge 0.85 Sn 0.15 mechanism to be closer to the one of the Ni / Ge system, due to the low incorporated Sn remaining in the intermetallic. In order to verify the stability of Ni / Ge 0.85 Sn 0.15 system, TEM / EDS analysis of samples annealed by RTA at 250 and 350 were obtained and are presented in Fig. 3.47. TiN, Ni and GeSn layers were well identified in the 250 image. In addition, a layer in which Ni, Ge and Sn atoms were mix was also observed. This layer was related to the Ni-rich phase (i.e. Ni 5 (Ge 0.85 Sn 0.15 ) 3 ). Sn was homogeneously distributed in the GeSn and the Ni-rich layer as observed in Fig. 3.48, with a comparable Ge / Sn ratio as occured for the Ni / Ge 0.90 Sn 0.10 samples. Then Sn incorporation in the Ni-rich phase was then also verified for this 15 at.% system. At 350 , the Ni(GeSn) was obtained. The layer lost its continuity due to grain boundary grooving and agglomeration (verified by atomic force microscopy (AFM) analyses in Fig. 3.12) at this temperature. Even if agglomeration occurred, part of the Sn still remained in the intermetallic and in the GeSn layer in very low proportions, however accumulation around the grain boundaries was important. Then, these Ge 0.85 Sn 0.15 layers were verified to be more unstable than the Ge 0.90 Sn 0.10 layers and strong agglomeration and segregation occurred at lower temperatures as hypothesized.

For this section, the impact of the Sn content and the behavior of Sn in the Ni / GeSn intermetallics during the solid-state reaction was studied. Sn addition modified the phase sequence during the solid-state reaction and impacted the morphological and electrical thermal stability ranges. The Ni(GeSn) phase formation was delayed with Sn addition. Meanwhile, morphological and electrical degradation was enhanced in high-Sn content layers.

It was verified that Sn incorporated in both the Ni-rich and the Ni(GeSn) phases. When the annealing temperatures were high enough to allow the agglomeration process to start, Sn segregation occurred. This segregation took place at the grain boundaries first. Then segregated Sn migrated to the surface. Accumulation of these atoms in the GB delayed Ni diffusion which delayed the mono-stanogermanide phase formation with the increasing Sn content. Nonetheless, when the incorporated Sn at.% was to high (15 at.% for our case), segregation occurred at lower temperatures (detected by XRD and TEM analyses). Not only the Sn that was incorporated in the intermetallics migrated to the surface, but Sn coming from the non-reacted GeSn underneath completely segregated as well.

The results of this section issued to a publication in the Journal of Applied Crystallography 53, 605, (2020).

Ni(GeSn) crystallographic texture: Impact of Sn content

The Ni(GeSn) phase corresponds to the less resistive phase in the Ni / GeSn system. This intermetallic is then the most promising in order to contact electronic and opto-electronic devices [START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF][START_REF] Gupta | GeSn Technology: Extending the Ge Electronics Roadmap[END_REF][START_REF] Yang | Towards direct band-to-band tunneling in p-channel tunneling field effect transistor (TFET): Technology enablement by germanium-tin (GeSn)[END_REF][START_REF] Tseng | GeSn-based p-i-n photodiodes with strained active layer on a Si wafer[END_REF][START_REF] Oehme | GeSn heterojunction LEDs on Si substrates[END_REF]. NiGe or Ni(GeSn) contacts are formed by a solid-state reaction from which Ni(Ge) or Ni(GeSn) polycrystalline layers can be obtained [13, 18, 20, 21, 45, 52-54, 73, 100]. In order to understand the properties of these films, it is essential to have a good understanding of their microstructure.

One important aspect of microstructure corresponds to the texture, which describes the orientations of the different grains from a statistical point of view and the frequency of occurrence of these various orientations within the film [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. For germanides rather complex and rich patterns can be obtained over a variety of substrates, which indicates the incidence of complex texture components [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. These specific orientations can generate specific features in a pole figure as presented in Fig. 3.49. In general, four types of texture can be present when thin films are used (i) random texture (considered as absence of texture), which occurs when there is a uniform distribution of grains orientations on the film. The pole figure (PF) will then be featureless, as observed in Fig. 3.49, due to the high probability of any (h k l) pole to occur.

(ii) Fiber texture, for which one crystallographic axis of the film is parallel to the substrate normal (denominated as the fiber axis), while there is a full rotational degree of freedom around the fiber axis. Concerning the PF features, a single spot in the center for that specific (h k l) plane will be observed. Due to the rotational degree of freedom around the fiber axis, the plane normal to any other crystal plane (different from the fiber axis) will describe a circle around the fiber axis as observed in Fig. 3.49.

(iii) Epitaxy, (or in-plane texture), in which an in-plane alignment fixes all three axes of the grain with respect to the substrate. This will lead to a set of well-defined spots on the pole figures as observed in Fig. 3.49.

Finally, (iv) axiotaxy, this particular texture can be understood as an off-normal fiber texture, i.e. the fiber axis is not perpendicular to the interface but slightly tilted to the normal axis. This generates circles around off-center positions on the pole figure, as displayed in Fig. 3.49 [START_REF] Detavernier | An off-normal fibre-like texture in thin films on single-crystal substrates[END_REF][START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF].

The reaction of a thin Ni film on Ge substrates has been widely studied to verify the phase formation and texture of the layer [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Balakrisnan | Texture of NiGe on Ge (0 0 1) and its evolution with formation temperature[END_REF][START_REF] Van Stiphout | Formation of ultrathin Ni germanides: solid-phase reaction, morphology and texture[END_REF][START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. In general, orthorhombic NiGe phase is characterized by very rich and complex texture, see Fig. 3.50 [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. [START_REF] Van Stiphout | Formation of ultrathin Ni germanides: solid-phase reaction, morphology and texture[END_REF]. On the other hand, annealing the film to temperatures higher than 400 caused the NiGe structure to relax, which was also reflected in its texture axiotaxial components became more intense after relaxation due to growth of axiotaxial grains.

Concerning the Ni / GeSn sample texture, literature is rather scarce. Different studies evaluated the crystalline nature of the Ni-rich phase either on GeSn or SiGeSn substrates. Results showed that the Ni-rich phase corresponded to the hexagonal ε-Ni 5 (GeSn) 3 metastable phase (see section 3.5.2). Meanwhile for the mono-stanogermanide phase, deep studies concerning texture are not available. Wirths et al. showed that the Ni(SiGeSn) layer formed was not in an epitaxial relationship with the SiGeSn layer underneath [START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. Liu et al. evidenced some variations in intensity for some diffracted peaks between the reference and carbon (C) implanted samples. Additionally the formation of certain preferred orientations like NiGeSn (1 1 2) was favored in the presence of carbon [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF]. Schulte-Braucks et al. analyzed Ni y (Ge 1-x Sn x ) z layers fabricated on several GeSn alloys with different Sn contents ranging from 0 at.% to 12.5 at.%. Their θ-2θ XRD structural analysis revealed the co-existence of several Ni y (Ge 1-x Sn x ) z phases, with different crystalline orientations, in which Sn content modified the phases in presence [START_REF] Schulte-Braucks | Schottky barrier tuning via dopant segregation in NiGeSn-GeSn contacts[END_REF].

We decided then, to study the impact of Sn addition on the mono-stanogermanide phase texture. NiGe and Ni(GeSn) films (with various Sn content) formed by the rapid thermal annealing (RTA) of sputter-deposited Ni on Ge (0 0 1) or GeSn (0 0 1) were evaluated. The Ni / Ge system was used as a reference sample while the focus was made on the effects of Sn addition on texture evolution.

Study of the out-of-plane and in-plane orientations of the NiGe and Ni(GeSn) phases: Results

As demonstrated in section 3.6.1, the same phase sequence was obtained for Ni / GeSn systems, with different Sn contents, analyzed during the solid-state reaction. This was determined from a straightforward comparison of in-situ XRD patterns realized under the same conditions. Differences were however noticed concerning the threshold temperatures to form the mono-stanogermanide phase. This trend was verified to be due to Sn segregation impact on Ni diffusion (see section 3.6.4). Additionally, a variation of the out-of-plane orientation of the mono-stanogermanide phase with Sn content was also evidenced. Fig. 3.51 presents patterns of all the samples analyzed by in-situ XRD after cooling down to room temperature.

In the case of the Ni / Ge system, diffraction peaks related to polycrystalline NiGe were observed at 32.5°, 34.5°, 43.5°, 45°, 53°and 54.1°which can be related to (2 0 0), (1 1 1)/(1 0 2), (1 1 2), (2 1 1)/(2 0 2), (0 2 0) and (0 1 3) orientations, respectively. For certain peaks two different reflections can share a similar d-spacing value, then indexation can not be done unambiguously (refer to Appendix A). The out-of-plane preferred orientation, i.e. the out-of-plane peak which presented the higher intensity, was the peak at 34.5°, which can correspond to (1 1 1)/(1 0 2) reflections. In the case of the Ni / Ge 1-x Sn x system, diffraction peaks related to the Ni(GeSn) phase were observed similarly to the Ni / Ge system case. A variation of the amount of out-of-plane peaks was however observed.

For the Ni / Ge 0.94 Sn 0.06 sample five peaks were present around 34.5°, 43.5°, 45°, 53°a nd 54.1°, while the preferred orientation also changed. An increase fn the (0 1 3) peak intensity was observed. Additionally, Sn segregation was verified by peaks at 30.7°, 32.3°, 44.9°and 55.5°related to β-Sn diffraction lines of respectively (2 0 0), (1 0 1), (2 1 1) and (3 0 1) orientations. The intensity of the peaks related to Sn diffraction lines increased with Sn content. The analyses were acquired under similar conditions, this intensity increase can then be related to the quantity of Sn matter which should be more important for samples with higher Sn content.

For the Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 samples, a further reduction in the amount of out-of-plane peaks was noticed, concerning the mono-stanogermanide phase, only three peaks could be related to the Ni(GeSn) phase and were observed around 32.5°, 53°and 54.1°. Then, a reduction of the number of out-of-plane peaks occurred with Sn content. In order to verify that the differences observed were related to a texture variation, in-plane RSM were performed for all the samples. Fig. 3.52 presents the integrated patterns in φ.

Common phases were indexed in the integrated in-plane RSM patterns for all the samples. Diffraction peaks located at 2θ = 62.5°and 2θ = 66.4°were linked to TiN (2 2 0) and Ge/GeSn (4 0 0) in-plane orientations, respectively. Various peaks related to NiGe or Ni(GeSn) in-plane orientations were also observed and identified with gray squares in Fig. 3.52. Sn diffraction peaks were additionally observed for Ni / GeSn samples, in which peaks intensity increased with Sn content as observed for θ-2θ results in Fig. 3.51. As observed for out-of-plane orientations, in-plane orientations varied with Sn addition. A reduction of the amount of peaks related to the mono-stanogermanide phase occurred with the increase of Sn at.%. In order to verify the differences between Ni / Ge and Ni / GeSn systems, full in-plane RSM were analyzed comparing Ni / Ge and Ni / Ge 0.90 Sn 0.10 samples, annealed in-situ up to 600 . Analyses were performed after the complete cooling down of the samples to room temperature, results are shown in Fig. 3.53. For Ni(GeSn) well defined spots were also observed, while the fiber texture was considerably lost. Spots related to in-plane diffraction of TiN (2 0 0) and (2 2 0), and Ge/GeSn (2 2 0) and (4 0 0), with the same epitaxial alignment than for the Ni / Ge system, were observed. On the other hand, spots related to the Ni(Ge 0.90 Sn 0.10 ) phase were evidenced at different 2θ values and are identified in Fig. 3.53 (b). Specific spots around 2θ = 30.5°a nd 2θ = 32.5°at φ = 0 and 90°related to (0 1 1)/(0 0 2) and (2 0 0) in-plane orientation respectively, showed an alignment to Ge/GeSn (2 2 0). Meanwhile a double spot around 2θ = 45.2°with a φ separation of 5°was observed to be slightly misaligned with the Ge/GeSn (4 0 0). Additionally spots at 2θ = 30.5°and φ = 45°, 2θ = 32.5°and φ = 15 and 75°, 2θ = 44.0°and φ = 0 and 90°, 2θ = 55.0°and φ = 12 and 78°and 2θ = 63.7°and φ = 45°, were observed. These spots were related to the diffraction lines of β-Sn with the following in-plane orienations (2 0 0), (1 0 1), (2 2 0), (3 0 1) and (4 0 0), and were due to Sn segregation. An alignment of Sn (2 0 0) with Ge/GeSn (4 0 0) and Sn (2 2 0) or ( 2

Discussion: NiGe and Ni(GeSn) texture

From out-of-plane (Fig. 3.51) and in-plane (Fig. 3.52) XRD analyses, it was observed that the NiGe phase was polycrystalline. Pole figure results verified that the different grains were not randomly oriented. Epitaxial alignments of NiGe grains were evident as well defined spots were present on the pole figures with a fourfold symmetry characteristic of the cubic Ge (0 0 1) substrate, as observed in previous studies [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Balakrisnan | Texture of NiGe on Ge (0 0 1) and its evolution with formation temperature[END_REF][START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. The orientation of the grains of the intermetallic film can be then directly determined by the orientation of the Ge substrate. In addition to the spots corresponding to the NiGe phase, signal from the Ge substrate (identified by white open circles in Fig. 3.54) was also identified in the pole figures at 2θ = 30.6°, 45.5°and 54.1°. These spots corresponded to the signal of (1 1 1), (2 2 0) and (3 1 1) planes respectively.

The complex and rich patterns obtained suggested that several texture components can be present in the film. The epitaxial alignments were determined by simulations with the CaRIne Crystallography 3.1 software. The Ge substrate pole figures were also simulated and are shown in Fig. 3.28, they were used as references for the orientation of the substrate.

For the NiGe phase, four different epitaxial components were identified, which are displayed in Table 3.6. The symbols in Table 3.6 were directly signaled in Fig. 3.54. We observed a preferential alignment of NiGe (0 1 0) and NiGe (1 2 4) with Ge (0 1 1) planes. This alignment was explained by the similar d-spacings of the following planes d N iGe(0

3 0) = 1.116 Å, d N iGe(1 2 4) = 1.084 Å compared to d Ge(4 4 0) = 1.000 Å.
For the Ni / Ge 0.90 Sn 0.10 system complete different pole figures were obtained (see Fig. 3.54). Then, a modification of the monogermanide texture with Sn addition was verified. Rich patterns with a four-fold symmetry, originated from the Ge / GeSn substrate were obtained as for the Ni / Ge system. In addition, spots related to Sn segregation were observed in the pole figures recorded at 2θ = 30.6 and 45.5°, together with the substrate contribution at 2θ = 30.6°, 45.5°and 54.1°. Similar pole figures were acquired for Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.85 Sn 0.15 , which presented similar features and are presented in Appendix B. Then for the rest of the discussion the Ni / Ge 0.90 Sn 0.10 will be addressed as the Ni / GeSn sample.

Symbol Number

Orientation × 1 Ni(GeSn) (0 1 0) GeSn (1 0 0), Ni(GeSn) (1 0 0) GeSn (0 1 1)

• 2 Ni(GeSn) (0 1 2) GeSn (1 0 0), Ni(GeSn) (1 0 0) GeSn (0 1 1)
Table 3.8: Summary of the orientations observed for the β-Sn phase. Each symbol corresponds to an orientation in Fig. 3.54

Symbol Number Orientation 1

Sn (1 0 0) Ge/GeSn (1 0 0), Sn (0 1 0) Ge/GeSn (0 1 0)

⊕ 2
Sn (1 0 1) Ge/GeSn (1 0 0), Sn (0 1 0) Ge/GeSn (0 1 0) * 3

Sn (1 1 0) Ge/GeSn (1 0 0), Sn (1 1 0) Ge/GeSn (0 1 1) 4 Sn (2 1 1) Ge/GeSn (1 0 0), Sn (2 0 1) Ge/GeSn (0 1 1)

For Ni / Ge 0.90 Sn 0.10 samples two texture components concerning the intermetallic and substrate (Ni(GeSn) / GeSn) alignments were observed, see Table 3.7 while four components could be related to Sn / GeSn alignment, see Table 3.8, and were identified in Fig. 3.54. Then, Sn addition modified the monogermanide texture. For this case, a preferential alignment of Ni(GeSn) (1 0 0) with Ge (0 1 1) planes was observed. This alignment was possible thanks to the similar d-spacings of the following planes d N i(GeSn)(5 0 0) = 1.094 Å compared to d GeSn(4 4 0) = 1.000 Å.

Concerning Sn alignment, in section 3.6, it was evidenced that Sn segregation occurred first at grain boundaries and then Sn was distributed in the surface. Islands or dots were observed in the surface of the Ni / GeSn samples, which were related to Sn accumulation (Fig. 3.34). Besides, multiple XRD results showed that Sn solidified in a crystalline phase (Fig. 3.46, 3.51, 3.52 and 3.53). It was however remarkable to identify that even under these conditions the segregated Sn crystallized following an epitaxial relation with the Ge/GeSn substrate instead of randomly oriented grains. This behavior can be related to Sn accumulation at the Ni(GeSn) / GeSn interface at 350 (see Fig. 3.41) which can allow Sn to follow an epitaxial alignment with the substrate. Similar phenomena was observed when Sn segregated from Si/SiSn/Si multilayers. Spots of the β-Sn phase revealing the orientation relations (2 2 0) Si (2 0 0) β-Sn and (2 0 0) Si (1 0 1) β-Sn were found from the studies of Gaiduk et al. [START_REF] Gaiduk | Suppression of tin precipitation in SiSn alloy layers by implanted carbon[END_REF].

For intermetallic thin films, a preferential orientation of grains with respect to the single-crystal substrate underneath will be driven by a minimization of the interfacial energy. A low interfacial energy can be achieved by the optimization of the bondings between the thin film and the substrate. The bonding can be optimized if the interface structure is periodic. Indeed, both the intermetallic film and the substrate will exhibit a periodic crystal structure (assuming that both are crystalline layers). Then, if a matching of the intermetallic lattice and the substrate lattice along a direction within the plane of the interface is possible, periodicity can occur.

For either of the cases presented before (NiGe and Ni(GeSn) phases), we evidenced NiGe / Ge and Ni(GeSn) / GeSn alignments thanks to the respectively alignments of NiGe (0 1 0) and Ge (1 1 0), NiGe (1 2 4) and Ge (1 1 0) and Ni(GeSn) (1 0 0) and GeSn (1 1 0) planes. These phases have different crystal structures, the mono-germanide or mono-stanogermanide phases are orthorhombic while the Ge/GeSn substrates are cubic. For a case like this, an interface matching is unlikely to occurr. However, if low-index planes (both from the intermetallic and the substrate) share similar d-spacing, the matching can take place. This latter corresponds to the principle of axiotaxy alignment.

It was verified, that for either NiGe and Ni(GeSn) phase, the alignment was possible thanks to the similar d-spacings between the planes d N iGe(0 3 0) = 1.116 Å, d N iGe(1 2 4) = 1.084 Å compared to d Ge(4 4 0) = 1.000 Å and d N i(GeSn)(5 0 0) = 1.094 Å compared to d GeSn(4 4 0) = 1.000 Å. Differences were highlighted concerning the texture evolution of Ni / Ge and Ni / GeSn system first, even if both intermetallics corresponded to polycrystalline samples the degree of texture for the Ni(GeSn) phase was greater. Indeed, for the NiGe system a higher incidence of different orientations with respect to the substrate was observed. Then, Ni(GeSn) layers consist of polycrystalline grains in which almost all the grains shared the same orientation. Second, the difference of d-spacing between the intermetallic and the substrate was smaller for the case of Ni(GeSn) / GeSn alignment. Then, the change in texture induced by Sn addition might reduce the interface energy between the film and the substrate.

On the other hand, when axiotaxy occurs, the grains oriented with this texture are left with one rotational degree of freedom around the normal to the axiotaxy axis. This results in circular patterns aligned off-center on the pole figure. Our results did not reveal these kinds of circular patterns (see Fig. 3.54). Nonetheless, the well defined spots obtained were arranged in a way that might project semi-circular lines off-center (see Fig. 3.54 for the PF recorded at for the Ni / Ge system). Due to the axiotaxy-related plane alignment, periodicity can be achieved in one direction. However, if a periodicity is well-defined in a second direction, a two-dimensional periodic interface structure can be obtained. This alignment will then generate well-defined epitaxial spots on the pole figures at certain positions along the axiotaxial lines [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. When comparing NiGe and Ni(GeSn) PF, it can be observed that this degree of well defined spots along axiotaxial lines is minimum or non-existent for the Ni(GeSn) sample compared to the NiGe system. Then, Sn addition might promote a better established epitaxial alignment between the Ni(GeSn) film and the GeSn substrate, while for the NiGe case a complex combination of axiotaxial and epitaxial alignments occurred.

The thin film texture can have an influence in different parameters phase formation, phase stability, morphological and electrical stability [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. Indeed, the threshold temperatures for Ni / Ge phases growth were verified to change when the substrate orientation changed (Ge (0 0 1) or Ge (1 1 1)) [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. Meanhile, Deduytsche et al. showed that agglomeration of a Ni / Si film can occur at lower temperatures on a single crystalline (0 0 1) oriented SOI (Silicon On Insulator) substrate compared to a polycrystalline Si substrate [START_REF] Deduytsche | Hightemperature degradation of NiSi films: agglomeration versus NiSi 2 nucleation[END_REF]. Finally, as seen in section 3.3.3, this morphological degradation can reduce electrical performance. These parameters are indeed critical for contact elaboration. Then, the impact of texture and other critical properties of the intermetallic can be a great insight when taking process step decisions.

For the kinds of samples under study, linking the texture to other properties evolution was not straightforward. Even if a variation of the threshold temperatures for the growth of the mono-stanogermanide phase with the Sn content was evidenced (see section 3.6.1), this phenomenon was related to Sn segregation behavior during the solid state-reaction. Meanwhile both the agglomeration process and further Sn segregation can degrade morphological and electrical properties (see sections 3.3.2 and 3.3.3), which made a direct correlation of the morphological or electrical evolution to a specific texture alignment impossible.

For this section, a comprehensive analysis of the texture variation of the monostanogermanide phase was addressed thanks to the combination of complementary combined XRD techniques, i.e. θ-2θ scans, in-plane RSM patterns and in-plane pole figures.

For the Ni / Ge system, four different epitaxial components were identified, while two texture components concerning Ni(GeSn) / GeSn alignments were observed. For both NiGe and Ni(GeSn) phases, the alignment was possible thanks to the similar d-spacings between the planes d N iGe(0 3 0) = 1.116 Å, d N iGe(1 2 4) = 1.084 Å compared to d Ge(4 4 0) = 1.000 Å and d N i(GeSn)(5 0 0) = 1.094 Å compared to d GeSn(4 4 0) = 1.000 Å. The degree of texture for the Ni(GeSn) phase was greater. Indeed, for NiGe a higher incidence of different orientations with respect to the substrate was observed. Then, Ni(GeSn) layers consist of polycrystalline grains in which almost all the grains shared the same orientation. Additionally, the difference of d-spacing between the intermetallic and the substrate was smaller for the case of Ni(GeSn) / GeSn alignment. Then, the change in texture induced by Sn addition might reduce the interface energy between the film and the substrate.

Chapter III. Ni / GeSn intermetallics properties

Conclusions of Chapter III

An extensive and comprehensive characterization of Ni / GeSn intermetallics was presented and the main results are summarized in Fig. 3.55. The films were obtained thanks to the growth of GeSn layers, (Sn content varied from 6, 10 and up to 15 at.%), by RP-CVD technique together with the use of PVD-magnetron sputtering for the metallization layers (10 nm of Ni capped by 7 nm of TiN).

The solid-state reaction (SSR) of a thin Ni film (10 nm) and a Ge 0.90 Sn 0.10 layer was detailed in terms of phase sequence the evolution of the morphological and electrical properties with the annealing temperatures. The phases appeared sequentially in which a Ni-rich phase was obtained first. This Ni-rich was identified as the metastable, non-stoichiometric, hexagonal ε-Ni 5 (GeSn) 3 phase. The unambiguous identification of the Ni-rich phase crystalline nature was achieved by the combined use of complementary XRD techniques (θ-2θ, in-plane RSM and pole figures). These phase allows stoichiometric changes, i.e Ge/GeSn concentration variations in a range between 35 and 44 at.%. This stoichiometric variations explained the shift in the Ni-rich phase observed when the annealing temperatures increased.

The Ni-rich phase was later consumed to the benefit of the Ni-monostanogermanide orthorhombic phase Ni(GeSn). The texture of this phase, (i.e. preferential grain orientation of the film with respect to the substrate underneath), was assessed. It varied with Sn addition. In general, Sn addition can promote a better established epitaxial alignment between the Ni(GeSn) film and the GeSn substrate underneath. The effect of the surface preparation prior to the metallization step was analyzed. Ar plasma treatment was efficient to eliminate to a great degree Ge and Sn oxides. Mean-while, the nature of the remaining oxides might have a role on the layers properties evolution.

Sn was evidenced to incorporate in both the Ni-rich and Ni(GeSn) phases. When the annealing temperatures increased, Sn segregated first at grain boundaries (GB) and then migrated to the surface. This segregation and distribution modified Ni diffusion rates. Then, the mono-stanogermanide growth was delayed when the Sn content increased. Sn segregation at low temperature for the Ni / Ge 0.85 Sn 0.15 system, allowed a faster Sn distribution on the surface which liberated the GB and enhanced again Ni diffusion.

The morphological degradation, during the SSR, occurred due to two phenomena first the layer agglomeration, second the Sn segregation. The latter enhanced surface degradation which was evidenced by a more abrupt increase of the root-mean-square roughness values. The beginning of surface degradation occurred concomitantly with the growth of the Ni(GeSn) phase, which was confirmed to be the less resistive phase. Indeed, the electrical properties evolution, with temperature, was strictly related to the phase in presence and the morphological evolution. The minimum of R sh values were achieved once the Ni(GeSn) phase was obtained. However, the surface degradation due to agglomeration and Sn segregation immediately triggered an abrupt increase of the R sh values which demonstrated the low thermal stability of Ni / GeSn layers.

Due to these results a part of our investigation were aimed to delay or avoid Sn segregation and layer agglomeration. Strategies concerning pre-amorphization of the GeSn layer, alloy elements and laser annealing were used and will be presented in the following chapter IV.

In the previous chapter, the correlation between the phase formed, the morphological properties, the electrical properties and Sn segregation mechanisms was deeply analyzed, for the Ni / GeSn system. The degradation of the Ni(GeSn) layers, by means of agglomeration and Sn segregation, was highly detrimental for the electrical properties and therefore the contact performances and reliability. In this chapter, three different alternatives targeting to both delay the agglomeration process and/or Sn segregation while ensuring low R sh values were evaluated, as schematized in Fig. 4.1. The first option that will be presented, corresponds to a modification of the GeSn layer prior to Ni metallization by means of ion implantation. More specifically, preamorphization by implantation (PAI), which is an out-of-equilibrium technique that allows to amorphized the semiconductor substrate. This step prior to metallization has shown to improve the properties of silicides [START_REF] Kittl | Ni-and Co-based silicides for advanced CMOS applications[END_REF][START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Rice | The effect of the dose and energy of a pre-silicide implant on nickel silicide formation[END_REF][START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF][START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF][START_REF] Lachal | Effects of Pre-amorphization Thickness and Carbon Implantation on NiPt/Si Silicidation Process[END_REF], germanides [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF][START_REF] Duan | Enhancing the thermal stability of NiGe by prior-germanidation fluorine implantation into Ge substrate[END_REF] or germanosilicides [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF][START_REF] Mao | Impacts of Ge Preamorphization implantation and Si capping on the specific contact resistivity of Ni(Pt)SiGe/p+-SiGe contacts[END_REF]. The impact of different ion implantation on the Ni / GeSn layers is presented in section 4.1.

On the other hand, the addition of elements such as platinum (Pt) or cobalt (Co), showed to enhance silicides, germanides and stanogermanides thermal stability as well [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Geenen | Controlling the formation and stability of ultrathin nickel silicides -an alloying strategy for preventing agglomeration[END_REF][START_REF] Cheng | A novel silicide and germanosilicide by NiCo alloy for Si and SiGe source/drain contact with improved thermal stability[END_REF][START_REF] Rodriguez | Phase formation sequence and cobalt behavior in the Ni 0.9 Co 0.1 system during the thin film solid-state formation[END_REF][START_REF] Wan | Improvement of Nickel-Stanogermanide contact properties by platinum interlayer[END_REF]. Then, the incorporation of either 10 at.% of Co or 10 at.% of Pt into the Ni metallization layer was evaluated and is presented in section 4.2.

Finally, the last alternative corresponds to an innovative annealing technique based on laser annealing. The Ni / GeSn intermetallics are generated thanks to UV-nanosecond laser irradiation. This short time thermal process is expected to prevent Ni(GeSn) degradation and/or Sn segregation to occurred. The results are presented in section 4.3.

Pre-amorphization by ions implantation (PAI)

Ions implantation is a technique in which a beam of ions is accelerated to penetrate into a solid surface. This strike generates physical, chemical or electrical changes into the surface. This can cause a structural change, in which the crystal structure can be damaged or even destroyed. This technique has been largely used for semiconductor processing and traditionally for doping [START_REF] Cristiano | Defect evolution and dopant activation in laser annealed Si and Ge[END_REF], enhance dopants distribution in semiconductor layers [START_REF] Farley | The role of defects in the diffusion and activation of impurities in ion implanted semiconductors[END_REF][START_REF] Lulli | Radiation damage evolution and its relation with dopant distribution during selfannealing implantation of As in silicon[END_REF], elaborate silicon on insulator [START_REF] Lam | Characteristics of MOSFETs fabricated in silicon-on-insulator material formed by high-dose oxygen ion implantation[END_REF][START_REF] Bruel | Siliconon-insulator by oxygen implantation: An advanced technology[END_REF] or enable mesotaxy [START_REF] White | Mesotaxy: Synthesis of buried single-crystal silicide layers by implantation[END_REF][START_REF] Golanski | Low temperature mesotaxy of ion beam synthesized ErSi 2[END_REF]. Moreover, preamorphization by ion implantation has been used, more recently, to increase the thermal stability of silicides [START_REF] Kittl | Ni-and Co-based silicides for advanced CMOS applications[END_REF][START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Rice | The effect of the dose and energy of a pre-silicide implant on nickel silicide formation[END_REF][START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF][START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF][START_REF] Lachal | Effects of Pre-amorphization Thickness and Carbon Implantation on NiPt/Si Silicidation Process[END_REF], germanides [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF][START_REF] Duan | Enhancing the thermal stability of NiGe by prior-germanidation fluorine implantation into Ge substrate[END_REF] or germanosilicides [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF][START_REF] Mao | Impacts of Ge Preamorphization implantation and Si capping on the specific contact resistivity of Ni(Pt)SiGe/p+-SiGe contacts[END_REF], while improving their electrical properties.

Similar thermal stability issues, as for GeSn layers (presented in chapter III, sections 3.1 and 3.3), were observed for silicides and germanides. For the case of thin films of Ni-based silicides, three phases are expected to appear during the solid-state reaction. First, the high resistance phase corresponds to a metal rich silicide (Ni 2 Si). After, a lowresistance monosilicide phase (NiSi) can be formed. This latter is consumed to the benefit of the higher resistance Si-rich phase (NiSi 2 ). Ni-monosilicide phase is the targeted phase for Si-based contact technology, thanks to its low specific resistivity, low contact resistivity, the low Si consumption and low formation temperatures. Unfortunately, NiSi resistivity collapses due to two phenomena (i) agglomeration of the NiSi layer (which can occur at low temperatures) and (ii) the nucleation of the higher resistive phase (NiSi 2 ) [START_REF] Kittl | Ni-and Co-based silicides for advanced CMOS applications[END_REF][START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Rice | The effect of the dose and energy of a pre-silicide implant on nickel silicide formation[END_REF][START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF][START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF][START_REF] Lachal | Effects of Pre-amorphization Thickness and Carbon Implantation on NiPt/Si Silicidation Process[END_REF].

For the case of Ni-germanides, degradation process goes slightly different than for silicides. Indeed, NiGe → NiGe 2 transformation is not thermodynamically stable [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]. This means that the principal mechanism that degrades Ni / Ge intermetallics corresponds to the agglomeration of the layers. This process is expected to occur at low temperatures (300-400 ) and can be related to the low Ni-Ge eutectic point (i.e. 760 [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]). The annealing temperatures can increase atomic mobility and diffusivity resulting in grain boundary grooving and the successive intermetallic agglomeration [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF][START_REF] Duan | Enhancing the thermal stability of NiGe by prior-germanidation fluorine implantation into Ge substrate[END_REF].

For stanogermanides, the causes of the degradation are similar. Indeed, as for the case of germanides, a GeSn rich phase is not expected to grow during the Ni / GeSn solid-state reaction. The main issue concerns then the Ni(GeSn) phase agglomeration followed by Sn segregation (see chapter III, sections 3.1.2, 3.1.3 and 3.6).

For this section, the alternative approach, to enhance Ni / GeSn thermal stability, concerns the modification of the semiconductor layer, by pre-amorphization by implantation (PAI). This additional process increased the thermal stability of Ni-silicides, delayed the high resisitivity NiSi 2 phase formation and enhanced smooth NiSi / Si interfaces [START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF][START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF][START_REF] Lachal | Effects of Pre-amorphization Thickness and Carbon Implantation on NiPt/Si Silicidation Process[END_REF]. Moreover, the use of a PAI process enhanced Ti-silicidation at lower temperatures and improved the electrical properties to some extent [START_REF] Mao | Impact of Ge preamorphization implantation on both the formation of ultrathin TiSi x and the specific contact resistivity in TiSi x /n-Si contacts[END_REF]. For the case of germanides, the use of PAI prior to metallization enhanced thermal and morphological stability, at the expenses of a general increase of the R sh values for a wide range of annealing temperatures. In addition and as evidenced for Ti-silicides, amorphization by implantation enhanced Ti-germanides phase growth at low temperatures and allowed to reach low R sh values [START_REF] Luo | Impact of Ge pre-amorphization implantation on forming ultrathin TiGe x on both n-and p-Ge substrate[END_REF]. Similar results were observed both for Ti-and Ni(Pt)germanosilicides ion implantation promoted the phase formation at lower temperatures and improved the electrical properties [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF][START_REF] Mao | Impacts of Ge Preamorphization implantation and Si capping on the specific contact resistivity of Ni(Pt)SiGe/p+-SiGe contacts[END_REF]. Moreover, Liu et al. studied the impact of C pre-stanogermanidation implant with a dose of 1x10 15 cm -2 at an energy of 2 keV on GeSn layers with 4 at.% substitutional Sn composition. They showed that the thermal stability can indeed be increased after C implantation and observed a delay on the surface degradation and an improvement on the electrical properties [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF].

Ion implantation was thus envisioned as a possible solution to improve stability in high Sn at.% content GeSn layers in order to ensure high quality metallic contacts. For this study, pre-amorphization by Si, Ge, C and Ge + C implantation on GeSn with 10 at.% of Sn content prior to the Ni metallization, was evaluated. The impact of PAI was verified in terms of the phase sequence, the morphological and the electrical properties evolution during the solid-state reaction. The results obtained are presented hereinafter. For these aims, GeSn layers with 10 at.% of Sn were epitaxially grown following the process parameters described in chapter III, section 3.2.1. The PAI process was realized in a conventional 300 mm ion implanter. The following parameters were used 2.8x10 14 at/cm 2 of Si at 10 keV, 1x10 14 at/cm 2 of Ge at 20 keV, 3.5x10 15 at/cm 2 of C at 5 keV and a combination of Ge + C. These parameters were adjusted to achieve an amorphous GeSn thickness of 25-30 nm by simulations. The simulations were made with the CTRIM code which has been properly detailed by Posselt et al. [START_REF] Posselt | Computer simulation of ion implantation into crystalline targets[END_REF]. A calibration was made by Koffel for pure Ge in which it is possible to calculate the profile of implanted ions, the profile of generated defects and even to give the position of the amorphous / crystal limit [START_REF] Koffel | Implantation, diffusion et activation des dopants dans le germanium[END_REF]. We considered for our case that the GeSn layers would have similar properties than pure Ge.

Prior to the metallization, a surface preparation with Ar plasma for 60 s was made. Then the deposition, at room temperature, of 10 nm of a Ni layer capped with 7 nm of TiN was performed by magnetron sputtering at room temperature. The phase formation sequence and crystalline evolution were monitored by in-situ X-ray diffraction (XRD) respecting the parameters specified in chapter II, section 2.1.1. In addition, samples were annealed using a rapid thermal annealing (RTA) treatment in a N 2 environment for 30 s at temperatures ranging from 150 up to 550 . Then, complementary θ-2θ scans were acquired. The evolution of the surface morphology and the electrical properties were measured by tapping mode atomic force microscopy (AFM) and sheet resistance (R sh ) measurements with a four-point probe meter. Transmission electron microscopy (TEM) coupled to energy-dispersive X-ray spectrometry (EDS) images were also acquired, following the parameters displayed in chapter III, section 3.6.

Study of the solid-state reaction

The Ni / Ge 0.90 Sn 0.10 , with beforehand C PAI, phase sequence followed by in-situ XRD is shown in Fig. 4.2 as a contour map. Characteristic peaks of TiN (1 1 1) at 36.4°, Ni (1 1 1) at 44.9°, Si (4 0 0) and Si/Ge/GeSn (4 0 0) around 67°were observed. ). This phase remained stable up to 680 . At this temperature, the dissipation of the diffraction peaks was related to a phase destabilization due to reaching temperatures close to the NiGe-Ge eutectic point [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF].

In order to study the potential impact of the PAI step, similar in-situ XRD analyses were made on Ni / GeSn systems with Si, Ge and Ge + C implantation prior to the metallization, together with a reference sample without implantation. Fig. 4.3 presents the main differences observed. The phase sequence is, overall, the same for all the systems the Ni layer was consumed to the benefit of a Ni-rich phase (Ni 5 (GeSn) 3 ) which was later consumed to form the Ni-monostanogermanide phase Ni(GeSn). Nevertheless, differences concerning the phase thermal stability were observed and are summarized in Table 4.1. [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF][START_REF] Turcotte-Tremblay | Thin film Ni-Si solid-state reactions: Phase formation sequence on amorphized Si[END_REF][START_REF] Chen | Investigation pre-amorphization implantation on nickel silicide formation[END_REF], which might be related to faster Ni diffusion. Indeed, performing PAI on semiconductor substrates modifies the crystalline ordered structure, in which the generation of defects (vacancies and interstitial sites, amorphous region) is promoted [START_REF] Fleetwood | Defects in Materials and Devices[END_REF]. Mass transport in solids is mediated by these kinds of defects (vacancies, interstitial or complex associations of these defects) [START_REF] Gas | Diffusion in Silicides: Basic Approach and Practical Applications[END_REF][START_REF] Jones | Diffusion in Silicon[END_REF]. Then, the increase of defects in the implanted area can influence the reaction kinetics. Changes could be noticed either through a direct impact on the nucleation of the phases or through a modification of the atomic diffusion characteristics, which might modify the phase sequence during the solid-state reaction [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. Then, Ni diffusion rates are expected to increase thanks to the diffusion paths generated by the amorphous-GeSn implanted area when PAI is used. This can explain the low intensity of the Ni peak evidenced in Fig. 4.4 for PAI samples and the consumption of this peak at lower temperatures for C and Ge + C PAI samples. Concerning the Ni(GeSn) growth, a variation of the threshold temperature for the phase formation (and Ni-rich phase dissipation) was observed, when the implanted ion varied (see Table 4.1). For the reference sample, the Ni(GeSn) phase appeared after annealing up to 220 . Ge PAI decreased the threshold temperature for the Ni(GeSn) growth ( 200). Meanwhile, Si PAI slightly increased the growth temperature (225 ). Finally, C PAI importantly delayed the phase growth up to 270 . In order to better un-Chapter IV. Alternatives to improve the Ni / GeSn thermal stability derstand the effect of each implanted species, each case is addressed hereafter separately first, then a comparative discussion is realized.

Impact of Ge PAI

Kittl et al. showed that Ge pre-amorphization accelerated Ni-silicide reaction, while N preimplantation slowed it down [START_REF] Kittl | Ni-and Co-based silicides for advanced CMOS applications[END_REF]. Additionally, Ge PAI can promote Ti-silicidation [START_REF] Mao | Impact of Ge preamorphization implantation on both the formation of ultrathin TiSi x and the specific contact resistivity in TiSi x /n-Si contacts[END_REF], Ti-germanidation [START_REF] Luo | Impact of Ge pre-amorphization implantation on forming ultrathin TiGe x on both n-and p-Ge substrate[END_REF] and Ti-germanosilicidation [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF] at lower temperatures. Ion implantation before the silicidation process can modify surface/interface energies which can modify the diffusion kinetics and diffusion rates for the phase growth. The Ge ion implantation is expected to produce an uniform amorphous layer (PAI process is an outof-equilibrium uniform process) [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. Then, the amorphous layers with an enhanced amount of defects, can allow a Ni uniform reaction at lower temperatures. For the case of silicides, the NiSi growth rate on an amorphous Si region is higher than the one on a crystalline Si region [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. These characteristics enabled the growth of the NiSi phases at lower annealing temperatures. Variations on the NiSi → NiSi 2 transformation were also encountered. A high accumulation of Ge in the NiSi / Si interface retarded the NiSi 2 growth rate. As Ni silicide grows, an amount of Ge atoms are repelled out of the silicide layer and piled up at the silicide / Si interface. The Ge concentration can auto-adjust the Ni-silicide growth rate to produce a smooth interface, change the interface energies and delay the phase transformation [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. Mao et al. also verified Ge accumulation along the TiSiGe / Ge interface, when studying Ti-germanosilicidation by means of TEM coupled with EELS analysis and secondary ion mass spectrometry (SIMS) depth profiles [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF].

For our case, Ge PAI enhanced Ni-stanogermanidation at lower temperatures, either when used alone or combined with C PAI. The discussion about this behavior will be realized after in section 4.1.1.4.

Impact of Si PAI

Concerning the use of Si PAI, Guihard et al. showed by time-resolved X-ray diffraction that after a certain implantation dose (1x10 14 Si/cm 2 ) and amorphized thickness, the silicide phase sequence can be importantly impacted. Additionally, the Ni deposited thickness showed to have an impact as well. First of all, a large temperature delay in the formation of the metal-rich phases was observed. Then, NiSi (obtained from 30 nm Ni deposition) growth temperature underwent an abrupt increase. Meanwhile, for samples realized with 10 nm of Ni deposition, the delay was less strong. Finally, any of the PAI treated samples showed a lower onset temperature for NiSi 2 phase transformation. In general, the NiSi phase stability temperature range was shortened and the NiSi 2 phase appeared at lower temperatures [START_REF] Guihard | Controlling nickel silicide phase formation by Si implantation damage[END_REF]. Ozcan et al. also verified a modulation on the phase sequence when Si PAI, realized either at room or at -100 , was used prior to Ni-silicides formation. The interest of performing PAI at low temperatures was to increase the damage to the crystal structure. Their results showed a NiSi formation at lower temperatures, by about 25 , compared to the non-amorphized substrate. Yet, they considered Si PAI effect on kinetics as negligible and stated that Si implantation into a Si substrate was not expected to have such a significant impact on the diffusion of Ni or Si [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF].

Pre-amorphization by ions implantation (PAI) For our case, Si PAI slightly delayed Ni-stanogermanidation. The discussion about this behavior was realized and presented in section 4.1.1.4. [START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF]. It is then likely that the presence of C in the growing silicide film delayed the diffusion of metal and Si atoms during the silicidation reaction [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. Verifying a delay on NiSi formation was however not possible. Indeed, Luo et al. studied phase sequences by means of ex-situ θ-2θ scans in which the first annealing was realized at 425 . At these temperatures NiSi formation was verified, yet the transformation temperature from the Ni-rich to the Ni-monosilicide phase was impossible to gather. Concerning germanides and stanogermanides, Liu et al. evidenced that C PAI can induce certain NiGe or Ni(GeSn) preferred crystalline orientations during the solid-state reaction [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF][START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF]. However, their ex-situ θ-2θ scans were only taken at high temperatures, then the impact of the phase sequence was difficult to determine.

For our case, C PAI importantly delayed the Ni(GeSn) phase formation up to 270 .

The discussion about this behavior is presented in section 4.1.1.4.

Discussion

For the early stage of the reaction, the enhancement in Ni diffusion for PAI samples can be associated to the increase in diffusion paths generated by the amorphous layer (lost of crystallinity and long-range order, and increase on defects). However, the physical and chemical effect of the implanted element on the substrate can not be neglected [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. The phase sequence can change strongly if the crystalline quality of the substrates is modified. Indeed, the orientation of the single crystalline substrate can have a strong effect on the texture of the phases, leading to different surface and interface energies. This can affect the phase formation sequence by (i) changing the formation temperature of a phase, (ii) enabling the growth of new phases or (iii) avoiding the growth of certain others [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. One important aspect to consider is thus the intermetallics texture and the evolution of this property during the phase sequence.

An influence of texture on the phase formation was evidenced when comparing the reactions for a thin Ni film on Si (0 0 1) and Si (1 1 1) substrates and when comparing the growth of Ni films on Ge (0 0 1) and Ge (1 1 1) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF][START_REF] Gaudet | The thermally-induced reaction of thin Ni films with Si: Effect of the substrate orientation[END_REF]. For the Ni / Ge case, Ni 5 Ge 3 and NiGe phases appeared consecutively on Ge (1 1 1) while they grew simultaneously on α-Ge and Ge (0 0 1). The phase formation temperatures changed strongly with the nature of the substrate being the lowest on α-Ge and the highest on Ge (1 1 1) [START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF]. Furthermore, the phase stability, i.e. the maximum temperature up to which an intermediate phase is stable, before it reacts with the substrate to form a phase that is more thermodynamically stable, can vary with texture variations [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF].

A crystalline difference between the reference GeSn (0 0 1) substrate and the amorphized one is unquestionable. Then a variation of the Ni / GeSn intermetallic phase texture, cannot be ruled out. Fig. 4.5 presents a comparison from the in-situ XRD patterns acquired for the reference sample and the C PAI sample. From a straightforward comparison of both patterns a difference on the out-of-plane orientations, for both the Ni-rich and the Ni(GeSn) phases was evidenced. For the reference sample, two peaks related to the Ni-rich phase were observed at 46.5°and 86.0°. They corresponded to (1 1 0) and (3 0 0) orientations, respectively. Later, three peaks of the Ni(GeSn) appeared at 42.3°, 44.9°and 53.6°, and were related to the (2 1 0), (2 1 1)/(2 0 2) and (0 1 3)/(0 2 0) orientations, respectively. For C PAI sample, two peaks were also related to the Ni-rich phase and were observed around 45.3°and 46.5°. The latter corresponded to (1 0 2) and (1 1 0) orientations, respectively. Meanwhile, only two peaks were detected for the Ni(GeSn) phase at 45.0°and 53.9°, related to (2 1 1)/(2 0 2) and (0 1 3)/(0 2 0) orientations, respectively. Then, a variation of the out-of-plane crystalline orientation of both Ni / GeSn intermetallics was confirmed. Similar behavior occurred for silicides and germanides when C PAI was used as mentioned earlier [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF][START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF][START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF].

The same crystalline orientations for both the Ni-rich and the Ni(GeSn) phases was observed to occur for samples with Si, Ge and Ge + C PAI, from a straightforward comparison of in-situ XRD patterns acquired under the same conditions (full in-situ XRD patterns are shown in Appendix C). The impact of crystalline orientation change, on both phases, could then impact the phase stability and the threshold temperatures for phase growth. Nevertheless, in order to correlate the variations of the phase thermal stability to the crystalline variations with absolute certainty, complementary XRD characterization techniques (i.e. in-plane RSM or pole figures) should be performed. Indeed, as expressed in chapter II, section 2.3, the limitations of θ-2θ configuration prevent a proper identification of the layer texture.

Another aspect to consider is related to the various interactions that the different implanted ions and the GeSn substrate can have during the amorphization. A previous study from Rice, already questioned the effects that the use of different ion implants, with differing masses and atomic sizes, might have on the formation of nickel-silicides [START_REF] Rice | The effect of the dose and energy of a pre-silicide implant on nickel silicide formation[END_REF]. In his study, Rice evaluated and compared the impact of different implanted species, fluorine (F), argon (Ar) and nitrogen (N 2 ), on Ni-silicides formation and properties evolution. A variation on (i) the amorphous surface, (ii) the amorphization extend of the Si substrate and (iii) Ni diffusion rates, was observed when the implanted ion changed [START_REF] Rice | The effect of the dose and energy of a pre-silicide implant on nickel silicide formation[END_REF].

For our case, the PAI parameters of the different ions were adapted to obtain the same amorphous GeSn thickness. It will be verified later on, that indeed the α-GeSn layer thickness was comparable for different samples (see Fig. 4.7). The damage on the GeSn layer can yet vary depending on the ion size or concentration. Dopant concentration can be controlled by the modification of the implantation parameters (dose and energies). Indeed, ion fluences and consequently the ion concentration in the layers, can have an impact on structural and morphological evolution of thin films [START_REF] Choudhary | Nitrogen ion implantation effects on the structural, optical and electrical properties of CdSe thin film[END_REF]. The implantation parameters varied from one species to the other to ensure a comparable amorphized thickness (see section 4.1). For a comparable thickness, the ion concentration can then vary from one system to the other. This variation on concentration can explain the differences observed due to a variation of segregation or distribution of the implanted ion.

On the other hand, lighter atoms (C for instance) can generate less damage than heavier ones (Ge for example) [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. The state of the amorphized layer and the degree of defects within it might thus vary from an implanted species to another which might modify diffusion kinetics. Concerning the size of the species used for this study (Si, Ge and C), the difference in atomic radii (a r ) goes as follow a r C = 67 pm < a r Si = 111 pm < a r Ge = 125 pm. This variation, can then generate a variation on the extent of damage which can enable a variation of diffusion behaviors. Ni diffusion was indeed enhanced (see Table 4.1 and Fig. 4.4) when PAI was used. Yet, the threshold temperatures for Ni consumption and the Ni-rich phase formation remained comparable for the different systems. As a whole, the early stage of the phase sequence remained comparable whatever the implanted species.

Finally, the chemical and/or physical interactions occurring between the implant elements and the intermetallics during the solid-state reaction, has to be considered. Differences on incorporation, distribution and/or segregation of these elements could modify the phase sequence. Once a crystalline intermetallic is formed (the Ni 5 (GeSn) 3 rich phase in our case) the mechanism for atom transport might vary. Indeed, the overall mass transport can be impacted by means of dislocations, grain boundaries, surface and interfaces diffusion of the atoms [START_REF] Gas | Diffusion in Silicides: Basic Approach and Practical Applications[END_REF]. Ge implanted atoms were observed to accumulate/segregate both at the NiSi / Si [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF] and Ti(SiGe) / Ge interfaces [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF]. C however, is expected to accumulate/segregate at the intermetallic / substrate interfaces and at the grain boundaries (GB) [START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. The occurence of this segregation/accumulation, might modify atom diffusion and thermal stability. SIMS depth profiles prove that part of the implanted C distributed in either Ni-silicides [START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF], Ni-germanides [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF] or Nistanogermanides [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF], and accumulated at the NiSi / Si [START_REF] Luo | Effects of carbon pre-silicidation implant into Si substrate on NiSi[END_REF], NiGe / Ge [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF] and Ni(GeSn) / GeSn interfaces [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF]. Moreover, implanted C was observed to segregate at the interfaces of Si/SiGe/Si multi-layered structures, while C coming from Co/C/Si and C/Co/Si interlayers [START_REF] Detavernier | CoSi 2 formation in the presence of carbon[END_REF] and Ni / SiC reactions [START_REF] Keyser | Phase formation and texture of nickel silicides on Si 1-x C x epilayers[END_REF] was expected to migrate at GB as well.

The C solubility in Ni, Ge or Sn, is low. Ni 3 C metastable phase can be only obtained at high temperatures (above 1600 ) and would need the incorporation of around 25 at.% of C to form the phase [START_REF] Hansen | Constitution of Binary Alloys[END_REF]. For the Si-C systems, only the silicon carbide (SiC) compound exist, while for the Ge case this does not occur [START_REF] Hansen | Constitution of Binary Alloys[END_REF]. The amount of C dissolved in Ge has been observed to be low and no evidence of a compound formation was acquired during Ge-C reactions [START_REF] Scaceand | Solubility of carbon in silicon and germanium[END_REF]. Meanwhile, Sn can dissolve only traces of carbon which precipitates on cooling in the form of graphite [START_REF] Hansen | Constitution of Binary Alloys[END_REF]. Then, incorporation of implanted C in the Ni / GeSn intermetallics should not be favorable and C atoms are expected to be expelled out from the lattice. On the other hand, Ge and Si atoms could be, to some extent, incorporated in the Ni-rich or Ni(GeSn) lattice. As displayed earlier, there is an important difference in the atomic radii of the different implanted ions (a r C = 67 pm < a r Si = 111 pm < a r Ge = 125 pm). This parameter might promote a preferential arrangement of the atoms during the phase sequence as schematically presented in Fig. 4.6.

The larger Ge atomic size might prevent it from accumulating at the GB. Implanted Ge will then pilled up around the interfaces (see Fig. 4.6) which could modify the interfaces states and interface energies which could be beneficial for the growth of the Ni(GeSn) phase at lower temperatures.

Thanks to the low dimension, C atoms would be expected to preferentially segregate at GB (see Fig. 4.6). Previously in chapter III, section 3.6, Sn segregation around GB was verified by means of TEM images coupled to EDS and EELS analysis and in-situ XRD patterns. Segregated Sn acted as a diffusion barrier, more precisely as a "stuffed barrier". This phenomenon hampered Ni diffusion delaying the Ni(GeSn) growth when Sn content increased. C is thus expected to equally distribute around the GB and act as a diffusion barrier which might reduce the rate of Ni diffusion to form the Ni(GeSn) phase. The diffusion barrier effect, whatever the nature of the ion, can be enhanced with the ion concentration. Indeed, a higher concentration of C compared to Ge or Si at GB can further prevent Ni diffusion.

The effect of Si PAI on phase nucleation is ambiguous [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF][START_REF] Guihard | Controlling nickel silicide phase formation by Si implantation damage[END_REF]. In our case, a delay on the Ni(GeSn) growth was observed when Si PAI was used. The intermediate size of Si compared to Ge and C, might allow it to homogeneously distribute in different areas of the intermetallic layer at grain boundaries and at the Ni(GeSn) / GeSn interfaces (see Fig. 4.6). Then, the phase sequence was only slightly impacted.

The effects on combining Ge + C PAI were quite interesting. Compared to the C PAI the temperature for Ni(GeSn) formation was reduced. Compared to Ge PAI, the Ni(GeSn) phase formation was delayed, however. We expected then, Ge and C to behave as when used alone. C segregation around GB can then reduce atom diffusion for the phase formation, while Ge accumulation at the interface might modify the interface energies helping to nucleate the Ni(GeSn) phase at lower temperatures. The result was then a competition phenomenon that allowed the system to be in an intermediate regime (see Finally, the temperature for the Ni(GeSn) decomposition was slightly observed to increase up to 680 no matter the PAI treatment. Then the fact of having these ions incorporated, segregated or agglomerated within the intermetallic might modify the eutectic point and increase its crystalline stability. Indeed, it was showed that significant variations in eutectic nucleation and growth dynamics can occur in alloys as a function of the type and amount of elements that are added [START_REF] Dahle | Eutectic modification and microstructure development in Al-Si alloys[END_REF]. However, the difference is small (5 ), then the systems behavior was still comparable.

Impact of PAI on Sn segregation

As previously highlighted, one of the biggest issues of Ni / GeSn intermetallics is their high metastability. Due to the low solid solubility of Sn in the Ge lattice, i.e. less than 1 at.%, Sn segregation can easily occur if annealing temperatures exceed the Ge-Sn eutectic point (231 ), the Sn melting point (232 ) or the GeSn epitaxy temperature. PAI is an out-of-equilibrium thermodynamic process. Then, the possibility of Sn to demix or separate might be low. However, due to the lost of long-range order, Sn might not remain in the Ge lattice during the phase sequence. Then, it is important to check if the PAI step results in Sn segregation since the as-deposited state for Ni / GeSn samples. Fig. 4.7

shows the mapping obtained by TEM coupled to EDS for Ni / Ge 0.90 Sn 0.10 samples with Ge (top row) and Ge + C (bottom row) PAI in the as-deposited state. TiN, Ni, amorphous α-GeSn, crystalline c-GeSn and Ge layers were well identified in the as-deposited images for both samples. Sn was observed to be homogeneously distributed in α-GeSn and c-GeSn layers. α-GeSn thickness for the Ge PAI samples was observed to be of 17 ± 1 nm thick, while for the Ge + C PAI samples it was of 19 ± 1 nm. Additionally, an intermixing layer with Ni, Ge and Sn atoms located between the Ni layer and the α-GeSn was present in both samples. No accumulation of Sn was observed at the interfaces and the surface. We have thus evidenced that the PAI step prior to metallization did not trigger Sn segregation. Ni diffusion happened since the as-deposited state. For Ge PAI, the thickness of the intermixing layer was about 8 ± 1 nm while for the Ge + C PAI sample it was about 7 ± 1 nm. A comparable initial state was thus verified for both samples.

As a complement, XRD analyses were systematically performed on samples at the as-deposited state and annealed, by RTA, from 150 up to 550 . The Sn segregation threshold temperature was extracted from the XRD patterns when peaks related to β-Sn were detected. A summary of the obtained results is plotted in Fig. 4.8.

In the as-deposited state, no Sn segregation was observed by ex-situ XRD scans on the implanted samples, no matter the implanted species. There was no Sn segregation either for the reference sample. Then, implantation at the energies used did not destabilize the metastable GeSn layers, in terms of Sn segregation, (see TEM / EDS analyses in Fig. 4

.7).

For the non-implanted samples, peaks due to β-Sn segregation were observed for annealing at 350 and above (refer to Fig. 3.10 in chapter III, section 3.3.1). Ge implantation did not modify the segregation threshold and peaks due to β-Sn segregation were observed at temperatures of 350 and above as well. Si implantation slightly delayed Sn segregation for 50-100 . Meanwhile, for samples with C PAI, Sn diffraction lines were not observed up to anneals at 500 . C PAI seemed to push up the Sn segregation threshold by 150 to 200 , which is remarkable for these kinds of layers. Gaiduk et al. showed that C implantation on Si/SiSn/Si multi-layers allowed the growth of good crystalline quality and almost complete (90-95%) incorporation of the Sn atoms into lattice positions of the Si matrix in which Sn segregation and precipitation was suppressed [START_REF] Gaiduk | Suppression of tin precipitation in SiSn alloy layers by implanted carbon[END_REF].

The differences concerning segregation and distribution for the different ion implanted explained earlier and schematized in Fig. 4.6 could explain the variations on the Sn segregation temperature thresholds. We previously showed that Sn segregation in Ni / GeSn intermetallics occurred in two steps first, Sn accumulated at the grain boundaries (GB) and near the intermetallic / GeSn interface, and later when the annealing temperatures were high enough, Sn migrated towards the surface (see chapter III, section 3.6). Ge accumulation at the Ni(GeSn) / GeSn interface would not modify the diffusion paths for Sn segregation. Then, the Sn segregation threshold was not impacted. Meanwhile, C preferential accumulation at grain boundaries could act as a diffusion barrier to Sn motion, slowing down segregation and postponing it to higher temperatures. Finally, a portion of Si atoms might accumulate at GB. This distribution is supposed to delay, up to certain extent, Sn segregation. Globally, the effects on Sn segregation could be related to the size of the implanted species as observed for the phase sequence during the solid-state reaction.

It was thus observed that PAI step, prior to Ni metallization, generated an amorphous GeSn layer in which Sn remained homogeneously incorporated. Additionally, a delay on the threshold temperatures for Sn segregation was observed for Si, C and Ge + C PAI samples. The accumulation of Si or C ions at GB can delay the threshold temperature for Sn segregation. This ions can hamper Sn diffusion and avoid its distribution towards the surface.

Impact of PAI on the surface morphology

Fig. 4.9 shows AFM images obtained for all types of samples annealed at various temperatures by RTA. Whatever the system studied, at low temperature (i.e. for T ≤ 300 ), samples exhibited the cross-hatch pattern along the 1 1 0 crystallographic directions, as observed in chapter III, section 3.2.1. Then, the PAI treatment did not modify the original surface morphology. The implantation doses and energies used are considered to be sufficiently low to prevent surface degradation. Such behavior was also observed by Gaiduk et al. when investigating the impact of C implantation on the stability of Si/SiSn/Si layers, in which the crystalline quality was preserved after implantation [START_REF] Gaiduk | Suppression of tin precipitation in SiSn alloy layers by implanted carbon[END_REF]. The cross-hatch morphology was preserved up to 300 . At 350 , the surface morphology for the Ni / Ge 0.90 Sn 0.10 reference system changed. Micro-cracks at the sub-micron scale were observed together with an increase of the surface RMS roughness values as observed in Fig. 4.10. This change was linked to the formation of the mono-stanogermanide phase, the start of the grain-boundary grooving and the subsequent agglomeration.

At higher temperatures, the roughness values gradually increased due to two phenomena further agglomeration and Sn segregation. The surface morphological evolution observed for samples with Si and Ge PAI was very similar and comparable to the reference sample. Thus, the implantation of these species did not modify the morphological evolution with temperature compared to the reference.

The Sn segregation threshold temperature was modified, when Si PAI was used however. The morphological evolution was therefore mainly due to the agglomeration of the Ni(GeSn) phase. Sn segregation aggravated the morphological degradation in a second instance. In contrast, Tsui et al. were able to effectively suppress NiSi agglomeration and maintain smooth interfaces when Ge PAI was used [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. One important aspect to consider is the difference on the degradation mechanisms between silicides and germanides/stanogermanides. For silicides, degradation occurs due to two phenomena (i) agglomeration of the NiSi layer and (ii) nucleation of the NiSi 2 phase. Accumulation of Ge in the NiSi / Si interfaces could (i) modify diffusion kinetics and interface energies, (ii) prevent atoms diffusion and (iii) auto-adjust the Ni-silicide growth rate to produce a smooth interface, which might delay the NiSi 2 formation [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. As a consequence, the Ni-silicides thermal stability was enhanced.

For the case of germanides/stanogermanides, the growth of such a NiGe 2 or Ni(GeSn) 2 is not expected to occur, due to thermodynamic impediments. Degradation of these kinds of films starts due to grain boundary grooving up to the agglomeration process. As illustrated in Fig 3 .13. Then, Ge accumulation at the Ni(GeSn) / GeSn interface is not expected to influence the agglomeration process.

Great differences were highlighted in samples with C implantation. The cross-hatch pattern was still visible up to 450 anneling together with low root-mean-squeare (RMS) roughness values. At 500 , the grain-boundary grooving started. Finally at 550 the layers strongly agglomerated. A similar increase in morphological stability was evidenced when studying the effects of C implantation in Si [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF], Ge [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF] or GeSn [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF]. For 125 Chapter IV. Alternatives to improve the Ni / GeSn thermal stability this case, C preferential segregation at grain boundaries might reduce the grain boundary grooving rates and delay the agglomeration of the layers. Indeed, C in the GB is likely to influence grain-boundary energy and grain boundary diffusion. One can expect that C will enhance the grain boundary cohesion and thus lower the grain-boundary energy delaying finally the agglomeration process [START_REF] Detavernier | CoSi 2 formation in the presence of carbon[END_REF].

In order to complement the AFM results, TEM-EDS analyses were performed on sam- For the Ge PAI sample, the Ni(GeSn) layer lost its continuity due to agglomeration at this temperature. Grains were mainly made of Ni and Ge atoms. Meanwhile, Sn started to accumulate around the grains and near the surface. Agglomeration of the GeSn layer underneath was also observed and Sn also started to segregate. Meanwhile, for the Ge + C PAI sample a very different scenario took place. A continous 29 ± 1 nm thick polycrystalline layer of Ni(GeSn) was observed in which no agglomeration or Sn segregation happened. C implantation was thus demonstrated to positively impact the surface morphology evolution by delaying agglomeration and Sn segregation. In addition, the GeSn layer stayed intact.

The reduction of the grain size of intermetallic phases, can be beneficial for reducing the agglomeration processes [START_REF] Srolovitz | Capillary instabilities in thin films[END_REF][START_REF] Miller | The instability of polycrystalline thin films: Experiment and theory[END_REF]. Indeed, models which explain the mechanisms of agglomeration in thin films stated that the grain boundary grooving was linked to grain size and that a reduction of the grain size might reduce agglomeration [START_REF] Srolovitz | Capillary instabilities in thin films[END_REF][START_REF] Miller | The instability of polycrystalline thin films: Experiment and theory[END_REF]. On the other hand, the average grain size is smaller for silicides formed on α-Si than that on single-crystal Si in samples subjected to the same heat treatments [START_REF] Mao | Impacts of Ge Preamorphization implantation and Si capping on the specific contact resistivity of Ni(Pt)SiGe/p+-SiGe contacts[END_REF]. Additionally, a reduction in crystallite size, when PAI was performed, was observed by different groups. Chen et al. observed this phenomenon when studying the effects of PAI on Ni-silicides formation [START_REF] Chen | Investigation pre-amorphization implantation on nickel silicide formation[END_REF], Duan et al. when using fluorine (F) PAI on Ge for NiGe formation [START_REF] Duan | Enhancing the thermal stability of NiGe by prior-germanidation fluorine implantation into Ge substrate[END_REF] and by Mao et al. when analyzing NiPt-germanosilicidation with Ge PAI [START_REF] Mao | Impacts of Ge Preamorphization implantation and Si capping on the specific contact resistivity of Ni(Pt)SiGe/p+-SiGe contacts[END_REF]. The full width at half maximum (FWHM) of the Ni-monostanogermanide, was therefore measured, from various XRD patterns obtained during in-situ measurements of the different systems. The results are listed in Table 4.2.

According to the Scherrer equation, the crystallite size is inversely proportional to the FWHM [START_REF]Elements of X-ray diffraction[END_REF]. All the FWHM values for Ni(GeSn) phases, in which the C PAI step was made, were higher than the ones obtained for the reference, Si PAI and Ge PAI systems. This means that the grain size of the mono-stanogermanide was reduced due to the im-

Impact of PAI on the electrical properties

In order to link the variations observed on the phase sequence and the morphological evolution of the layers to the electrical properties, the evolution of the sheet resistance (R sh ) after various RTA from 150 up to 550 was followed. The results obtained for all the analyzed systems are presented in Fig. 4.12. Whatever the system observed, from the as-deposited state and up to 200 , the R sh values remained constant at around 28-35 Ω/sq. At 250 , a slight increase of the R sh values was observed. This increase was related to the Ni-rich phase growth. After this, the resistivities decreased due to the growth of the Ni(GeSn) phase (less resistive phase of the Ni / GeSn system). At 350 , minimum values were reached for all the samples. Beyond 350 , a strong increase of the R sh values occurred for the samples without PAI and with Ge or Si PAI. This behavior can be directly linked to the degradation of the Ni(GeSn) layers, observed in Fig. 4.9 and due to agglomeration. Later on, Sn segregation additionally degraded the electrical properties. Thus, as observed for the morphological thermal stability, Si or Ge PAI process did not significantly impact the electrical properties evolution with the annealing temperatures.

Si PAI on Si substrate to form NiSi, was observed to degrade electrical properties. R sh values increased more importantly at lower temperatures than for the reference sample [START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF]. On the other hand, Ge PAI allowed to obtain, generally, lower R sh for Tigermanides and Ti-germanosilicides [START_REF] Mao | Improved Ti germanosilicidation by Ge preamorphization implantation (PAI) for advanced contact technologies[END_REF][START_REF] Luo | Impact of Ge pre-amorphization implantation on forming ultrathin TiGe x on both n-and p-Ge substrate[END_REF]. Meanwhile, Ge PAI on Si substrates to form Ni-silicides improved electrical thermal stability [START_REF] Tsui | Improvement of the Thermal Stability of NiSi by Germanium Ion Implantation[END_REF]. Indeed, Tsui et al. evidenced a delay on NiSi agglomeration and NiSi 2 formation. These latter correspond to the main reasons for NiSi contacts degradation. Then, Ge PAI prevented, for that case, the R sh degradation. As evidenced earlier, Ge or Si PAI, are not expected to modify the grain boundary grooving and the agglomeration process during Ni-stanogermanidation. Then, the degradation of the electrical properties can occur as for the reference sample.

On the other hand, an increase on the electrical thermal stability was evidenced when using C PAI in different systems such as Ni-silicides [START_REF] Nakatsuka | Improvement in NiSi/Si contact properties with C-implantation[END_REF][START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF], Ni-germanides [START_REF] Liu | Effects of carbon pre-germanidation implant into Ge on the thermal stability of NiGe films[END_REF][START_REF] Liu | Effects of Carbon pre-Germanidation implantation on the thermal stability of NiGe and dopant segregation on both n-and p-type Ge Substrate[END_REF] and Ni-stanogermanides [START_REF] Liu | Improvement of the thermal stability of Nickel Stanogermanide by Carbon prestanogermanidation implant into GeSn substrate[END_REF]. The Ni / GeSn samples with C PAI step showed an enhancement in the process window in which the R sh values remained low and stable. Indeed, a plateau was observed from 300 up to 500 for the Ge + C PAI sample and from 350 up to 500 for the C PAI sample. The earlier reduction of R sh values for Ge + C PAI system, can be related to the Ni(GeSn) phase growth at lower temperatures explained in section 4.1.1. Beyond 500 , the sheet resistance importantly increased due to the agglomeration and Sn segregation evidenced in the previous sections 4.1.2 and 4.2.2.3.

It was thus evidenced that the electrical properties deterioration was linked in a first order to the layer agglomeration process and then Sn segregation can worsen the degradation. On the other hand, C implantation had the ability to delay specifically the agglomeration and Sn segregation thanks to the preferential segregation/accumulation at the grain boundaries. This phenomena was then highly beneficial to preserve the electrical stability.

For this section, the impact of Si, Ge and C PAI, on Ni / GeSn phase sequence, Sn segregation, morphological and electrical properties during the solid-state reaction was analyzed. As far as the phase sequence was concerned, the overall evolution, was comparable the Ni 5 (GeSn) 3 phase was obtained first, followed by the mono-stanogermanide (Ni(GeSn)) phase, for all the systems.

Ni diffusion was enhanced in the pre-amorphized layers. This, was related to the increase of defects on the amorphous GeSn layer. Concerning the threshold temperatures for the Ni(GeSn) growth, some variations were encountered from one system to the other. Ge PAI slightly reduced the temperature for Ni(GeSn) growth. Meanwhile Si PAI slightly increased this threshold temperature and C PAI importantly delayed it. Ge accumulation at the Ni(GeSn) / GeSn interface might modify the interface energy for the phase growth, enabling Ni(GeSn) nucleation at lower temperatures. C on the other hand, might accumulate around GB hampering Ni diffusion and delaying phase formation. This same phenomenon might also be the cause of the delay for Sn segregation evidenced for C and Ge + C PAI samples. Indeed, C accumulation at GB can hamper Sn diffusion towards the surface.

Additionally, C implantation positively impacted the surface morphological evolution by delaying the Ni(GeSn) phase agglomeration and by enhancing the process window in which the R sh values remained low and stable. C accumulation at grain boundaries might prevent grooving to occur which prevented agglomeration.

The combination of Ge and C PAI step prior to GeSn metallization can then be considered as an interesting option to obtain the less resistive phase, Ni(GeSn), at low temperatures and with enhanced thermal stability. This alternative could be a solution to ensure high quality and stable GeSn-based metallic contacts that could be CMOS-compatible to be fully integrated in electronic or opto-electronic components.

Addition of an alloy element (Co or Pt) to the Ni layer

As presented previously, (see chapter III and the previous section 4.1), Ni / GeSn intermetallics present low thermal stability like is the case of different silicides and germanides. Ni / GeSn intermetallics degrade when reaching certain annealing temperatures. Generally, the layers agglomerated and this process can be aggravated due to Sn segregation.

Once the layers agglomerate, the electrical properties were observed to be compromised (see chapter III, section 3.3), which represents a great issue for the integration of these contacts in electronic and opto-electronic devices. NiSi compounds, for example, have also shown low thermal stability morphological and electrical degradation can occur due to agglomeration and the phase transition to the high resistive NiSi 2 phase [START_REF] Geenen | Controlling the formation and stability of ultrathin nickel silicides -an alloying strategy for preventing agglomeration[END_REF][START_REF] Cheng | A novel silicide and germanosilicide by NiCo alloy for Si and SiGe source/drain contact with improved thermal stability[END_REF][START_REF] Lee | Improved electrical and thermal properties of nickel silicides by using a NiCo interlayer[END_REF]. The addition of a small amount of an alloying element to the Ni film, such as Pt, Pd, Mo or Re, helped to reduce agglomeration and delayed the NiSi 2 formation [START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Lavoie | Towards implementation of a nickel silicide process for CMOS technologies[END_REF][START_REF] Adusumilli | Silicide-phase evolution and platinum redistribution during silicidation of Ni 0.95 Pt 0.05 /Si (1 0 0) specimens[END_REF].

Furthermore, the addition of Co in Ni silicides gained popularity since the 2010s thanks to the observed increase thermal stability that these alloys present [START_REF] Rodriguez | Phase formation sequence and cobalt behavior in the Ni 0.9 Co 0.1 system during the thin film solid-state formation[END_REF]. Additionally, Co addition enhanced the thermal stability for germanosilicides [START_REF] Cheng | A novel silicide and germanosilicide by NiCo alloy for Si and SiGe source/drain contact with improved thermal stability[END_REF] and germanides [START_REF] Shin | A Study on Thermal Stability Improvement in Ni Germanide/p-Ge using Co interlayer for Ge MOSFETs[END_REF].

Meanwhile, Pt addition was verified to increase the thermal stability of the lowresistivity NiSi phase, improving also the morphological stability and delaying both agglomeration process and the nucleation of high-resistive NiSi 2 as well [START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Mangelinck | Effect of Co, Pt, and Au additions on the stability and epitaxy of NiSi 2 films on (1 1 1) Si[END_REF][START_REF] Lee | On the Morphological Changes of Ni-and Ni(Pt)-Silicides[END_REF][START_REF] Demeulemeester | On the growth kinetics of Ni(Pt) silicide thin films[END_REF]. For germanosilicides, the addition of more than 10 at.% of Pt to Ni, can delay Ge outdiffusion and agglomeration up to 100 [START_REF] Bourjot | Thermal stability of Ni 1-u Pt u (0 < u <0.15) germanosilicide[END_REF]. On the other hand, Wang et al. have previously showed that the addition of Pt to Ni by co-sputtering (the molar ratio of Ni to Pt was about 3) improved thermal stability in a wide range of formation temperatures, with a broader process window for stanogermanidation while suppressing, to some extent, Ni(GeSn) agglomeration in GeSn layers with a Sn composition of 5.3 at.% [START_REF] Wang | Thermally stable multi-phase nickel-platinum stanogermanide contacts for germanium-tin channel MOSFETs[END_REF][START_REF] Wang | Metal stanogermanide contacts with enhanced thermal stability for high mobility germanium-tin field-effect transistor[END_REF].

The use of Ni 0.9 Co 0.1 or Ni 0.9 Pt 0.1 layers, which are classical CMOS compatible alloys, were then envisioned as a solution to increase the Ni / GeSn system thermal stability. Yet, Co or Pt influence on the phase formation sequence of Ni / GeSn system and on the morphological or the electrical properties, has not been fully addressed. We considered, then, highly important to study the evolution of these properties with temperature.

For these aims, a systematic and comprehensive analysis of the Ni (10 at.% of Co or Pt) / Ge 1-x Sn x evolution during the solid-state reaction was evaluated and is presented hereinafter. High quality Ge and GeSn layer with 6, 10 and 15 at.% of Sn were grown following the same conditions as explained in chapter III, section 3.2.1. A surface preparation procedure, was performed using 1% diluted HF during 1 min sequentially rinsed in deionized water and dried with a N 2 gun. Then, 10 nm of Ni 0.9 Co 0.1 or Ni 0.9 Pt 0.1 followed by a 7 nm TiN capping, were realized by physical vapor deposition (PVD-magnetron sputtering) at room temperature in a 300 mm tool.

Impact of Co addition on the Ni / Ge system during the solid-state reaction

Information concerning NiCo on Ge or GeSn layers is scarce. Cheng et al. evidenced an enhancement of the electrical thermal stability when using NiCo (10 at.% of Co) alloy on silicides and germanosilicide. The presence of low-temperature CoSi (for their case) could contribute to the stability of the thin film at high temperatures. Additionally, the uniform distribution of Co and Ge indicated the suppression of Ge precipitation [START_REF] Cheng | A novel silicide and germanosilicide by NiCo alloy for Si and SiGe source/drain contact with improved thermal stability[END_REF]. On the other hand, Shin et al. evidenced an improvement of the NiGe thermal stability when using a Co interlayer (TiN / Ni / Co on Ge). The formation of CoGe compounds, could have promoted the thermal stability improvements. Moreover, the use of a Co interlayer enhanced the current and ohmic characteristics relative to the intermetallic. For these reasons, this section will be divided in two parts. First, the NiCo / Ge system will be analyzed and compared to Ni-germanides. Then, the impact of 10 at.% Co addition on the Ni / GeSn system will be also presented. Co impact was measured by means of phase sequence and the contact properties evolution in terms of surface morphology and sheet resistance (R sh ).

Co impact on the phase sequence

The Ni / Ge and Ni 0.9 Co 0.1 / Ge solid-state reaction were followed by in-situ XRD respecting the parameters described in chapter II, section 2. From 95 , a new diffraction peak appeared around 2θ = 46.6°. This was linked to the formation of the ε-Ni 5 Ge 3 and ε-(NiCo) 5 Ge 3 rich hexagonal phase with (1 1 0) orientation (refer to chapter III, section 3.5 for the full crystallographic analysis of this phase). Co 5 Ge 3 phase presents a B8 2 -type hexagonal close-packed crystal structure, which allows a composition range of 32.5-43.0 at.% of Ge variation [START_REF] Zeng | Thermodynamic description and solidified microstructure of the Co-Ge system[END_REF][START_REF] Dayer | Les systèmes binaires cobalt-germanium et nickelgermanium: étude comparée[END_REF]. As Ni 5 Ge 3 and Co 5 Ge 3 share the hexagonal structure, Co is thus expected to be soluble in the Ni-rich phase lattice.

At 120 , Ni or NiCo layer were totally consumed. After this, the Ni-rich or NiCo-rich phase shifted towards higher 2θ angles. This shift was linked to stoichiometry variations, i.e. lattice constant fluctuation of the hexagonal metastable Ni-rich phase when the Ge composition varied. Around 135 , the Ni-rich or NiCo-rich phases were consumed to the benefit of the Ni-germanide or NiCo-germanide phases and finally disappeared at 150 .

Five diffracted peaks were related to the NiGe phase, while only three peaks were observed for the (NiCo)Ge phase. Then, Co addition might modify the mono-germanide preferential orientation of the grains, i.e. the texture of the film. The NiGe and (NiCo)Ge phase were stable up to 670 . The destabilization of this phase was related to the achieving of the eutectic Ni-Ge temperature [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF].

Around 450-470 a shift of the NiGe and (NiCo)Ge peaks positioned at 53.6°occurred and can be more distinguished in Fig. 4.14. At low temperature (i.e. 150 ) we can actually observed the incidence of two diffraction lines that can be related to the (0 1 3) and (0 2 0) orientations. When the annealing temperatures increased, these peaks entangled while the intensity increased. Finally, at 450-470 a shift towards higher 2θ angles occurred together with an increase of intensity. This shift could be related to (i) a change in the mono-germanide preferential crystalline orientation, (ii) variations on the lattice parameter due to thermal expansion or (iii) the growth of a new phase. tion, so it could be less probable. On the other hand, the thermal expansion of the lattice is expected to occur when the annealing temperatures increase. However, this expansion is expected to increase the lattice size and then shift the diffracted peaks towards lower 2θ angles. In our case, we evidenced a shift towards higher 2θ angles. On the other hand, the growth of a Ge-rich phase (i.e. NiGe 2 or (NiCo)Ge 2 ) is not expected due to thermodynamic constraints [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]. This shift was observed both for the Chapter IV. Alternatives to improve the Ni / GeSn thermal stability Ni / Ge and the NiCo / Ge system. However, the growth of a Co x Ge y compound for the NiCo / Ge system can additionally occur and explain the stronger variation in this system. In order to verify this behavior, in-plane RSM were acquired for Ni / Ge and NiCo / Ge systems for samples annealed up to 700 . The patterns were acquired once the samples were cooled down to room temperature. The full in-plane RSM are presented in Fig. For the in-plane RSM of the Ni / Ge sytem presented in Fig. 4.15 (a), the spot located at φ = 0°and 2θ = 66.5°corresponded to Ge (4 0 0) diffraction. Meanwhile the spots at φ = -45 and 45°and 2θ = 45.3°were related to Ge (2 2 0). Additionally, lines along all φ values were observed at 2θ = 42.4 and 62.0°, which were related to TiN (2 0 0) and (2 2 0), respectively. For this case, TiN layer presented a fiber texture instead of an epitaxial alignment regarding the Ge substrate. This samples were heated up to 700 . At this temperature the NiGe phase was observed to degrade due to reaching the eutectic temperature. Then during the cooling down of the samples, the NiGe phase did not recrystallized.

A different scenario was observed for NiCo / Ge system in Fig. 4.15 (b). The same spots related to Ge (4 0 0) and (2 2 0) were observed at φ = 0°and 2θ = 66.7°, and φ = -45 and 45°and 2θ = 45.1°, respectively. A line at 2θ = 61.8°related to TiN (2 2 0) and another one quite weak at 2θ = 43°related to TiN (2 0 0) were verified. Additionally, spots at different values of φ and 2θ and highlighted by black squares in Fig. 4.15 (b), were observed. These spots were tentatively related to CoGe phase formation in Fig. 4.16.

The phase sequence of a thin Co film deposited on Ge (0 0 1) is expected to evolve from a Co-rich phase towards a Ge-rich phase as follows. The CoGe phase was identified to appear as the first forming phase. After the complete consumption of Co, CoGe phase transformed into the Co 5 Ge 7 one. Finally, CoGe 2 phase was found to form from the reaction between the Co 5 Ge 7 phase and the Ge substrate [START_REF] Opsomer | Influence of Ge substrate crystallinity on Co germanide formation in solid-state reactions[END_REF]. Shin et al. considered that Co 5 Ge 7 formation during TiN / Ni / Co / Ge reaction can occur [START_REF] Shin | A Study on Thermal Stability Improvement in Ni Germanide/p-Ge using Co interlayer for Ge MOSFETs[END_REF]. Meanwhile, Cheng et al. observed a contribution of Co-germanides in diffraction patterns acquired from planview TEM images [START_REF] Cheng | A novel silicide and germanosilicide by NiCo alloy for Si and SiGe source/drain contact with improved thermal stability[END_REF]. Then, all CoGe, Co 5 Ge 7 and CoGe 2 phases were considered as possible candidates to fit the observed peaks. From a straightforward comparison of PDF diffraction files from the ICDD, CoGe phase fits better to the observed peaks in Fig. 4.15 and 4.16. However the presence of different Co x Ge y compounds can not be ruled out. In order to go further on the identification of these phases complementary XRD techniques could be used as in-plane pole figures.

On the other hand, both the incidence of a negative contraction on the smallest axis of the lattice of the orthorhombic monogermanide phase and the growth of a Co x Ge y compound can explain the behavior of the (NiCo)Ge peak at high temperature.

In order to complement the in-situ XRD results, NiCo / Ge samples, obtained under the same process conditions, were ex-situ annealed for 30 s at temperatures ranging from 150 up to 550 in a rapid thermal annealing (RTA) furnace under N 2 environment.

Then, θ-2θ scans were performed in a X'Pert Pro PANalytical X-ray diffractometer. Results are presented in Fig. 4.17. Whatever the system studied and at low temperature (i.e. for T ≤ 300 ), the samples exhibited a cross-hatch pattern along the 1 1 0 crystallographic direction as described in chapter III, section 3.3. At 300 a small disruption of the surface morphology was evidenced for both samples (cracks at the sub-micron scale). This was related to the growth of the NiGe or (NiCo)Ge phase and the start of the grain boundary grooving. Indeed, the growth of the (NiCo)Ge phase was verified at this temperature by XRD diffraction in Fig. 4.17.

At 400 , the grain boundary grooving became more important, together with an increase of the root-mean-square (RMS) roughness values, as observed in Fig. 4.19. At 500 , the surface morphology was degraded for both systems (see Fig. Then, Co addition did not modify importantly the morphological evolution, besides from generally increasing the morphological degradation (verified by the higher surface RMS roughness values). Meanwhile, the formation of a transient phase might promote a morphological improvement of the layers. From the as-deposited state and up to 200 , the R sh values remained stable and around 26-30 Ω/sq. At 250 a slight increase of R sh values was observed for both samples. This increase was related to the growth of the Ni 5 Ge 3 or (NiCo) 5 Ge 3 phases. Indeed, the presence of this phase at this temperature was confirmed for the NiCo / Ge system (see Fig. 4.17). After this temperature, the R sh values importantly decreased. This decrease was related to the growth of the mono-germanide phase, which was verified by ex-situ XRD patterns, (see Fig. 4.17). R sh values were slightly higher for NiCo / Ge system, minimum values were obtained at 350 and were about 7.4 Ω/sq and 12 Ω/sq for NiGe and (NiCo)Ge, respectively. This difference could be related to the higher increase in roughness for the NiCo / Ge system compared to Ni / Ge one when the annealing temperature increased (see Fig. 4. 18 and 4.19). R sh values remained stable up to 500 for the Ni / Ge system and increased after annealing at 550 due to the degradation of the layers. For (NiCo)Ge, the evolution was slightly different. R sh values remained stable up to 400 . At 450 , the R sh values increased. This increase could be related to agglomeration. However, at 500 R sh values decreased and further increased at 550 . Two explanation were considered for this behavior as for the RMS roughness variations observed in Fig. 4.19, (i) the incidence of an artifact or (ii) the formation of Co x Ge y transient phases that could reduce R sh in a small range of temperature. This phase should however be low resistive to counterbalance the morphological degradation at high annealing temperatures. The incidence of diffraction spots that could be related to Co x Ge y compounds formation were indeed identify by in-plane RSM analyses displayed in Fig. 4.15 and 4.16. For Co-Ge system, CoGe phase formation was related to an increase of R sh values up to a maximum and further decrease to stabilize the Co 5 Ge 7 phase [START_REF] Opsomer | Influence of Ge substrate crystallinity on Co germanide formation in solid-state reactions[END_REF]. A similar behavior was observed in the case of the NiCo / Ge samples presented in Fig. 4.20. Even if Ni / Ge and NiCo / Ge morphological evolution remained comparable, R sh values importantly increased at 450 followed by a further decrease at 500 , with a final increase due to layers degradation at 550 for NiCo / Ge samples. However, no peaks related to any of these phases were evidenced in the θ-2θ out-of-plane patterns in Fig. 4.17. The formed phases might present a low thermal stability, these kinds of transient phases are then difficult to identify by ex-situ XRD as described in chapter II, section 2.1.1 or be highly textured, then the θ-2θ configuration did not allowed to detect the phases in which the diffraction did not fulfilled the Bragg's law. A proper identification of these compounds is thus difficult. On the other hand, the fact that at 500 , not only a reduction of the R sh values but also a reduction of the roughness values was observed might indicate that the phenomena are related. The formation of transient phases that might improve the surface state, (improve the surface morphology) could be responsable of the electrical improvement at this temperature. This aspect will be further discussed in section 4.2.2.4.

Contrary to what has been found for silicides and germanides, Co addition to the Ni metallization layer did not generate an remarkable improvement of the electrical or morphological stability in our case. The use of a Co interlayer [START_REF] Lee | Improved electrical and thermal properties of nickel silicides by using a NiCo interlayer[END_REF] or Ni and Co codeposition [START_REF] Deduytsche | Formation and stability of NiSi in the presence of Co and Fe alloying elements[END_REF] on Si helped to improve the R sh thermal stability, while a relatively smoother interfacial morphology was obtained [START_REF] Lee | Improved electrical and thermal properties of nickel silicides by using a NiCo interlayer[END_REF]. For Ni-germanides, the same beneficial effect was evidenced when a Co interlayer was deposited between Ni and Ge (0 0 1) substrate [START_REF] Shin | A Study on Thermal Stability Improvement in Ni Germanide/p-Ge using Co interlayer for Ge MOSFETs[END_REF]. Shin et al. evidenced a concentration increase of Co atoms at the surface of the NiGe layer. The distribution of Co atoms may explain the improved thermal stability [START_REF] Shin | A Study on Thermal Stability Improvement in Ni Germanide/p-Ge using Co interlayer for Ge MOSFETs[END_REF]. Indeed, this accumulation of Co at the surface can modify the surface or the interface energies and delay agglomeration. In our case, Co is expected to incorporate both into the Ni-rich and the Ni-monogermanide phase. The surface or interfaces energies are expected then to remain comparable to the Ni system. On the other hand, Rodriguez et al. evidenced that compared to pure Ni silicides, the addition of 10 at.% of Co extended the thermal stability of (NiCo)Si. This improvement was explained by the reduction of NiSi 2 nucleation temperature when Co was added to the system. The disilicide phase was therefore formed at lower temperature than for the reference Ni film. This early formation of disilicide might avoid film agglomeration and helped to preserve the integrity of the layer thus a low sheet resistance can be maintained at much higher temperatures. For NiCo / Ge, a Ge-rich phase is not expected to appear. Degradation of the layer would occur due to grain boundary grooving followed by the layers agglomeration. Then, Co incorporation in the Ni-rich and Ni-germanide phase seems not to modify the agglomeration process of the layers. The formation of Co x Ge y transient phases at high temperatures might however improve the surface state and help to reduce the R sh values in a small temperature window.

The phase sequence for both Ni / Ge and NiCo / Ge system was comparable. Ni or NiCo layers were consumed to the benefit of the growth of the Ni 5 Ge 3 or (NiCo) 5 Ge 3 . Later, these phases were consumed to the benefit of the Ni or NiCo mono-germanide phases.

Contrary to what has been found for silicides and germanides, Co addition to the Ni metallization layer did not generate an improvement of the electrical or morphological stability in our case. The morphological evolution remained comparable for both samples. Meanwhile, the electrical evolution remained comparable up to a threshold temperature. The growth of the Ni or NiCo-rich phases slightly increased the R sh values. This values later decreased thanks to the growth of the mono-germanide phase. For the Ni / Ge system, R sh values remained low and stable up to 550 , temperature at which the agglomeration of the layers became important. For NiCo samples, an increase of R sh after 400 followed by a decrease at 500 was noticed. This behavior can be related to the growth of Co x Ge y phases. Meanwhile, R sh values further increased at 550 as for Ni / Ge system due to the agglomeration of the layers.

Study of the NiCo / GeSn system 4.2.2.1 Study of the phase sequence

The NiCo / Ge 0.90 Sn 0.10 solid-state reaction followed by in-situ XRD is shown in Fig. 4

.21.

A similar sequence as for Ni / Ge, NiCo / Ge and Ni / GeSn system was observed. Characteristic peaks of TiN (1 1 1) at 36.4°, NiCo (1 1 1) at 44.9°, Si (4 0 0) and Si/Ge/GeSn (4 0 0) around 68°were observed. Overall, the phase sequence was comparable to the previous systems the Ni 0.9 Co 0.1 was consumed to the benefit of a NiCo-rich phase (NiCo) 5 (GeSn) 3 (diffraction peak at 46.5°and 85.6°), in which Co is expected to be soluble. Then, this latter was in turn consumed to form the NiCo-monostanogermanide phase (NiCo)(GeSn) (diffraction peak at 2θ = 53.9°). Nevertheless, few discrepancies are highlighted. First, the NiCo-rich phase was obtained since the as-deposited state. Ni diffusion might then be increased thanks to Co addition. This behavior was not observed when comparing Ni / Ge and NiCo / Ge systems (see Fig. 4.13). For the previous case analyzed, only the metallization layer was modified (NiCo compared to Ni metallization). For this case, the substrate layer was modified thanks to Sn introduction in substitution sites. Then, the surface energies for NiCo / GeSn are expected to be modified in a way that NiCo diffusion can be enhanced.

At 120 , the NiCo alloy was totally consumed. Additionally, the Ni-rich peaks shifted towards higher 2θ angles, which was linked to stoichiometry variations, i.e. lattice constant fluctuation of the hexagonal metastable ε-Ni 5 (GeSn) 3 phase when the Ge(Sn) composition varied, as described in chapter III, section 3.5.

A new diffraction peak appeared at 53.4°related to the (NiCo)(GeSn) phase, while at 210 , the diffraction peaks linked to the Ni-rich phase dissipated. The thermal stability of the NiCo-rich phase increased compared to the NiCo / Ge system. Indeed, the (NiCo) 5 Ge 3 dissipated at 150 . This can be related to the phenomenon described in chapter III, section 3.6 Sn segregation at grain boundaries is expected to act as a diffusion barrier which can hamper Ni diffusion, which can increase the Ni-rich phase stability while delaying the Ni-monogermanide growth.

On the other hand, a change on the preferential orientation of the NiCo-monogermanide phase was noticed. Indeed, only one out-of-plane peak related to the (NiCo)(GeSn) phase was noticed. This behavior was also evidenced for Ni / GeSn systems (refer to chapter III, section 3.7). Then Sn addition helped to induce a preferential orientation for both Ni-and NiCo-stanogermanides. The (NiCo)(GeSn) phase remained stable up to 655 . Compared to NiCo / Ge and Ni / GeSn, the range of thermal stability of this phase was slightly reduced, see Table 4 The dissipation of the mono-stanogermanide diffraction peaks could be related to a phase destabilization linked to the reach of the NiGe-Ge eutectic point [START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]. The addition of Sn in the Ge lattice could indeed modify this temperature. However, no bigger difference was noticed for the Ni / GeSn system compared to the Ni / Ge or NiCo / Ge ones. On the other hand, after the dissipation of the (NiCo)(GeSn) phase, the appearance of a weak diffraction peak at 34.7°was observed (highlighted by a dotted circle in Fig. 4

.21).

This peak is more noticeable in Fig. 4.22.

This diffraction peak was not detected during comparable in-situ XRD measurements on Ni / GeSn and NiCo / Ge systems. Thus, this peak can be due to the concomitant presence of Co and Sn and could be linked to the formation of Co x Sn y compounds. The formation of these kinds of compounds suggest the presence of Sn in the metallic state. However, no diffraction peaks were related to Sn phase during in-situ analysis (see Fig. 4.21). A first step to identify the nature of this compounds and Sn segregation will be then drawn hereinafter. However, the growth of CoGe compounds, like for the case of the NiCo / Ge system, can not be totally ruled out.

Impact of Co addition on Sn segregation

Beyond 230 , Sn is expected to be in a melted state and no diffraction peak could then be detected with in-situ XRD analyses (refer to chapter III, section 3. The same phase sequence as the one observed from in-situ XRD was evidenced. Indeed, NiCo layer was consumed to the benefit of the NiCo-rich phase. Then, the NiCo-Addition of an alloy element (Co or Pt) to the Ni layer monostanogermanide was obtained. Additionally, a peak at 55.2°was observed after annealing at 400 . This latter can be related to the diffraction line of β-Sn (3 0 1) and due to Sn segregation. Thus, for each system, namely Ni / and Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 , the Sn segregation threshold temperature was extracted from the XRD patterns, the latter is plotted in Fig. 4.24. Concerning the Ni / Ge 0.90 Sn 0.10 system, no diffraction peaks linked to the β-Sn phase were discernible at 300 . Diffraction peaks related to the β-Sn phase were however observed since 350 . Thus, the Sn segregation threshold for this system ranges between 300 and 350 . In the same way, the Sn segregation threshold for the Ni 0.9 Co 0.1 were ranging between 350 and 400 . Thus, the addition of Co slightly delayed Sn segregation temperature. A deep analysis of Sn segregation will be presented in section 4.2.4 to be able to compare Co and Pt addition effects.

On the other hand, Sn orientation seemed to change compared to the results displayed in Fig. 3.51 and 3.52. Indeed, only one peak at 55.2°was evidenced. An in-plane RSM pattern from a NiCo / GeSn sample annealed up to 700 was then acquired, after cooled down to room temperature, and is presented in Fig. 4.25.

A complete different pattern to the one obtained for NiCo / Ge system in Fig. 4.15 was evidenced. A peak at φ = 2.5°and 2θ = 45.5°was related to Ge (2 2 0), while the peak at φ = 47.4°2θ = 65.7°was linked to Ge (4 0 0). Lines along all φ values were observed at 2θ = 30.5°, 31.9°, 36.1°, 42.3°, 44.7°, 55.0°and 61.9°. The peaks obtained at 2θ = 30.5°, 31.9°, 44.7°and 55.0°can be related to the diffraction of β-Sn and correspond to the following orientations (2 0 0), (1 0 1), (2 1 1) and (3 0 1) respectively. Meanwhile, the lines located at 2θ = 36.1°, 42.3°and 61.9°were related to TiN diffraction lines with the following orientations (1 1 1), (2 0 0) and (2 2 0), respectively.

A change on the preferential texture of the crystallized Sn was thus evidenced. Indeed, in chapter III, section 3.7, an epitaxial alignment of Sn to the Ge/GeSn substrate was evidenced in which four epitaxial components were highlighted. One first hypothesis, for this change on preferential texture, could be related to Co interactions with segregated Sn. Different peaks were additionally identified (highlighted with a black square in Fig. 4.25) and could be related to Co x Sn y compounds (although the formation of Co x Ge y compounds as for the NiCo / Ge system, can not be ruled out). Additionally, an impact on the Sn segregation temperature was highlighted. Then, an effect of Co addition on

Impact of Co addition on the Ni / GeSn surface morphology

Fig. 4.26 shows the AFM images obtained for samples annealed at various RTA temperatures for different systems. The goal was to verify the impact of Co addition on the surface morphological evolution with annealing temperatures.

Whatever the system studied and at low temperatures, the samples exhibit a crosshatch pattern along the 1 1 0 crystallographic direction described in chapter III, section 3.3. This morphology was preserved until 300 , for the different samples. Yet, a slight modification of the NiCo / Ge samples were observed at the sub-micron scale. This was related to the growth of the (NiCo)Ge phase, verified by XRD and R sh analysis, see Fig. 4.17 and 4.20, respectively. At 350 , the surface morphology for Ni and Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 systems drastically changed with a strong increase of the surface roughness values as observed in Fig. 4.27.These changes were linked to the formation of the mono-stanogermanide phase and its subsequent agglomeration. At higher temperatures, the roughness gradually increased due to the amplification of both phenomena the mono-stanogermanide phase agglomeration and the Sn segregation. The surface morphology evolution observed for both Ni and Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 systems was similar, however higher RMS roughness values were reached with Co addition. At 550 a slight reduction on the roughness systems the decrease occurred at 500 . Meanwhile, for the NiCo / Ge 0.94 Sn 0.06 system, the roughness values remained comparable after 400 and for the NiCo /Ge 0.90 Sn 0.10 system this decrease occurred at 550 . As explained in section 4.2.1.2, this behavior could be related to the formation of transient phases that can modify the surface states.

Impact of Co addition on the Ni / GeSn electrical properties

Fig. 4.29 shows the evolution of the sheet resistance after various RTA for 30 s in N 2 environment. The objective was to verify the impact of Co addition on the electrical properties evolution of the Ni / GeSn system.

Whatever the system observed, from the as-deposited state and up to a threshold temperature, the sheet resistance values (R sh ) remained constant at around 28-35 Ω/sq. This temperature range corresponds to the Ni(Co)-rich phase domain. At 300 , R sh values decreased for the NiCo / Ge systems due to the growth of the (NiCo)Ge phase.

Both for Ni / GeSn and NiCo / GeSn system the reduction was delayed up to 350 which correlated well to the delay on the morphological surface modification in Fig. 4.26 and the delay on the mono-stanogermanide phase formation verified on section 4.2.2.

For the NiCo / Ge system R sh values remained stable up to 450 , at which a slight increase of the values occurred followed by a further decrease. Later on, the R sh increased at 550 due to the agglomeration of the layers. The increase and further decrease observed in the 450-500 temperature range could be related to Co x Ge y compounds formation as explained in section 4.2.1.3. Meanwhile, beyond 350 and for both Ni and Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 systems, a strong increase of the R sh values was observed. This behavior was directly linked to the degradation of the intermetallic layer and to the subsequent Sn segregation verified by XRD patterns (see section 4.2.2.2). Thus, as observed for the morphological properties, the addition of Co did not significantly impact the electrical properties of the Ni / GeSn system.

On the other hand, a decrease of the R sh values at high temperature (i.e. 550 )

was observed for the Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 system and might be linked to the formation of Co x Sn y or Co x Ge y compounds as described in section 4.2.1.3 and 4.2.2.2. In order to verify the incidence of this R sh values reduction on samples with Co addition, NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.85 Sn 0.15 samples were also annealed by RTA and analyzed thanks to four-point probe meter. The electrical evolution is presented in Fig. 4.30 together with NiCo / Ge and NiCo / Ge 0.90 Sn 0.10 samples for comparison.

A similar evolution was then noticed for the Ni / Ge and NiCo / Ge 0.85 Sn 0.15 systems, and NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.90 Sn 0.10 systems, respectively. The R sh values remained stable up to 300 (for the Ni / Ge system) and 350 (for GeSn-based systems) at which the sheet resistance importantly decreased due to the growth of the mono-germanide or stanogermanide phase. After this temperature, the R sh values increased due to agglomeration of the layers. However, a R sh decrease at high temperatures was evidenced at 500 for the Ni / Ge and NiCo / Ge 0.85 Sn 0.15 systems, and at 550 The R sh values increase at high temperatures occurred at 550 for NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.90 Sn 0.10 systems and at 500 for NiCo / Ge and NiCo / Ge 0.85 Sn 0.15 systems. This temperature variation can be related to Sn segregation at low temperature for the NiCo / Ge 0.85 Sn 0.15 system. Indeed, in chapter III, section 3.6, it was observed that Sn segregation modified the phase sequence. Sn segregation at low temperatures for the Ni / Ge 0.85 Sn 0.15 system and migration towards the surface freed the grain boundaries from Sn accumulation and enhanced Ni diffusion to grow the Ni(GeSn) at lower temperatures than for the Ni / Ge 0.94 Sn 0.06 or Ni / Ge 0.90 Sn 0.10 systems. The Ni / Ge 0.85 Sn 0.15 behavior approached then the one of the Ni / Ge system. For this case, Sn segregation at lower temperatures for the NiCo / Ge 0.85 Sn 0.15 system might promote Sn segregation towards the surface at lower temperatures, while Sn might remain incorporated on the NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.90 Sn 0.10 systems. Then, NiCo / Ge and NiCo / Ge 0.85 Sn 0.15 systems could be prone to the formation of Co x Ge y (which modified the morphology and the electrical properties at 500 ). Meanwhile, the remaining Sn in the intermetallics could rather promote the formation of Co x Sn y compounds for NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.90 Sn 0.10 systems at 550 . To properly identify the reasons of these behavior, further analysis to identify the formation and nature of tran-sient phases at high temperatures is necessary.

For this section, the impact of Co addition to the Ni metallization layer for contacting Ge or GeSn layers was evaluated. As far as the solid-state reaction was concerned, the overall phase sequence remained the same either when comparing Ni / Ge and NiCo / Ge systems or Ni / GeSn and NiCo / GeSn systems. At low temperature, a Ni-rich phase was obtained first, the latter being consumed to form the mono-stanogermanide phase.

The addition of Co showed a positive impact on reducing the formation temperature for the (NiCo) 5 (GeSn) 3 and (NiCo)(GeSn) phases. Meanwhile, the Sn segregation threshold was delayed up to 400 . The change on the preferential orientation of β-Sn induced for this case or the formation of Co x Sn y / Co x Ge y compounds might change Sn diffusion kinetics towards the surface. Co addition might also contribute to lower the surface / interface energies which positively impacted the Sn segregation for this case.

Regarding the morphological or electrical properties, it was highlighted that Co addition had a weak influence. Indeed, the surface morphology and the sheet resistance evolution as a function of annealing temperature were quite similar either when comparing Ni / Ge and NiCo / Ge systems or for the Ni / Ge 0.90 Sn 0.10 and the NiCo / Ge 0.90 Sn 0.10 systems. However, a reduction of R sh values at high temperatures was observed for NiCo-systems. Then, the growth of transient phases at high temperatures might enable an improvement of the electrical properties. Ni 0.9 Co 0.1 alloy was finally verified to be suited to contact GeSn layers. The phase sequence, morphological and electrical evolution was comparable to Ni systems. Additionally, a delay on Sn segregation was noticed. However, the degradation of the layers due to the low thermal stability is still a problem. The effect of Pt is analyzed hereinafter to verify if these stability range can be increased.

Impact of Pt addition on the Ni / GeSn systems during the solid-state reaction

As presented previously, Ni / GeSn intermetallics present low thermal stability and degrade when reaching certain annealing temperatures. Generally, the layers will agglomerate and this process can be aggravated due to Sn segregation. Once the layers agglomerated the electrical properties can be compromised. Limitations on the stability process window of the monosilicide have been also encountered. Indeed, the thermal stability of this intermetallic is low, degradation occurs due to severe agglomeration above a certain temperature and the transformation of NiSi phase into the more resistive NiSi 2 [START_REF] Adusumilli | Silicide-phase evolution and platinum redistribution during silicidation of Ni 0.95 Pt 0.05 /Si (1 0 0) specimens[END_REF][START_REF] Panciera | Direct epitaxial growth of h-Ni 2 Si by reaction of a thin Ni (10 at.% Pt) film with Si (1 0 0) substrate[END_REF][START_REF] Liu | Comparison of the thermal stability of NiSi films in Ni/Pt/(1 1 1)Si and Ni/Pt/(1 0 0)Si systems[END_REF][START_REF] Lee | In situ XRD analysis of Ni(Pt)/Si(100) reactions in low temperature regime ≤ 400[END_REF][START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF][START_REF] Demeulemeester | The influence of Pt redistribution on Ni 1-x Pt x Si growth properties[END_REF][START_REF] Bockstael | Effect of Pt addition on growth stress and thermal stress of NiSi films[END_REF]. In order to fulfill the growing demands of the microelectronic industry and continue the down-scaling of devices, research to tailor different properties of the silicide thin film increased and showed that the use of a NiPt alloy can be beneficial [START_REF] Demeulemeester | The influence of Pt redistribution on Ni 1-x Pt x Si growth properties[END_REF][START_REF] Bockstael | Effect of Pt addition on growth stress and thermal stress of NiSi films[END_REF]. The use of Ni 0.9 Pt 0.1 alloy for contacting GeSn layers was then analyzed and is presented hereinafter.

4.2.3.1 Ni 0.9 Pt 0.1 / Ge 1-x Sn x phase formation sequence

The phase sequence between a Ni (deposited either in a 200 or 300 mm tool) thin film and Ge 0.90 Sn 0.10 layer, shown in chapter III, can be summarized as follows (i) A Ni-rich phase (i.e. Ni 5 (GeSn) 3 ) was formed first, (ii) the total consumption of Ni occurred around 120 , (iii) then the Ni-rich phase dissipated and the growth of the mono-stanogermanide phase Ni(Ge 0.90 Sn 0.10 ) initiated. Overall, the same phase sequence was verified to occur for the reaction of NiCo film with a Ge 0.90 Sn 0.10 layer (see section 4.2.2.1). In order to compare these previous results to Pt-added systems, the Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 phase sequence was followed by in-situ XRD, respecting the same conditions as for previous analyses (see chapter II, section 2.1.1), the resulted XRD pattern is shown in Fig. 4.31.

Characteristic peaks of TiN (1 1 1) at 36.4°, Ni 0.9 Pt 0.1 (1 1 1) at 43.9°, Si (4 0 0), and Si/Ge/GeSn (4 0 0) around 68°were observed. At 120 , the Ni 0.9 Pt 0.1 alloy peak shifted towards lower 2θ angles, suggesting an increase in the alloy lattice parameter. Indeed, as observed in Fig. 4.32 the 2θ values shifted progressively while the intensity of the peak decreased. This behavior will be explained later in section 4.2.3.2.

At 140 , the Ni 0.9 Pt 0.1 alloy was totally consumed. The addition of 10 at.% of Pt delayed the total consumption of Ni. Indeed, comparable in-situ XRD measurements presented in chapter III, showed that the Ni layer was totally consumed at 120 . Pt redistribution during thermal reaction might explain this delay in total Ni consumption. Indeed, the creation of a Pt-enriched layer at the NiPt / Ni-rich phase interface [START_REF] Cojocaru-Mirédin | Snowplow effect and reactive diffusion in the Pt doped Ni-Si system[END_REF] and/or at grain boundaries [START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF] may act as a diffusion barrier for Ni and decreased the Ni mobility for phase formation. This phenomenon will be further described in section 4.2.3.2.

After the Ni 0.9 Pt 0.1 total consumption, two peaks appeared at 26.4°and around 87°. Based on previous results on the Ni / GeSn system presented in chapter III, these two peaks were due to the growth of the Ni-rich hexagonal phase of the type (Ni 0.9 Pt 0.1 ) 5 (Ge 0.90 Sn 0.10 ) 3 . For the case of Ni-silicides, the flux of Pt is limited in the Ni-rich phase (i.e. Ni 2 Si). This could be due to a low diffusion coefficient of Pt in Ni 2 Si and/or a low solubility of Pt in the Ni 2 Si phase [START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF][START_REF] Cojocaru-Mirédin | Snowplow effect and reactive diffusion in the Pt doped Ni-Si system[END_REF]. Pt solubility in the Ni 5 (GeSn 3 ) phase is discussed then in section 4.2.3.3. The two peaks, at 26.4°and around 87°, related to the NiPt-rich phase, shifted progressively towards higher 2θ angles as temperature increased. They disappeared at 240 . This shift, equally as for the Ni / GeSn systems, was linked to stoichiometry variations, i.e. lattice constant fluctuations of the metastable ε-Ni 5 (GeSn) 3 hexagonal phase when the GeSn concentration varied (see chapter III).

At 240 , the Ni-rich phase disappeared and a new peak emerged at 53.2°. This peak was due to the growth of the (Ni 0.9 Pt 0.1 )(Ge 0.90 Sn 0.10 ) mono-stanogermanide phase. Contrary to results obtained for the Ni / Ge 1-x Sn x systems from literature [START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Nishimura | Formation of Ni(Ge 1-x Sn x ) layers with solid-phase reaction in Ni/Ge 1-x Sn x /Ge systems[END_REF] and in-situ XRD data collected on Ni / Ge 0.90 Sn 0.10 , (see chapter III) only one out-of-plane orientation was evidenced for NiPt system with a peak around 53.2°. Similar result was obtained when Co was added (see section 4.2.2.1). Thus, Pt addition might modify the mono-stanogermanide texture. A similar phenomenon was reported for the texture of NiSi on Si (0 0 1). The addition of Pt induced an epitaxial alignment with the substrate, while the axiotaxy alignment in the layers decreased [START_REF] Detavernier | Influence of Pt addition on the texture of NiSi on Si (0 0 1)[END_REF].

When reaching 350-360 , a slight shift of the mono-stanogermanide phase peak towards higher 2θ angles was observed. It is highlighted by the circular dotted symbol in Fig. 4.31 and verified in Fig. 4.33.

The shift observed for the mono-stanogermanide phase indicated a lattice parameter reduction and was concomitant to the appearance of two other peaks at 39.2°and 42.0°. In the 345-355 temperature range, only the peak already described at 53.2°was visible.

Then at 360 , a second diffraction peak appeared at about 53.6°. As the temperature increased, the intensity of the latter increased while the original peak disappeared. This phenomenon might be related to a phase change, compounds segregation or thermal expansion and will be discussed later on in section 4.2.3.4. In order to complement the in-situ XRD results in Fig. 4.31, Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 samples obtained under the same process conditions described earlier, were ex-situ annealed for 30 s at temperatures ranging from 150 up to 550 in a rapid thermal annealing (RTA) furnace under N 2 environment. Then, θ-2θ scans were acquired under the same conditions described in chapter II, section 2.1.1. The results obtained are shown in Fig. 4.35.

From the as-deposited state and up to 200 , characteristic peaks of Ni, TiN and the signal coming from the Si and Ge/GeSn substrate were observed at 44.0°, 36.5°and around 68°, respectively. At 250 , the metal Ni 0.9 Pt 0.1 layer was consumed and the Ni(Pt)-rich phase started to form. At 300 , the peaks, related to the Ni(Pt)-rich phase shifted towards higher 2θ angles. This was due to stoichiometry variation as discussed earlier. At 350 , the Ni(Pt)-rich phase disappeared and the formation of the mono-stanogermanide phase was observed. This phase was still present for the highest annealing temperatures (i.e. 550 ). Unlike in-situ annealed samples, it was difficult to identify diffraction peaks related to the formation of transient phases for samples which underwent RTA. A peak observed at 350-400 and highlighted by a black dotted rectangle could however correspond to one of those phases (Sn, Pt x Sn y , Pt x Ge y or Ni x Sn y ). Two hypotheses can be drawn to explain the lack of such diffraction peaks at the same 2θ positions observed by in-situ XRD (i) the transient compounds being only stable in a narrow temperature range, the risk of overlooking transient phases is higher with RTA samples. (ii) The thermal budgets being different, kinetics might play a role in the observed phases. In this case, it means that, although nucleation of the compounds may occur, no diffusion could take place in order to grow the phase and make it observable by XRD. On the other hand, small diffraction peaks present in the 65-75°range could be attributed to several phases; because the broad Si/Ge/GeSn peak interferes with their analysis, their exact assignment remains unclear.

At 450 , β-Sn diffraction peaks linked to Sn segregation were visible. These kinds of peaks were not observed during in-situ XRD analyses. A similar phenomenon was observed for Ni / GeSn and NiCo / GeSn samples and was discussed in detail in chapter III. Indeed, during annealing Sn can stay incorporated in the intermetallics at temperatures beyond the Ni-Ge eutectic point or Sn melting point. Then, no peaks corresponding to Sn diffraction lines were observed at low temperatures, for in-situ XRD patterns, when heating up. Sn segregation can occur beyond this critical point, at a temperature in which Sn will be melted thus no diffraction peak could be detected. Sn diffraction lines can be detected for the ex-situ annealed samples in Fig. 4.35, because for these samples, XRD patterns were obtained once the samples reached room temperature. This allowed segregated Sn to solidify in a crystalline phase. It was evidenced previously in chapter III, that Sn segregation aggravated the morphological and electrical degradation of the layers. This phenomenon is highly detrimental for the contacts quality and reliability cannot be ensured. Additionally, as observed in chapter III, Sn coming from the GeSn layer that did not react with the Ni layer can also segregate. Pt addition delayed Sn segregation up to 450 , while for Ni / GeSn and NiCo / GeSn systems diffraction lines related to Sn segregation were observed after annealing at 350 and 400 , respectively. The reason for the delay on Sn segregation is then important to address and will be further analyzed in section 4.2.3.5.

In order to complement the results concerning NiPt / GeSn phase sequence, in-situ XRD patterns for systems with various Sn at.% content were also analyzed. The same phase sequence as the one verified for the Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 sample was obtained, from a straightforward comparison of in-situ XRD patterns. A Ni(Pt)-rich phase was formed first. Then, this phase disappeared to the benefit of the growth of mono-germanide or mono-stanogermanide phase. This phase appears to be stable up to a threshold temperature at which the peak suffered a shift towards higher 2θ values (except for the NiPt / Ge system). Additionally, peaks that could be related to either Sn segregation, Pt x Sn y , Pt x Ge y or Ni x Sn y compounds were observed (the nature of these peaks will be explained hereinafter). Differences were however observed concerning the temperature threshold for phase formation, which are summarized in Table 4.4. The increase in Sn content increased the themal stability of the Ni-rich phase (see Table 4.4) which, as consequence, delayed the formation of the (NiPt)(GeSn) phase as verified for the Ni / GeSn in chapter III and NiCo / GeSn systems in section 4.2.2.1. In order to better analyzed this variation, a focus on the low temperature area for different Ni(Pt) / Ge 1-x Sn x sample is presented in Fig. 4.36. Since the as-deposited state and up to 140 , the NiPt peak was observed around 43.9°for all the samples. The 2θ position of this peak compared to Ni 2θ position (44.5°) was shifted toward lower angular values due to Pt addition. Ni and Pt atoms can form a solid solution in the as-deposited layer [START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF]. Due to the higher atomic radii of Pt compared to Ni, 139 pm vs. 124 pm respectively, the diffraction peak can be positioned at lower 2θ values.

Pt snowplow effect

Around 120 , the NiPt peak progressively shifted from 43.8°at 50 up to 41.9°at 140 , (see Fig. This phenomenon was related to the snowplow effect. This effect consists in the rejection of an alloy element during the growth of a phase, i.e. an addition or impurity element is pushed away by a moving interface and accumulates at one of the sides [START_REF] Cojocaru-Mirédin | Snowplow effect and reactive diffusion in the Pt doped Ni-Si system[END_REF].

The Pt snowplow effect in silicides has been already observed in several cases [START_REF] Demeulemeester | On the growth kinetics of Ni(Pt) silicide thin films[END_REF][START_REF] Lee | In situ XRD analysis of Ni(Pt)/Si(100) reactions in low temperature regime ≤ 400[END_REF][START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF][START_REF] Demeulemeester | The influence of Pt redistribution on Ni 1-x Pt x Si growth properties[END_REF][START_REF] Demeulemeester | Pt redistribution during Ni(Pt) silicide formation[END_REF][START_REF] Mangelinck | Atom probe tomography of Ni silicides: First stages of reaction and redistribution of Pt[END_REF][START_REF] Putero | First silicide formed by reaction of Ni (13% Pt) films with Si (1 0 0): Nature and kinetics by in-situ X-ray reflectivity and diffraction[END_REF]. This Pt behavior might be linked to a poor solubility and limited flux of Pt in Ni 2 Si phase [START_REF] Cojocaru-Mirédin | Snowplow effect and reactive diffusion in the Pt doped Ni-Si system[END_REF][START_REF] Demeulemeester | Pt redistribution during Ni(Pt) silicide formation[END_REF]. The solubility of Pt in the Ni 2 Si is expected to be limited since the Ni 2 Si and Pt 2 Si phases do not share the same structure (Ni 2 Si has an orthorhombic structure and the structure of Pt 2 Si is rhomboedric). The solubility limits of Pt in this phase was verified to be low about 8 at.% at 800 and is estimated to be about 1 at.% at 300 by thermodynamics calculations [START_REF] Hoummada | Effect of Pt addition on Ni silicide formation at low temperature: Growth, redistribution, and solubility[END_REF]. Then, Pt segregation would be preferential to occur than the incorporation into the phases. Assuming that Ni is the fastest diffusive species to react with the Ge 0.90 Sn 0.10 layer, a difference on Pt and Ni diffusion flux during the phase sequence will enable the metal alloy lattice to become richer in Pt as the intermetallic phase grew. As Pt has a bigger atomic radii than Ni, it can increase the lattice size, shifting the diffraction NiPt peak towards lower 2θ values. Additionally, the low solubility of Pt in the Ni-rich phase (Ni 5 (GeSn) 3 ) might promote Pt segregation. Indeed, for Pt-Ge system, three Pt-rich phases are thermodynamically stable, i.e. Pt 3 Ge (monoclinic), Pt 2 Ge (cubic) and Pt 3 Ge 2 (orthorhombic). Then, Pt solubility in the hexagonal ε-Ni 5 (GeSn) 3 phase is expected to be low. This low solubility might promote Pt expulsion and accumulation towards the surface in a snowplow process as for the case of Ni(Pt)-silicides.

Pt solubility in the Ni / GeSn intermetallics

Information on the evolution of the NiPt-rich phase, was obtained thanks to the in-situ XRD analysis realized. At 140 , Ni(Pt)-rich peaks were observed at 26.4°and around 87°, for all the samples. Previously, we verified a variation of 2θ for the Ni-rich phase, when comparing GeSn-samples to pure Ge samples, when Ni metallization was used (see chapter III). Indeed, Ni 5 Ge 3 peak was located at 31.9°for the Ni / Ge sample while Addition of an alloy element (Co or Pt) to the Ni layer Ni 5 (GeSn) 3 peaks were around 31.6°for Ni / GeSn samples. This variation was earlier related to Sn incorporation and later verified by TEM-EDS analysis shown in Fig. 3.41 and 3.42. For NiPt-samples, NiPt-rich peaks were broad and diffuse, then a variation between Ge-based and GeSn-based samples was difficult to verify.

The NiPt-rich phase peak shifted towards slightly higher 2θ angles, with the annealing temperatures. This shift was observed for all the samples, as observed for Ni-samples, and was entirely related to stoichiometry variations in the ε-Ni(Pt) 5 (Ge(Sn)) 3 phase as presented in chapter III. After the Ni(Pt)-rich phase was totally consumed, the Ni(Pt)monogermanide or Ni(Pt)-monostanogermanide phase was sequentially obtained and evaluated by in-situ and ex-situ XRD analysis, see Fig. At 200 , TiN, Ni, Pt and GeSn layers were well identified together with a layer of 7 ± 1 nm consisting of the Ni-rich phase. Pt was observed to slightly diffused in a small extent of about 1-2 nm. The EDS profiles in Fig. 4.38 indicated four important zones for this sample. Zone 1 corresponded to the GeSn layer, in which Sn was homogeneously distributed. Then, in zone 2, the Ni 5 (GeSn) 3 phase was evidenced with a low quantity of Pt. Ge / Sn ratio in the GeSn layer and the Ni 5 (GeSn) 3 phase was comparable. Then, Sn solubility in the Ni-rich phase was verified as for the Ni / GeSn system (see chapter III). Zone 3 was related to the NiPt alloy which had not reacted yet and in which Pt was homogeneously distributed. Finally, zone 4 presents the TiN capping layer.

At 350 , a 34 ± 1 nm thick layer was well identified from Fig. 4.37, due to the Ni(GeSn) phase growth. Meanwhile, Pt was accumulated mostly in the top surface of this phase. The EDS profiles in Fig. 4.38 show the incidence of 5 distinctive zones or layers. Zone 1 corresponded to the unreacted GeSn layer in which Sn was homogeneously distributed. Then, zone 2 corresponded to a Sn depletion area of the Ni(GeSn) phase. Similar behavior was observed for the Ni / GeSn system (see Fig. It was then verified that the shift on the NiPt peak towards lower 2θ position was related to Pt snowplow effect. This phenomenon occurred due to faster Ni diffusion and the low solubility of Pt on the Ni-rich phase. This effect can also explain the higher thermal stability of the NiPt. The accumulation of Pt in the surface, low diffusion and incorporation in the phases allowed the alloy to remain stable up to 140 .

PtSn x compounds formation

From Table 4.4 and Fig. 4.36, a variation of the threshold temperature for the mono-(stano)germanide phase growth (which corresponds to the Ni-rich consumption temperature in Table 4.4) when the Sn content varied was observed. This phase was obtained at 165 for the NiPt / Ge system. This temperature increased up to 210 , 240 and 245 , for NiPt / Ge 0.94 Sn 0.06 , NiPt / Ge 0.90 Sn 0.10 and NiPt / Ge 0.85 Sn 0.15 systems, respectively. For the Ni / GeSn systems (see chapter III), it was verified that the Sn behavior during the solid-state reaction had an impact on the diffusion mechanisms during the phase formation. Sn distribution and segregation along grain boundaries (GB) delayed the Ni(GeSn) phase growth. Sn is expected to act as a diffusion barrier, more precisely as a "stuffed barrier", which reduced the Ni mobility, delaying the mono-stanogermanide phase growth. Then, a similar effect is expected to occur for NiPt / GeSn samples. For this case however, both Pt and Sn accumulation in GB might hamper Ni diffusion.

For the case of silicides, high concentration and distribution of Pt was verified to occur in different areas of the silicide first a high concentration of Pt was found in the Ni(Pt) layer that has not reacted during the isothermal treatment due to snowplow effect, second at Ni 2 Si grain boundaries and third at the interfaces [START_REF] Imbert | Impact of surface preparation on nickel-platinum alloy silicide phase formation[END_REF][START_REF] Hoummada | Composition measurement of the Ni-silicide transient phase by atom probe tomography[END_REF][START_REF] Mangelinck | Chapter 9 -The Growth of Silicides and Germanides[END_REF]. Then Pt accumulation in the Ni-rich GB could hamper Ni diffusion and delay the Ni(GeSn) formation. The fact that the mono-stanogermanide phase was further delayed with a further increase of Sn content implies that the phenomenon has to be related to the concomitant effect of Pt and Sn. In a first instance, Pt is expected to mostly accumulate at the surface and, while the intermetallics growth, Pt can segregate either at GB or at the NiPt / Ni 5 (GeSn) 3 interface. This behavior could hamper to some extent Ni diffusion. Meanwhile the reaction continues, Sn segregation can also occur. Indeed, a depletion Sn zone occurred in the Ni(GeSn) phase while a slight increase was observed near the surface (see Fig. 4.38). Then a further accumulation of Sn at GB could also reduce Ni mobility and delay the mono-stanogermanide growth when Sn content increase.

On the other hand, Pt and Sn accumulation could further modify the phase sequence due to the growth of new phases. Indeed, a shift was observed in the mono-stanogermanide phase after a certain threshold temperature that also varied with Sn content (see Table 4.4). This shift was accompanied by the incidence of new diffraction peaks. Indeed, for the Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 system, a peak at 39.2°appeared from 350 up to 495 , then another peak grew from 360 up to 455 at 42.0°and finally these ones were consumed to the benefit of one peak positioned at 42.6°since 495 . For this quaternary system different possibilities could explain the diffraction peaks encountered Sn segregation and the formation of Pt x Sn y or Ni x Sn y or Pt x Ge y compounds. In order to understand these peaks and complete the characterization of the phases in the 350-600 temperature range, ex-situ in-plane RSM analyses were performed for the NiPt / Ge 0.90 Sn 0.10 system. Two samples were annealed, respecting the annealing conditions used for in-situ XRD measurements (same ramp up rate and scanning time), up to 400 and 600 , respectively. Integrated patterns obtained from in-plane RSM are shown in Fig. 4.39.

The analysis of in-plane RSM integrated patterns is highly instructive. All diffraction peaks were indexed and common peaks were highlighted for both samples. First, diffraction peaks at 36.5°, 42.3°, and 62°were, respectively, related to TiN (1 1 1), (2 0 0), and (2 2 0) in-plane orientations. Then, diffraction peaks at 45.2°and 66.6°were linked to Ge/GeSn (2 2 0) and (4 0 0) in-plane orientations. Finally, numerous peaks coming from the Ni(GeSn) phase have also been indexed, highlighting the fact that the monostanogermanide phase has numerous in-plane orientations. However, it exhibited a strong out-of-plane texture (only one orientation was observed in Fig. 4.31 for this phase). The remaining diffraction peaks were attributed to β-Sn. Only the most intense (2 0 0) orientation located at 30.6°was clearly detected at 400 whereas numerous other orientations are present at 600 .

Diffraction peaks related to β-Sn would not properly fit the diffracted peaks obtained at 39.2°, 42.0°and 42.6°during the heating up. On the other hand, segregated Sn is expected to be in a molten state at these temperatures during the in-situ heating up. It is only when the samples were cooled down that Sn was detected. This phase can be then ruled out to fit the peaks at 39.2°, 42.0°and 42.6°.

In a previous study, Wang et al. showed the formation of Pt x (GeSn) y phases from 350-450 when depositing Ni-Pt (by co-sputtering in a magnetron sputtering tool) on NiSn [START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF]. In the current study, the diffraction peaks observed in the 350-600 temperature range during in-situ XRD analyses or at 400 and 600 during in-plane RSM studies were not attributed to Pt x (GeSn) y or Ni x Sn y phases but to PtSn x compounds.

First, in the present study, the diffraction peaks observed at 39.2°, 42.0°and 42.6°during the in-situ XRD analyses (see Fig. 4.31) were not detected during comparable in-situ XRD measurements on Ni 0.9 Pt 0.1 / Ge or Ni / Ge 1-x Sn x systems. Thus, the formation of phases in this range of temperatures is due to the concomitant presence of Pt and Sn. Nevertheless, given the study of Wang et al. and Liu et al., all Pt-Sn, Pt-Ge and Ni-Sn systems were taken into account for diffraction peaks assignment. At 39.2°, the diffraction peak can be linked to the following phases Pt 3 Sn, PtSn 4 , PtSn 2 or Ni 3 Sn, Ni 3 Sn 4 or Pt 3 Ge 2 . At 42.0°, the diffraction peak can be related to PtSn, Pt 2 Sn 3 , PtSn 3 , or Pt 3 Ge 2 phases. Finally at 42.6°, the diffraction peak can be linked to the following phases PtSn 4 , PtSn 3 or NiSn, Ni 3 Sn, Ni 3 Sn 2 , Ni 3 Sn 4 or PtGe 2 , PtGe, Pt 2 Ge 3 . PtGe phase is less probable to occur since the difference between the theoretical position of the most intense diffraction peaks and the experimental one is higher than for the other phases. The Pt 3 Ge 2 phase has a very low range of thermal stability and its existence domain is quite short [START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]. Thus, the probability to observe this phase in the 350-600 temperature range is weak. To go further, the enthalpies of formation for Pt-Sn, Ni-Sn and Pt-Ge systems were gathered from literature data and are presented in Except for PtSn 4 and PtSn 3 (metastable) phases, the enthalpies of formation for the PtSn x phases are lower than the ones of the Pt-Ge or Ni-Sn systems. Thus, as far as thermodynamic is concerned, the formation of PtSn x compounds is more favorable. This hypothesis was verified thanks to TEM images coupled to EELS maps for atom identification in Fig. 4.40. The analysis of the in-plane RSM integrated patterns for samples annealed at 400 and 600 presented in Fig. 4.39 can be then helpful to cautiously determine the exact nature of the phases observed. The diffraction peaks not assigned to TiN, Ni(GeSn) or Ge/GeSn phases were related to β-Sn and PtSn x phases, namely PtSn 4 , PtSn 3 , PtSn 2 , and, to a lesser degree, Pt 2 Sn 3 . If some of the peaks are common to both annealing temperatures, a trend was identified. At 400 , the main phase is expected to be PtSn 4 , whereas at 600 , the metastable PtSn 3 phase is predominantly present. Indeed, from Table 4.4 and Fig. 4.36 it can be highlighted that the threshold temperature for Ni(GeSn) shift and the consequent formation of PtSn x compounds decreased with Sn content (the full in-situ XRD patterns can be consulted in Appendix D). Indeed, peaks related to PtSn x were observed since 405, 350 and 305 for Ni 0.9 Pt 0.1 / Ge 0.94 Sn 0.06 , Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 and Ni 0.9 Pt 0.1 / Ge 0.85 Sn 0.15 systems, respectively. Pt was expected to accumulate in the Ni-rich phase grain boundaries, while Sn was observed to be homogeneously distributed in the Ni(GeSn) phase (see Fig. 4.37 and 4.38). However, due to the high metastability of these kind of intermetallics, Sn segregated after reaching certain annealing temperatures. This segregation was observed to occur preferentially at GB. Then, Sn was available to form PtSn x compounds with the accumulation of Pt at GB, and can evolve from a Snrich compound to a Pt-rich one due to further Sn segregation to form β-Sn. Then as observed for the Ni / GeSn system in chapter III, Sn segregation is expected to occur at lower annealing temperatures for higher Sn content layers. Then PtSn x compounds can be formed at lower temperatures as well.

Impact of Pt addition on Sn segregation

As for the case of the NiCo / GeSn system (see section 4.2.2.2), Pt addition seemed to delay Sn segregation to higher temperatures. XRD analyses were performed on samples annealed by RTA were presented in Fig. 4. [START_REF] Mcguire | Characterization of semiconductor maerials: Principles and Methods[END_REF] The patterns highlight the impact of the alloying element on the Sn segregation threshold. Thus, for each system, namely Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 metallizations on Ge 0.90 Sn 0.10 layers, the Sn segregation threshold temperatures were extracted from the XRD patterns, the latter are plotted in Fig. 4.41. As explained earlier, Sn segregation threshold for the Ni / GeSn system was ranging between 300 and 350 . In the same way, the Sn segregation threshold for the Ni 0.9 Pt 0.1 / Ge 0.9 Sn 0.1 systems was ranging between 400 and 450 , respectively. Thus, the addition of Pt meaningfully delayed Sn segregation. A further discussion and com-Addition of an alloy element (Co or Pt) to the Ni layer parison with NiCo / GeSn system is given in section 4.2.4.

NiPt / GeSn morphological evolution with temperature

As mentioned in chapter III, Ni / GeSn intermetallics exhibit a poor thermal stability linked to the precipitation of Sn atoms and the agglomeration of the Ni(GeSn) layer. This phenomenon was observed to be Sn content dependent and enhanced for GeSn layers with high Sn concentrations. Issues for contacting high Sn content stacks might arise and process steps might need further adjustments. Pt addition is positioned then as a solution in order to increase the thermal stability of these intermetallics. As observed earlier (section 4.2.3.5), Sn segregation can be delayed. Atomic force microscopy images were acquired in order to verify surface stability, both when increasing Sn content and the annealing temperature. Fig. 4.42 shows the AFM images obtained for samples annealed at various temperatures by RTA for 30 s in N 2 atmosphere for NiPt / Ge and NiPt / GeSn systems with various Sn content. Whatever the system studied and at low temperature (i.e. for T < 350 ), samples exhibited a cross-hatch pattern along the 1 1 0 crystallographic direction, as observed in chapter III, section 3.2.1 when analyzing the epitaxial GeSn layers. The cross-hatch pattern was preserved up to 400 for the NiPt / Ge system. For these samples, some variation at the sub-micron scale of the surface was evidenced at temperatures as low as 250 . This could be related to the growth of the intermetallics. Indeed, for this system, Chapter IV. Alternatives to improve the Ni / GeSn thermal stability 4.41. This threshold was reduced down to 400 for the NiPt / Ge 0.85 Sn 0.15 . Indeed, this sample had the highest Sn content and as a consequence the range of stability was reduced.

Briefly, this section covered the systematic and comprehensive analysis concerning the impact of Pt addition on the Ni / GeSn phase sequence, the morphological and the electrical properties evolution during the solid-state reaction. It was encountered that the phase sequence remained comparable to the Ni / GeSn reference a Ni-rich phase was formed first followed by the Ni(GeSn) phase. The main difference concerned the formation of PtSn x phases which was verified by both in-plane RSM and TEM / EELS analyses.

Pt also played a role in Ni diffusion due to the snowplow effect and Pt distribution around grain boundaries. Pt solubility in both the Ni-rich and the Ni(GeSn) phases was low. The low solubility and the accumulation at GB can be one of the reasons of the temperature delay for the Ni-rich phase formation compared to the Ni / GeSn system. On the other hand, Pt delayed Sn segregation, grain boundary grooving and agglomeration. This enhancement of the morphological properties benefited directly the stability of the R sh values.

Despite the complexity of the phase formation sequence due to PtSn x compounds formation, the addition of Pt in the Ni-stanogermanide was clearly beneficial and allowed to extend the process window in which satisfactory contacts, in terms of morphology, roughness and electrical properties, can be obtained. Ni 0.9 Pt 0.1 CMOS compatible alloy is then suited to contact GeSn-base devices and is envisioned as an viable option to maintain the thermal stability of the Ni / GeSn layers.

Comparison of Co and Pt effect on the Ni / GeSn system

In order to contrast the results obtained in this section and summarize the main findings, a global comparison on the use of NiCo and NiPt alloys to enhance the properties of the Ni / GeSn layers is addressed hereinafter in terms of phase sequence, Sn segregation and evolution of the morphological and the electrical properties during the solid-state reaction.

Impact on the phase sequence

The study of the solid-state reaction for the Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems highlighted that, even if the overall phase sequence was the same, discrepancies existed. The Fig. Concerning the Ni consumption, the addition of 10 at.% of Co did not impact the temperature at which the Ni was totally consumed, indeed the Ni full consumption was observed at 120 for both metallizations. On the other hand, the addition of 10 at.% of Pt pushed up the Ni total consumption up to 140 . Pt redistribution during thermal reaction, in particular the Pt snowplow effect, can explain this delay in the total Ni consumption.

The formation of the Ni-rich phase was impacted by the nature of the alloying element. For Ni metallization, the Ni 5 (GeSn) 3 phase appeared at 85 and was totally consumed at 210 . With the addition of Co, the (NiCo) 5 (GeSn) 3 phase was observed since the as-deposited state. This phase was totally consumed at 210 . On the other hand, the addition of Pt delayed the Ni-rich phase growth up to 140 and its total consumption up to 240 . Similar trends were observed for the temperatures at which the mono-stanogermanide phase appeared. Co addition slightly lowered this temperature (200 vs. 210 for Ni) whereas the Pt addition pushed this temperature up to 240 .

The impact of the alloying element on the formation temperatures of the intermetallics can be discussed according to various lines. From a solubility point of view, elements forming germanides that are soluble in Ni The NiGe phase can coexist with the CoGe, Co 5 Ge 7 and CoGe 2 phases on one hand and with Pt 2 Ge, Pt 3 Ge 2 , PtGe, Pt 2 Ge 3 and PtGe 2 on the other hand [START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]. NiGe being orthorhombic, all the orthorhombic phases might be soluble in the NiGe phase. For the Pt / Ge system, all the intermetallic phases, except Pt 2 Ge which is hexagonal, are orthorhombic. For the Co / Ge system, only CoGe 2 is orthorhombic, CoGe being monoclinic and Co 5 Ge 7 tetragonal. Pt incorporation in the phase should then be possible. However, a low incorporation of Pt on the Ni(GeSn) was observed (see Fig. 4.37 and 4.38).

Concerning the Ni-rich phase, the Ni 5 Ge 3 coexists only with the transition metals. The fact that the Co and Co 5 Ge 3 have an hexagonal structure, like Ni 5 Ge 3 , and that Pt has a cubic structure and was observed to have a low solubility in the Ni / GeSn intermetallics might explain why the addition of Co decrease the energy threshold for the Ni-rich phase and Pt addition delayed it to higher temperatures.

For the mono-germanide phase, if the effect of the nucleation barrier is not self-evident regarding the nature of the alloying element, one should consider surface / interface effects. The bond strengths for Co-Ge, Ni-Ge and Pt-Ge are 230, 290 and 400-500 kJ.mol -1 , respectively [START_REF] Gingerich | Experimental and predicted stability of diatomic metals and metallic clusters[END_REF][START_REF] Shim | Electronic states and nature of bonding of the molecule NiGe by all electron ab initio Hartree-Fock (HF) and configuration interaction (CI) calculations and mass spectrometric equilibrium experiments[END_REF][START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF]. If the bonding between atoms is weaker (case of Co), the surface, interface and grain boundary diffusions are more easy. As a consequence, the atomic diffusion will be faster and the formation of the intermetallic phases can be shift towards lower temperatures with the addition of cobalt. The opposite reasoning can be made with the addition of platinum. In addition, concerning the platinum effect, it can be observed in Fig. 4.31 and 4.32 that at around 120 , the NiPt alloy peak shifted towards lower 2θ angles, suggesting an increase of the alloy lattice parameter. The phenomenon is known as Pt snowplow and was already discussed. This phenomenon might explain the latency observed for the appearance of the Ni-rich phase with the addition of Pt as explained earlier.

Sn segregation thresholds

The GeSn layers being realized by RP-CVD under kinetic conditions far from equilibrium, they are metastable. Indeed, GeSn is known for its poor thermal stability and Sn has been reported to segregate at the surface with increasing annealing temperature [START_REF] Takase | Behavior of Sn atoms in GeSn thin films during thermal annealing: Ex-situ and in-situ observations[END_REF][START_REF] Tsukamoto | Investigation of Sn surface segregation during GeSn epitaxial growth by Auger electron spectroscopy and energy dispersive X-ray spectroscopy[END_REF], which was verified in chapter III. Sn segregation is linked to the formation of Sn clusters / precipitates in the GeSn layers and the subsequent migration of atoms towards the surface [START_REF] Takase | Behavior of Sn atoms in GeSn thin films during thermal annealing: Ex-situ and in-situ observations[END_REF][START_REF] Tsukamoto | Investigation of Sn surface segregation during GeSn epitaxial growth by Auger electron spectroscopy and energy dispersive X-ray spectroscopy[END_REF][START_REF] Groiss | Free-running Sn precipitates: an efficient phase separation mechanism for metastable Ge 1-x Sn x epilayers[END_REF]. Even if reducing the density of point defects limits the initial formation of Sn clusters / precipitates, Sn segregation was mainly governed by the migration / diffusion of atoms. Thus, the addition of alloying elements might (i) affect the diffusion of Sn atoms and (ii) impact the surface / interface energies involved in Sn segregation and distribution mechanisms. As reported earlier, the addition of Co or Pt could lead to the formation of Co x Sn y compounds and was verified to enable the formation of PtSn x compounds. After the formation of Sn clusters / precipitates in the GeSn layers, Sn atoms diffuse towards the surface, the diffusion occurred within the grains or at grain boundaries of the Ni(GeSn) phase. In the presence of alloying elements, such as Co or Pt, the Sn diffusion might be slowed down by the formation of Co x Sn y / PtSn x compounds and thus, the Sn surface segregation might be delayed.

A second hypothesis concerns the modification of the surface / interface energies by the alloying elements. It is well known that the addition of alloying elements can impact the texture of silicides and germanides [START_REF] Schutter | Phase formation and texture of thin nickel germanides on Ge (0 0 1) and Ge (1 1 1)[END_REF]. As can be seen in Fig. 3.8, 4.21 and 4.31, the number of out-of-plane orientations for the mono-stanogermanide phase decreases with the addition of an alloying element from 3 various orientations with pure Ni to only one orientation with the addition of Co or Pt. The formation of highly textured films has already been reported for Pt addition in Ni-based silicides [START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Detavernier | Influence of Pt addition on the texture of NiSi on Si (0 0 1)[END_REF]. They are expected to have a significantly lower interfacial energy compared to polycrystalline films. Thus, the formation of more textured mono-stanogermanide films by the addition of an alloying element can contribute to lower the surface / interface energies and have a positive impact on Sn segregation.

Finally, Sn segregation has also been linked to the strain at the interface [START_REF] Takeuchi | Tensile strained Ge layers on strain-relaxed Ge 1-x Sn x / virtual Ge substrates[END_REF]. The covalent radius of Ni (124 pm) is lower than the ones of Co (126 pm) and Pt (139 pm), the addition of these elements might modify the local strain and positively impact the Sn segregation threshold. As can be seen in Fig. 4.47, the bigger the element radius is, the higher the Sn segregation threshold. Meanwhile, the NiPt / Ge 0.90 Sn 0.10 layer remained smooth and continous. For these kinds of intermetallics, agglomeration starts with the grain boundary grooving and results in agglomeration, grains separation and island formation. The driven force for a film to agglomerate is the total surface / interface energy of the intermetallics and its substrate [START_REF] Gambino | Silicides and ohmic contacts[END_REF]. The impact of Pt on Ni stanogermanide is thus consistent with results obtained for Ni(Pt) / Si and Ni(Pt) / Ge systems [START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF][START_REF] Zhang | Improvement of thermal stability of Ni germanide using a NiPt(1%) alloy on Ge-on-Si substrate for nanoscale Ge MOSFETs[END_REF][START_REF] Kang | Thermally robust Ni germanide technology using cosputtering of Ni and Pt for high-performance nanoscale Ge MOSFETs[END_REF]. NiGe and PtGe have the same orthorhombic crystalline structure. Thus, Pt could be soluble in the monogermanide phase, i.e. NiGe. Even if Pt solubility was observed to be low, some diffusion in the upper part of the Ni(GeSn) intermetallic was evidenced in Fig. 4.37. Pt-Ge bond energy (400-500 kJ/mol) is higher than the NiGe bond energy (290 kJ/mol) [START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF]. As a result, Pt will reduce the breakdown rates of chemical bonds which will reduce the grain boundary grooving development and could explain the morphological stability improvement at high temperatures. This assumption on the grain boundary grooving kinetics is endorsed by AFM pictures. For the Ni / GeSn system, agglomeration is too fast to observe grain boundary grooving. On the other hand, due to Pt addition, the groove development being slower, the phenomenon is observed at 350 and 400 for the NiPt / GeSn system, see Fig. 4.48.

On the other hand, as for the case of carbon accumulation at grain boundaries described in section 4.2.2.3, Pt accumulation in these sites could influence grain-boundary energy and grain boundary diffusion. This could enhance the grain boundary cohesion and thus lower the grain-boundary energy delaying finally the agglomeration process. It has also been mentioned that the addition of Pt in Ni(SiGe) germanosilicide might suppress Ge out diffusion due to the high melting temperature of Pt and to the stronger Pt-Ge bonds compared to NiGe ones [START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF]. The suppression or limitation of Ge out diffusion could then be beneficial for morphology stability.

Additionally, models proposed to explain the mechanisms of agglomeration in thin films also state that the grain boundary grooving which is the very first step of agglomeration is linked to the grain size and that a reduction of the grain size might reduce agglomeration [START_REF] Srolovitz | Capillary instabilities in thin films[END_REF][START_REF] Miller | The instability of polycrystalline thin films: Experiment and theory[END_REF]. This point was verified in section 4.2.2.3 for PAI samples. The full width at half maximum (FWHM) of the Ni(Pt) mono-stanogermanide from various XRD patterns were obtained during in-situ measurements (Fig. 4.31) and compared to the ones gathered for Ni and NiCo systems under the same conditions. Results are listed in Table . 4.6. According to the Scherrer equation, the crystallite size is inversely proportional to the FWHM [START_REF]Elements of X-ray diffraction[END_REF]. All the FWHM values for Ni(GeSn) phases, in which Pt was added, were higher than the ones obtained for Ni and NiCo systems. This means that the grain size of the mono-stanogermanide was reduced due to Pt addition, which was clearly beneficial to delay agglomeration.

Electrical properties

In addition to compare the different systems analyzed, the R sh evolution with temperature was obtained for Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems. The results are presented in Fig. 4.50.

Whatever the system observed, from the as-deposited state and up to 200 , the sheet resistance values remained constant at around 28-35 Ω/sq. A slight increase was observed at 250 which was related to the Ni-rich phase growth. At 350 , minimum values were reached for all the systems 23.1, 16.3 and 10.4 Ω/sq for Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.9 Sn 0.1 systems respectively. The differences are significant. Therefore, the mono-stanogermanide layer obtained at 350 with the addition of Pt as an alloying element was less resistive than the one obtained with Co and even less resistive than the one without any alloying element (pure Ni). Beyond 350 and for Ni and Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 systems, a strong increase of the R sh values was observed. This behavior was directly linked to the agglomeration of the Ni(GeSn) layer and to Sn seg- regation. Thus, as observed for the morphology, the addition of Co did not significantly impact the electrical properties. The decrease of the R sh values at high temperature observed for the Ni 0.9 Co 0.1 system might be linked to the formation of Co x Sn y or Co x Ge y compounds described in section 4.2.2.4. On the other hand, the addition of Pt as alloy element enhanced the process window in which the R sh values remain low. Indeed, a plateau was observed between 350 and 400 and for 450 , the increase was relatively low. Beyond 450 , the sheet resistance increased and followed the trends and values observed for the Ni / Ge 0.90 Sn 0.10 system. The positive effect of Pt addition on surface morphological thermal stability and the delay on Sn segregation positively impact the electrical properties. The Ni(GeSn) morphological and electrical degradation was verified to occur due to two phenomena first, the agglomeration of the layer and second, Sn segregation. Both occurred when the annealing temperatures increased. However, thermal annealing is an unavoidable step when elaborating electronic or opto-electronic devices. Indeed, annealing is one critical process step used, among others, to achieve targeted device performances. During the period of traditional Moore's law scaling, the thermal budget of key annealing steps, typically for junction, gate and contact formation modules, had to be continuously reduced to find the best trade-off between activation, controlled species diffusion, crystallization quality and defects removal [START_REF] Huet | Laser thermal annealing for low thermal budget applications: From contact formation to material modification[END_REF]. Annealing Ni / GeSn stacks allows the formation of the Ni(GeSn) less resistive phase, i.e. the most interesting phase to contact GeSn-based devices.

Laser annealing technique was considered as a thermal process when the use of an intense laser pulse of nanosecond duration (from a ruby laser) was verified to be able to anneal the lattice of silicon [START_REF] Huet | Inducing and probing non-thermal transitions in semiconductors using femtosecond laser pulses[END_REF]. The interest in laser annealing process grew since 1980's, initially, to achieve ultrafast heating by intense laser or electron beams, which can melt and recrystallize, in a few hundred nanoseconds, semiconductor layers. This rapid solidification of semiconductors allowed to discover novel regimes of phase formation and has served as a means of removing implantation damage. It was this latter aspect that importantly stimulated the field [START_REF] Poate | Laser Annealing of Semiconductors[END_REF]. For CMOS production, the potential benefits of using laser annealing are large (i) highly localized dopant activation above the solid phase equilibrium limit, (ii) near-perfect liquid-phase induced recrystallization and (iii) pattern variations at transistor level. Currently, this technique is on the verge of being integrated in key annealing process steps in next generation CMOS, from contact formation to material modification and 3D sequential integration [START_REF] Huet | Laser thermal annealing for low thermal budget applications: From contact formation to material modification[END_REF].

Since the 1980's, studies considering the use of laser pulses to induce thin-film reactions were evocated [START_REF] Poate | Mechanisms of laser annealing[END_REF][START_REF] Liau | Laser annealing for solid-phase thin-film reactions[END_REF][START_REF] Wittmer | The thermal stability of laser-produced silicide layers[END_REF][START_REF] Lue | Formation of nickel and palladium silicides by a short-pulse light-flash and its application in the metallization of solar cells[END_REF][START_REF] D'anna | Metal silicides with energetic pulses[END_REF]. These studies allowed to predict the thicknesses of resultant thin-films for M-Silicide compounds formation and compare those results to existing furnace-annealing data. Poate et al. and D'Anna et al. showed that metal-Si reaction by laser irradiation was possible, which can be of high interest for contact elaboration [START_REF] Poate | Mechanisms of laser annealing[END_REF][START_REF] D'anna | Metal silicides with energetic pulses[END_REF][START_REF] Poate | Laser induced reactions of platinum and other metal films with silicon[END_REF]. Poate et al. indicated that the reacted layers were formed by nucleation and growth from a greatly super-cooled liquid instead of solid-state diffusion [START_REF] Poate | Laser induced reactions of platinum and other metal films with silicon[END_REF]. Meanwhile D'Anna et al. observed that reactions occurred whenever the energy fluence on the sample produces a sufficient heating for the interface temperature to exceed the lowest eutectic temperature of the binary system [START_REF] D'anna | Metal silicides with energetic pulses[END_REF]. Most of current uses of laser annealing for contact formation (for group IV materials) is focalized on silicides. In general, this technique was useful to (i) obtain ohmic contacts with proper electrical properties [START_REF] Silva | Formation of epitaxial Ti-Si-C Ohmic contact on 4H-SiC C face using pulsed-laser annealing[END_REF][START_REF] Wu | Formation of Ohmic contacts on laser irradiated n-type 6H-SiC without thermal annealing[END_REF][START_REF] Rascunà | Morphological and electrical properties of Nickel based Ohmic contacts formed by laser annealing process on ntype 4H-SiC[END_REF][START_REF] Badala | Ni/4H-SiC interaction and silicide formation under excimer laser annealing for ohmic contact[END_REF], (ii) improve the thermal stability of silicides [START_REF] Venkataraman | Formation of interface silicides at room temperature in pulsed laser deposited Ti thin films on Si (1 0 0)[END_REF], (iii) increase interfacial morphological stability [START_REF] Chen | Nickel silicide formation using pulsed laser annealing for nmosfet performance improvement[END_REF][START_REF] Lee | Multiple-pulse laser thermal annealing for the formation of Co-silicided junction[END_REF][START_REF] Setiawan | Laser-induced Ni(Pt) germanosilicide formation through a self-limiting melting phenomenon on Si 1-x Ge x /Si heterostructure[END_REF] and even (iv) modify the intermetallic texture [START_REF] Setiawan | Laser-induced Ni(Pt) germanosilicide formation through a self-limiting melting phenomenon on Si 1-x Ge x /Si heterostructure[END_REF][START_REF] Feautrier | Impact of laser anneal on NiPt silicide texture and chemical composition[END_REF]. On the other hand, the use of laser annealing was shown to produce uniform Ni / Ge contacts with a remarkably smooth substrate interface [START_REF] Huet | Laser thermal anneal formation of atomically-flat low-resistive germanide contacts[END_REF]. For these reasons, laser annealing is envisionned as a novel technique that can be used to properly develop Ni / GeSn contacts with low thermal budgets.

Ni / Ge and Ni / GeSn thermal regimes using laser pulse as annealing method

Numerical simulations were performed using a TCAD package, called LASSE Innovation and Application Booster (LIAB) which has been deeply explained by Huet et al [START_REF] Huet | Pulsed laser annealing for advanced technology nodes: Modeling and calibration[END_REF], to determine the melting threshold of TiN / Ni / Ge samples as presented in Fig. 4.51. Indeed, TiN / Ni / Ge stacks laser annealing was simulated from 0.1 to 0.8 J/cm 2 . The melt onset was expected between 0.40-0.45 J/cm 2 , at which a plateau in the curves was observed. At these values, the expected temperature of the stack should be close to the Ge melting temperature (around 940 ). According to these simulations, morphological variation onset (at the surface level) corresponded to the Ge melt threshold (Fig. 4.51). Then, the interesting energy density range to be explored seemed to be located between 0.1 and 0.6 J/cm 2 . Laser annealing technique can be beneficial for applications in which extremely fast, high temperature and localized annealing are needed. In order to use laser annealing technique to elaborate reliable Ni(GeSn) contacts, Ni / Ge and Ni / Ge 0.90 Sn 0.10 (referred as Ni / GeSn for this section) stacks were elaborated following the steps described in chapter III, section 3.2.1 (TiN and Ni deposition at room temperature in a 300 mm tool), (see Fig. 4.52 (a)). Then, UV-NLA was performed on both kinds of samples with a SCREEN-LT3100 tool based on a XeCl excimer laser (308 nm wavelength) with a pulse duration of 160 ± 2 ns. 15 x 15 mm 2 fields were irradiated with laser energy densities ranging from 0.08 up to 0.66 J/cm 2 with steps of 0.02 J/cm 2 . Single pulses were performed with standard laser conditions in 69 positions and were specified in Fig. 4.52 (b). Some locations out of the R sh mapping were used for melt identification (identified as green squares in Fig. 4.52 (b)) while cumulative pulses with the same energy (3 and 10 pulses) were also performed. These specific ranges of energy densities were chosen to investigate the annealing regime in which the Ge melting threshold was achieved, (see simulations in Fig. 4.51). The impact of the laser annealing was verified in terms of phase sequence, morphological and electrical properties evolution during the solid-state reaction, thanks to ex-situ XRD patterns, AFM and SEM images, and R sh evolution analyses, respectively. The results are presented hereinafter. During the laser annealing, in-situ time-resolved reflectivity (TRR) at 635 nm were acquired to detect the melting threshold and estimate the melting duration. For this, surface reflectivity evolution as a function of time was recorded for each laser shot at different energy densities (ED), for samples treated with one pulse of 160 ± 2 ns. Then, all profiles were gathered to obtain the maps shown in Fig. 4.53. A color variation on the TRR intensity implies a change on the morphological or structural properties of the samples. Thanks to this technique, the melting threshold was identified when a variation of signal intensity, proportional to the surface reflectivity, occurred. The dotted white lines in each pattern highlighted the critical energies for each sample (see Fig. 4.53). For Ni / Ge samples, a slight decrease on refletometry for ED > 0.4 J/cm 2 was observed, which was more clearly for ED > 0.46 J/cm 2 . Meanwhile for the Ni / GeSn system, the decrease was observed at ED > 0.36 J/cm 2 and more clearly visible for ED > 0.40 J/cm 2 . Then, to further verify the onset energy for reflectometry variation, the average of TRR intensity (before and after the pulse shot for 30 samples) was acquired for each sample. The evolution is presented in Fig. This analysis confirmed that the TRR signal was reduced in two steps first a slight decrease can be noticed (ED > 0.42 J/cm 2 for Ni / Ge system and ED > 0.36 J/cm 2 for Ni / GeSn system), then a clear decrease follows (ED > 0.56 J/cm 2 for the Ni / Ge system and ED > 0.52 J/cm 2 for the Ni / GeSn system). The threshold energy to reach the melt state was lower for Ni / GeSn samples. This can be related to intrinsic properties of the material and the physical interactions between the irradiate energy and the samples. In general, when the irradiated beam strikes the surface, part of the laser energy is transmitted through the layers while the rest is reflected. The radiation energy which is transferred into the layers, passes to the electrons and the electrons transfer their energy to the lattice on the time scale of the excitation. The electrons and the lattice remain in equilibrium and the laser simply heats the solid to the melting temperature within the duration of the laser pulse [START_REF] Huet | Inducing and probing non-thermal transitions in semiconductors using femtosecond laser pulses[END_REF]. The capacity to reflect this energy will thus vary from one sample to another depending on the intrinsic characteristics of the material. Indeed, samples with lower reflection %, at the laser wavelength studied, can absorb a higher amount of photons transported by the beam. Then the thermalization of the lattice is higher for a same value of energy density and higher temperatures might be reach at lower ED. For the two kinds of stacks analyzed (Ni / Ge and Ni / GeSn) the important difference is the presence of a 60 nm thick Ge 0.90 Sn 0.10 layer. The evolution of reflection % at different wavelengths for both kinds of samples was then acquired and is presented in Fig. 4.55. Ni / GeSn samples presented lower reflection % for the wavelength range between 280-320 nm of wavelength (laser anneals were performed at 308 nm). The absorption % is then greater for these stacks at lower ED. These can be related to Sn addition which will increase the metallic characteristic of the layer, increasing their energy absorption. It was observed previously, (see chapter III and chapter IV), that Sn addition actually delayed Ni(GeSn) formation towards higher temperatures. This tecnique could then be suited to form the Ni(GeSn) phase using low thermal budgets.

Ni / Ge and Ni / GeSn phase sequence

The ultimate goal of annealing a metal / semiconductor system is to promote its evolution towards a more stable thermodynamic condition. This annealing allows diffusion of the atoms to take place and the intermetallic phases to be formed. For the case of Nigermanides or Ni-stanogermanides, the NiGe or Ni(GeSn) phase corresponds to the less resistive phase, then the most suitable for contacting GeSn-based devices. For this reason, it is important to verify if the laser annealing technique is well suited to properly form the mono-(stano)germanide phase. Ex-situ XRD analysis were then realized, following the parameters described in chapter II, section 2.1.1, for either Ni / Ge and Ni / GeSn In general, the same phases were observed for the Ni / GeSn system. At 0.32 J/cm 2 , Ni, TiN and Ni 5 (GeSn) 3 peaks were observed. A broad low-intensity peak around 32.3°w as also present. This could be related to the growth of the Ni(GeSn) phase. The peak related to this phase was latter well evidenced at 0.48 and 0.64 J/cm 2 . In addition, a peak was observed at 26.0°for the highest ED analyzed (i.e. 0.64 J/cm 2 ).

A peak at this angular position cannot be related to Sn segregation. The fast annealing at high temperatures and the high solidification rates that this technique allows might allow the formation of metastable phases that have not been achieved by rapid thermal annealing techniques [START_REF] Poate | Laser Annealing of Semiconductors[END_REF]. For these reasons, the formation of Ni x Sn y compounds can be considered. From the ICDD PDF files analysis concerning various Ni x Sn y compounds, the Ni 3 Sn 4 phase could fit the peak observed. On the other hand, Lim et al. and Huet et al. reported the growth of a metastable epitaxial nickel digermanide (NiGe 2 ) phase, formed by laser annealing of Ni on (1 0 0) germanium-on-silicon substrates, when using 10 laser pulses at 0.3 J/cm 2 for 23 ns each at a wavelength = 248 nm or one laser pulse in a range of 0.25-0.55 J/cm 2 for 200 ns at a wavelength = 308 nm, respectively [START_REF] Huet | Laser thermal anneal formation of atomically-flat low-resistive germanide contacts[END_REF][START_REF] Lim | Formation of epitaxial metastable NiGe 2 thin film on Ge (1 0 0) by pulsed excimer laser anneal[END_REF]. The growth of the NiGe 2 phase is not thermodynamically favorable in bulk reactions, as can be seen in Fig. 3.2 and Table . 3.1. However, Lim et al. associated the growth of this phase to the NiGe 2 film ability to be constructed epitaxially on Ge (1 0 0) with a low interfacial energy [START_REF] Lim | Formation of epitaxial metastable NiGe 2 thin film on Ge (1 0 0) by pulsed excimer laser anneal[END_REF]. A further analysis considering in-plane RSM or in-plane pole figures can be highly adequate in order to properly identify the nature of the peak observed at 26.0°.

The ex-situ θ-2θ scans displayed in Fig. 4.56 can be compared to the ones obtained from samples annealed by rapid thermal annealing (RTA) (see Fig. 4.58). This comparison allowed to verify that (i) the phase sequence was comparable, (ii) no peak related to Sn segregation was observed from out-of-plane XRD analysis for laser annealed samples and (iii) laser annealing might modify the mono-stanogermanide texture. Indeed, con-cerning the last aspect, only one out-of-plane peak was observed related to the Ni(GeSn) phase on laser annealed samples, while a polycrystalline film, with various out-of-plane orientations, was obtained from RTA samples (see Fig. 4.58). A further analysis of this behavior will be given in section 4.3.3. The phase sequence for both RTA-samples and laser annealed samples goes as follows The Ni-rich phase (i.e. Ni 5 (GeSn) 3 ) is the first phase to appear. This latter is consumed to the benefit of the Ni(GeSn) phase. However, for RTA samples, Sn segregation can occur at low temperatures (i.e. 350 ), which was verified by the occurrence of β-Sn diffraction lines at 30.6°and 32.0°. Meanwhile, similar peaks were not observed for laser annealed samples.

Two aspects have to be kept in mind to justify or understand this phenomenon. First, laser annealing is actually a commonly use technique to ensure a proper distribution of ion implanted dopants on Si. Ion implantation is used thanks to the high control over both the depth of the implantation and the precise amount of dopant introduced into the lattice, that can be achieved. However, this process also creates damage in the layer, which can promote a random accumulation of the implanted dopant throughout the lattice. A high temperature thermal cycle is then necessary to remove the damage and allow dopant atoms enough energy to move to a site where they can be electrically active. It is during this thermal annealing cycle that many of the challenges are revealed. Excessive diffusion or precipitation of the dopant (mostly when the solubility of the implanted atom is low in the semiconductor material) and deactivation of dopants can be encountered. However, the non-uniform thermal gradients from laser annealing can control dopants diffusion and activation. The metastable condition of the laser process can give rise to very characteristic impurity segregation phenomena [START_REF] Poate | Laser Annealing of Semiconductors[END_REF]. Substitutional Sn in Ge is introduced under metastable conditions, far from equilibrium, due to the low solid solubility of Sn on Ge. For these reasons, Sn segregation can be promoted at high temperatures. However, the specific conditions of laser annealing might control Sn distribution, as occurs for dopants on Si, which prevents segregation to occur and promotes a homogeneous distribution in the intermetallic.

Second, processing times in heat pulse annealing are usually small fractions of a second (160 ± 2 ns four this study), thereby making the solid-state diffusion negligible. However, the atoms in the melted zone can have high mobilities. This can allow mass transport over tens of nanometers even for heat pulses in the nanosecond range. When this occurs, the different atoms in the layer are necessarily redistributed [START_REF] Poate | Laser Annealing of Semiconductors[END_REF]. Thanks to these rapid growth velocities achieved in pulsed laser annealing, recrystallization of the melted region takes place under conditions that are far from equilibrium. Substitutional impurities can be incorporated into the lattice at concentrations that exceed equilibrium solubility limits during crystallization of a phase. The high growth velocity achieved by laser annealing provides then the opportunity to achieve high-speed, non equilibrium crystal growth phenomena under well-controlled experimental conditions. The Ni-rich and Ni(GeSn) phases growth was verified by ex-situ XRD in Fig. 4.56 and 4.57. Under these specific conditions, kinetics might not be favorable then for Sn segregation. This atom might instead re-arrange in the Ni(GeSn) lattice during the phase growth. Indeed, Sn incorporation was verify for these kind of intermetallics, see chapter III.

The addition of a RTA step at 200 after laser annealing was also studied and is presented in Fig. 4.59. The objective was to verify if the Ni(GeSn) growth could be enhanced at lower temperatures when combining laser and rapid thermal annealings as observed for Ti-silicides. Indeed, the combination of a first UV-nanosecond laser annealing (UV-NLA) step followed by RTA was efficient to reduce the transformation temperatures for Ti / Si system [START_REF] Esposito | Impact of nanosecond laser energy density on the reduction of C54-TiSi 2 temperature[END_REF]. The phase sequence remained comparable as the one described in Fig. 4.56. The Nirich phase was formed first, followed by the growth of the NiGe or Ni(GeSn) phases. The threshold temperature for the Ni-rich phase growth varied however. The Ni 5 Ge 3 phase was verified to appear after only 0.36 J/cm 2 pulse and Ni 5 (GeSn) appeared at 0.32 J/cm 2 pulse. Once the RTA step was added, the Ni-rich peak for both kinds of samples was observed since the lowest annealing energy (i.e. 0.08 J/cm 2 ). Additionally, no further peak related to Sn segregation was observed. The combination of a first step of UV-NLA An increase in roughness due to complex wrinkling was observed in TiN structures [START_REF] Gence | Wrinkled titanium nitride nanocomposite for robust bendable electrodes[END_REF]. The reason of this surface structure modification was related to thermally induced compressive stress of the TiN layer which originated from the difference between the layer coefficient of thermal expansion with the substrate [START_REF] Gence | Wrinkled titanium nitride nanocomposite for robust bendable electrodes[END_REF]. Wrinkling of thin films is a well known and generally an unwanted phenomenon and occurs when the compressive stress overpass a critical value (depending on the film / substrate system). The schematic behavior of TiN / Ni / GeSn sample is presented accordingly in Fig. 4.62. The Ni / GeSn layers underneath the TiN capping layer is expected to act as a "flexible" substrate when melted. Then, compressive stress due to the difference in thermal expansion can lead to the TiN wrinkling to liberated the internal stress. The occurrence of this particular stress conditions in the film might be one of the reasons for the texture variations on the Ni(GeSn) phase observed in Fig. 4.59. Indeed, texture development in silicides or germanides films are sensitive to variations in lattice constants (and dspacings) of the substrate or the film. The varying stress conditions might change the lattice parameter during crystal growth and impact the mono-stanogermanide phase texture [START_REF] Schutter | Texture in thin film silicides and germanides: A review[END_REF]. The use of high temperature millisecond laser annealing on NiPt / Si samples was proven to properly grow a NiPt-monosilicide in which smaller grains and increase epitaxial orientations, were obtained. One of the possible causes for this change was indeed related to the variation of residual stress in the layers [START_REF] Feautrier | Impact of laser anneal on NiPt silicide texture and chemical composition[END_REF]. Meanwhile the growth of NiPt-germanosilicides, by laser annealing, also generated a highly textured (NiPt)(SiGe) layer [START_REF] Setiawan | Laser-induced Ni(Pt) germanosilicide formation through a self-limiting melting phenomenon on Si 1-x Ge x /Si heterostructure[END_REF].

On the other hand, wrinkling phenomenon varied with the number of pulses for the same energy density, see Fig. Wrinkles amplitude increased when the laser pulse further increased. In general, the evolution of wrinkle patterns involves three stages an initial growth, coarsening and equilibration. They are characterized by a gradual increase and saturation of the wrinkle wavelength and amplitude, resulting in the formation of equilibrium wrinkle patterns [START_REF] Shugurov | Wrinkling of the metalpolymer bilayer: the effect of periodical distribution of stresses and strains[END_REF]. Then, both the increase of energy density or use of multipulses can promote wrinkles coarsening. In order to evaluate the advantage of using laser annealing, compared to the traditional RTA, the surface morphological evolution of Ni / GeSn from both kinds of thermal budgets is presented in Fig. 4.64.

For both kinds of samples and up to low temperatures, (< 300 for RTA samples) or low energy densities (< 0.32 J/cm 2 for laser annealed samples), the cross-hatch pattern was observed. The initial conditions were thus comparable for both kinds of samples treated with different annealings. At 350 , the surface morphology drastically changed with a strong increase of the surface roughness, for RTA samples. This change was linked to the formation of the mono-stanogermanide phase and its subsequent agglomeration, aggravated by Sn segregation. At higher temperatures, the roughness values gradually increased due to the amplification of both phenomena, see Fig. 4.65. For laser annealed samples, the mechanism of degradation changed however. The increase of RMS roughness values noticed in Fig. 4.65 were related to the onset of wrinkles formation and the following increase in the amplitude of these defects when the energy density increased. Furthermore, coupling both annealing treatment, i.e. annealing the samples with 1 pulse laser followed by RTA at 200 for 30 s in N 2 environment, did not further modify the morphological evolution. In Fig. 4.66 it can be verified that the trend is the same as for the laser annealed samples. The cross-hatch pattern was preserved up to 0.32 J/cm 2 , UV-nanosecond laser annealing followed by wrinkles formation in the surface at 0.34 J/cm 2 , which amplitude increased with the energy density. However, the RTA step at 200 did not seem to modify wrinkles amplitude or increase the overall roughness of the layers. Both, the surface patterns and the RMS roughness values remained comparable.

Ni / Ge and Ni / GeSn electrical evolution

The sheet resistance (R sh ) was measured with a four-point probe meter in order to follow the electrical evolution with laser annealing treatment. Fig. 4.67 presents the R sh evolution with the energy densities (ED) for Ni / Ge and Ni / GeSn samples.

The overall trends for the R sh evolution obtained for both systems were comparable. At low energy densities, the R sh values remained constant at around 28-31 Ω/sq (the Ni / GeSn system was slightly less resistive than the Ni / Ge one). After a certain energy density, 0.32 J/cm 2 for the Ni / GeSn system and 0.36 J/cm 2 for the Ni / Ge, the R sh values slightly increased. This increase was related to the growth of the Ni 5 Ge 3 or Ni 5 (GeSn) 3 phases. Indeed, these phases were verified to grow at these energies by XRD patterns in Fig. 4.56. Additionally, the onset for wrinkles formation observed in Fig. 4.60 was highly linked to the growth of the intermetallic phases. Later on, the R sh values progressively decreased when reaching the threshold energies for both NiGe and Ni(GeSn) phases growth at 0.44 J/cm 2 and 0.38 J/cm 2 , respectively as verified by XRD analyses in Fig. 4.56. Minimum R sh values of 22.1 Ω/sq at 0.54 J/cm 2 and 19.7 Ω/sq at 0.50 J/cm 2 were encountered for the Ni / Ge and Ni / GeSn systems, respectively. The sheet resistance decrease at lower energy densities for the Ni / GeSn system was linked to the formation of the Ni(GeSn) phase at lower energies compared to the Ni / Ge system. Modifying the amount of shots for a constant energy value did not induce a significant variation of the electrical evolution. Fig. 4.68 presents the R sh evolution for the Ni / GeSn systems for laser annealed samples with either 1, 3 up to 10 striking pulses. No matter the amount of pulses, the electrical evolution was comparable for all the samples. R sh values remained constant around 28-30 Ω/sq up to energy at which the Ni 5 (GeSn) 3 phase was obtained. R sh slightly increased due to the formation of this latter, R sh of 23.1 Ω/sq was achieved at 350 , while the minimum of laser annealed samples was of 19.1 Ω/sq. After reaching these minimum values, the R sh values increased. This increase was related to the degradation of the morphological properties. However, when comparing RTA and laser annealed samples, it can be noticed that the reached R sh values (once morphological degradation occurred) diferred importantly. R sh were importantly higher for RTA samples as observed in Fig. 4.69 (b).

The agglomeration and Sn segregation was verified to occur at 350 for RTA samples (see chapter III, while a complete lost of continuity of the layers was observed at 400 (see Fig. 3.41). This corresponded to the main cause of the lost of electrical stability. Meanwhile, the degradation of the top surface due to wrinkles growth was observed for laser annealed samples in section 4.3.3. Even with the apparition of these wrinkles, the Ni(GeSn) layer underneath the TiN capping is expected to remain continuous. The electrical stability can thus be maintained. Additionally, an increase of the Ni(GeSn) texture was observed in Fig. 4.58 and was related to stress variation in the layer in section 4.3.3. Indeed, the increase of epitaxial texture can lead to better defined interfaces, which are expected to decrease the intrinsic contact resistivity and thus would be consistent with the improved electrical performance compared to RTA samples [START_REF] Feautrier | Impact of laser anneal on NiPt silicide texture and chemical composition[END_REF].

A reduction on the crystallite size was also related to an improvement of the morphological stability [START_REF] Srolovitz | Capillary instabilities in thin films[END_REF][START_REF] Miller | The instability of polycrystalline thin films: Experiment and theory[END_REF], which was verified to occur when PAI or Pt alloy were used for Ni / GeSn samples. Meanwhile the use of laser annealing indeed promoted a reduction of grain size for silicides [START_REF] Feautrier | Impact of laser anneal on NiPt silicide texture and chemical composition[END_REF]. The full width at half maximum (FWHM) of the Ni monostanogermanide, was therefore measured, from various XRD patterns obtained both for samples annealed by RTA or by laser pulses. The results obtained are listed in Table 4.7 and displayed in Fig. 4.70. According to the Scherrer equation, the crystallite size is inversely proportional to the FWHM [START_REF]Elements of X-ray diffraction[END_REF]. Most FWHM values for laser annealed samples (after 0.48 J/cm 2 ) were higher than the ones obtained for RTA samples. This means that the grain size of the mono-stanogermanide were then smaller for laser annealed samples, which can be beneficial to maintain a continous Ni(GeSn) layer and prevent important electrical degradation.

On the other hand, the impact of combining laser annealing to a RTA step at 200 , on the electrical properties evolution, was also verified. The results, concerning both Ni / Ge and Ni / GeSn systems, are presented in Fig. 4.71.

The trend remained comparable, between each system and compared to the samples that were only laser annealed. From 0.08 J/cm 2 up to 0.42 J/cm 2 for the Ni / Ge system and 0.36 J/cm 2 for the Ni / GeSn system, the R sh was about 34-38 Ω/sq. These values were slightly higher than the ones obtained for laser annealed samples only (around 28-30 Ω/sq). This behavior was related to the growth of the Ni 5 Ge 3 or Ni 5 (GeSn) 3 phases. Indeed, XRD analyses in Fig. 4.59, verified that coupling the RTA step at 200 allowed the growth of the Ni-rich phase at lower energy densities. Later on, R sh values decreased until reaching minimum values, thanks to the presence of the NiGe or Ni(GeSn) phase, verified by XRD analysis in Fig. 4.59. After this, the same trend was fulfilled by the Ni / Ge system comparing laser annealed samples with and without the RTA step. Meanwhile, a difference scenario took place for Ni / GeSn samples. Indeed, an important R sh decreased was evidenced at 0.42 J/cm 2 and minimum values of 11.4 Ω/sq were obtained at 0.44 J/cm 2 . This suggested an improvement on the electrical properties which was not possible to achieve with RTA or laser annealing alone. Neither, a variation on the phase sequence (see Fig. 4.59) or a variation on morphological properties (see Fig. 4.66) was detected. Then, this improvement could be related to a modification on morphological interfaces states than might promote better electrical conduction within the layer. Finally, the increase on R sh values due to morphological surface degradation (wrinkles formation) enabled R sh to increase. However, the RTA step at 200 did not promoted further electrical degradation and the evolution remained comparable.

To summarize the main results of this section, the impact of thermally annealing Ni / GeSn layers by means of laser pulses was analyzed. As far as phase sequence was concerned, the overall evolution was comparable to traditional RTA samples the Ni 5 (GeSn) 3 phase was obtained first, followed by the mono-stanogermanide (Ni(GeSn)) phase. A change on texture was however evidenced (only one out-ofplane orientation was detected for laser annealed samples) while the crystallite size of the Ni(GeSn) phase seemed to be reduced with this thermal budget. This latter aspect could contribute to thermal stability improvements.

The morphological evolution of the samples differed importantly. Indeed, at low energy densities, the common cross-hatch pattern was observed. Then, initial conditions compared to RTA samples were verified. Yet, agglomeration or Sn segregation (island formation in the surface) was not noticed for laser annealed samples contrary to RTA samples. Instead, wrinkles were formed in the top surface due to thermal induced compressive stress relaxation and a difference on the thermal expansion between the TiN layer and the melted layer underneath. These morphological changes induced an important increase on root-mean-square roughness values and a slightly increase on R sh values. These values remained however quite low compared to the ones observed for RTA-samples once the agglomeration onset was passed.

Laser annealing was verified to be a suited annealing technique to form Ni contacts on GeSn layers, with low thermal budgets, in which the electrical performance was optimal. In addition, coupling laser annealing with a RTA step, at low temperatures, allowed to achieve very low R sh that were not possible to obtain with either laser annealing or RTA only.

Conclusion of Chapter IV

To summarize this chapter, three main alternatives to enhance the Ni / GeSn system thermal stability and/or delay Sn segregation were analyzed. These techniques corresponded to (i) pre-amorphization by ion implantation in which the us of Si, Ge or C atoms were tested. (ii) The addition of an alloy element to the Ni metallization layer, for this case the addition of 10.% of Co or Pt was tested. (iii) The final option, included the use of UV-nanosecond laser annealing as the thermal process to obtain the Ni(GeSn phase), as schematized in Fig. . The impact of these alternatives was measured in terms of phase sequence, morphological and electrical properties evolution during the solid-state reaction.

Concerning the phase sequence, the phase evolution was comparable to the one obtained for the Ni / GeSn system and described in chapter III. The hexagonal Ni 5 (GeSn) 3 was obtained first. This phase was later consumed to the benefit of the Ni(GeSn) phase. However, the different PAI used (Si, Ge or C) and the alloy elements used (Co or Pt) modified the threshold temperatures for the phase formation. In general, Ge PAI enhanced the phase growth at lower temperatures, Si and C PAI slightly and strongly delay the Ni(GeSn) formation to higher temperatures, respectively. This phenomenon was related to the ion segregation / accumulation at the intermetallic / substrate interfaces or at grain boundaries due to size and concentration variations. On the other hand, the use of a NiCo alloy enhanced the phase formation at lower temperatures while the use of a NiPt alloy did the opposite. This phenomenon was related to the alloy solubility in the intermetallic phases and segregation or accumulation phenomena and the formation of transient compounds. Concerning laser annealing, the phase sequence followed the trend explained before, however it followed a simultaneously sequence (that means that one phase was formed while the previous was still present).

On the other hand, Co and Pt alloys modified the phase sequence due to the formation of transient phases related to the presence of Co or Pt. For the NiCo samples the presence of Co x Ge y or Co x Sn y compounds was identified by in-situ XRD and in-plane RSM, however the exact nature of these compounds was not identify with certainty. On the other hand, the formation of PtSn x compounds was verified by means of XRD and TEM analysis. Laser annealed samples with multiple pulse shots present the incidence of a peak that could be related to Ni x Sn y or the metastable NiGe 2 phases. These kinds of Conclusion of chapter IV phases have not been observed for the multiple analysis realized during this PhD thesis or comparable analysis from literature. Then the proper identification of these compounds can be highly instructive to the understanding of the Ni / GeSn system when annealed by laser annealing and will need further work.

The Sn segregation threshold was positively impacted by the use of all the alternatives listed earlier. For PAI samples, the use of C PAI importantly delayed Sn segregation (up to 550 ), Co addition delayed this phenomenon up to 400 and Pt addition up to 450 . Meanwhile, samples annealed by laser pulses di not show the incidence of β-Sn diffraction peaks for all the range of energy densities analyzed (0.8-0.66 J/cm 2 ), not even when multiple shots were realized or when the thermal budget consisted in the use of a thermal pulse coupled to RTA at 200 . Then, C PAI, Co and Pt addition where well suited to delay Sn segregation while laser annealing suppressed this phenomena.

Concerning the morphological and electrical evolution, both properties were impacted by the different alternatives tested. C PAI enhanced both morphological en electrical thermal stability by delaying the agglomeration and Sn segregation up to higher temperatures. Co addition did not impact the morphological or the electrical properties compared to Ni / GeSn systems. Meanwhile, Pt addition showed a similar effect than C PAI treatment enhancing the morphological and electrical properties stability by delaying the agglomeration and Sn segregation up to higher temperatures. For laser annealed samples, the electrical properties evolution was comparable to the Ni / GeSn samples treated by classical RTA, however the degradation of these properties with high energy density was less severe than when using RTA. However, the morphological evolution was completely different. Wrinkles appeared in the surface once the melting onset of the layer underneath the TiN layer was reached. This effect allowed the stress relaxation from the TiN capping layer which enabled it to form this morphology.

The results of this chapter were highly instructive in order to understand and propose alternatives to improve the Ni / GeSn contacts performance and ensure high quality, stability and reliability. Based on the combination of the experimental results menationed earlier, the best compromise for Ni-based contact technology in terms of contact resistivity and thermal stability can be proposed. First the GeSn semiconductor can follow a PAI process combining Ge and C atoms, later Ni metallization with 10 at.% of Pt can be realized and finally thermal annealing combining laser pulse (around 0.4 J/cm 2 ) and RTA at low temperatures (around 150-250 ) can be performed. Previously, a comprehensive study concerning the Ni / GeSn intermetallics properties and a range of possibilities to improve the Ni / GeSn thermal stability were presented in chapter III and chapter IV, respectively. A combination of structural, morphological and electrical characterization techniques (described in chapter II) were used in order to assess the phase sequence, the evolution of the morphological and the electrical properties during the solid-state reaction. These analyses enabled to identify the nature of the phases appearing when Ni-based thin films were deposited on GeSn layers, describe the morphological evolution and acquire the sheet resistance evolution during heat treatments, which allowed a proper identification of thermal ranges of stability, see chapter III. Additionally, the impact of pre-amorphization by implantation, alloying the Ni layer with Pt or Co and the use of laser pulses as annealing treatment was discussed (see chapter IV). However, previous chapters were strictly focused on Ni-based metallizations.

One alternative to these kinds of metallization corresponds to the use of titanium (Ti) as metallization layer. Ti corresponds to a CMOS-compatible metallization that can be considered to contact GeSn based devices. Indeed, ultra-low values of specific contact resistivity were achieved with the use of Ti when contacting GeSn layers with 5 at.% of Sn [START_REF] Wu | Ultra-low specific contact resistivity (1.4x10 -9 ω.cm 2 ) for metal contacts on in-situ Ga doped Ge 0.95 Sn 0.05 film[END_REF][START_REF] Wu | Formation of Ohmic contacts on laser irradiated n-type 6H-SiC without thermal annealing[END_REF][START_REF] Wu | Record low specific contact resistivity (1.2x10 -9 ω-cm 2 ) for p-type semiconductors: Incorporation of Sn into Ge and in-situ Ga doping[END_REF]. In order to understand the Ti / GeSn solid-state reaction, a full assesment of the phase sequence and the properties evolution with annealing temperatures is necessary.

For these reasons, an overview of the use of Ti will be given in this chapter as schematized in Fig. The Ti / Ge phase sequence during the solid-state reaction (SSR) has already been studied [START_REF] Thomas | Reaction of titanium with germanium and silicon-germanium alloys[END_REF][START_REF] Thomas | Nucleation and growth in the reaction of Titanium with Germanium and some Silicon-Germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF][START_REF] Han | Studies of Ti-and Ni-germanide schottky contacts on n-Ge (1 0 0) substrates[END_REF][START_REF] Janardhanam | Annealing temperature dependency of the electrical and microstructural properties of Ti and Pt contacts to n-type Ge substrates[END_REF][START_REF] Pelleg | A note on the reactions in the Ti-Ge system[END_REF], however certain debate concerning the phase sequence was evidenced. In general, the TiGe 2 phase can be obtained after annealing at high temperatures, around 650-800 [START_REF] Thomas | Reaction of titanium with germanium and silicon-germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF]. Meanwhile, the first phase appearing was indexed as the high resistivity Ti 6 Ge 5 phase and the Ti 5 Ge 3 phase was not observed in certain cases [START_REF] Thomas | Reaction of titanium with germanium and silicon-germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF]. The coexistence of the Ti 5 Ge 3 and Ti 6 Ge 5 phases was however reported [START_REF] Han | Studies of Ti-and Ni-germanide schottky contacts on n-Ge (1 0 0) substrates[END_REF][START_REF] Janardhanam | Annealing temperature dependency of the electrical and microstructural properties of Ti and Pt contacts to n-type Ge substrates[END_REF]. Later, Chapter V. Alternative to Ni-based contacs Ti metallization Pelleg et al. showed that the first phase to appear was the Ti 6 Ge 5 phase, the presence of the Ti 5 Ge 3 phase was obtained upon annealing at higher temperatures and finally the TiGe 2 phase was later obtained [START_REF] Pelleg | A note on the reactions in the Ti-Ge system[END_REF].

On the other hand, the Ti / GeSn phase sequence during the SSR has not been thoroughly investigated. Galluccio et al. indicated that no reaction occurred between Ti / Ge 0.92 Sn 0.08 layers after 30 s annealing at 300 and 500 [START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF]. Meanwhile, Wu et al. confirmed the existence of a Ti(GeSn) x layer after annealing at 450-500 for 30 s [START_REF] Wu | Formation of Ohmic contacts on laser irradiated n-type 6H-SiC without thermal annealing[END_REF]. No clear consent concerning the nature and identification of the phases appearing during the Ti / GeSn SSR was evidenced. For these reasons and in agreement with the general steps realized to characterize Ni-based contacts, we have initiated the analysis of the Ti / Ge and Ti / GeSn phase sequences. Additionally, the evolution of the layers morphological and electrical properties with temperature was evaluated. The potential benefits of using Ti to contact GeSn-based devices against Ni-metallization have been discussed accordingly.

For this aim, high quality Ge and GeSn layers with 10 at.% of Sn were epitaxially grown on Ge-buffered Si (1 0 0) substrates in a reduced pressure chemical vapor deposition (RP-CVD) tool. A surface preparation was performed using an Ar plasma treatment for 60 s. Then, a deposition at room temperature of 20 nm of Ti with 7 nm of TiN capping layer was performed by magnetron sputtering. The phase formation sequence and crystalline evolution were monitored by in-situ and ex-situ X-ray diffraction, as described in chapter II, section 2.1.1. In addition, samples were annealed using a rapid thermal annealing (RTA) treatment in a N 2 environment for 30 s at temperatures ranging from 150 up to 650 . Then, complementary analyses of the surface morphology and the electrical properties evolution were also acquired by atomic force microscopy and four-point probe analyses, respectively.

5.1 Ti / Ge system phase sequence Fig. 5.2 shows the phase sequence during the solid-state reaction followed by in-situ XRD for the Ti / Ge system as a contour map, where the y-axis shows the temperature and the x-axis the diffraction angles. At 390 , the Ti layer was consumed and a peak appeared around 34.3°. This peak could be related to the growth of either Ti 5 Ge 3 or Ti 6 Ge 5 phases. This peak remained stable in a short range of temperature, i.e. 390-470 . At 470 , a peak appeared around 39.3°. This latter could only be indexed as the Ti 6 Ge 5 phase. Then the previous peak at 34.3°was considered to be related to the Ti 5 Ge 3 phase. Around 850 , the Ti 6 Ge 5 peak intensity importantly decreased (highlighted by a red rectangle in Fig. 5.

2).

In order to complement the in-situ XRD results, Ti / Ge samples, obtained under the same process conditions, were ex-situ annealed for 30 s at temperatures ranging from 150 up to 650 in a rapid thermal annealing (RTA) furnace under N 2 environment. Then, θ-2θ scans were acquired respecting the conditions described in chapter II, section 2.1.1. The results obtained are presented in Fig. 5.4. From the as-deposited state and up to 450 peaks of TiN (1 1 1), Ti (0 0 2) and Ti (1 0 1) were observed at 36.5°, 38.3°and 40°, respectively. At 400 , Ti peaks slightly shifted to lower 2θ angles, similar behavior was observed previously from the in-situ XRD At 450 three peaks were observed at 34.4°, 37.0°and 40.0°. These latter could be related either to the Ti 5 Ge 3 or the Ti 6 Ge 5 phases. The phase sequence is expected to remain comparable to the one observed by in-situ XRD results in Fig. 5.2. Then, these three peaks observed at 450 are expected to be realated to the Ti 5 Ge 3 phase. This phase remained stable up to 650 . Peaks related to Ti 6 Ge 5 or TiGe 2 were not detected with further annealing. The thermal budgets difference in ex-situ compared to in-situ measurements might modify the threshold temperatures for phase formation due to variations in diffusion kinetics. This phenomena will be discussed later in section 5.3.

Ti / GeSn system phase sequence

The Ti / GeSn system, (GeSn layer with 10 at.% of Sn content), was also analyzed thanks to both in-situ and ex-situ XRD diffraction techniques, respecting the conditions described in chapter II, section 2.1.1. The results are presented in Fig. 5.6 and 5.9.

Similar as for the Ti / Ge system in Fig. 5.2, at low temperatures, characteristic peaks of TiN (1 1 1) at 36.4°, Ti (0 0 2) and (1 0 1) at 38.4°and 40.1°, respectively, Si (4 0 0) and Ge/GeSn (4 0 0) around 68°, were observed. At 260 , the Ti peaks intensity decreased significantly. Additionally a slight shift of these peaks towards lower 2θ angles was observed together with an important decrease of intensity, as for the case of Ti / Ge samples, see Fig. 5.7. This can be related to Ti consumption and phase transition. The These latter remained stable up to 530 . After this temperature, the peaks disappeared to the benefit of the growth of a peak located at 39.3°. Like for the Ti / Ge system case, this latter could only be indexed as the Ti 6 (GeSn) 5 phase. Then, the previous three peaks obtained at lower temperatures were considered to be related to the Ti 5 (GeSn) 3 phase.

Around 800 the Ti 6 (GeSn) 5 peak shifted towards lower 2θ angles while its intensity decreased. Then, the growth of a peak at 37.8°, related to the Ti(GeSn) 2 phase, was noticed. A difference was observed concerning the threshold temperatures for phase formation between Ti / Ge and Ti / GeSn systems as displayed in Table 5.1 and schematized in Fig. 5.8. Sn addition might then play a role on diffusion kinetics, surface or interface energies which promoted phase nucleation at different temperatures. For GeSn-based samples, the Ti peak shift and total consumption, together with the T 5 (GeSn) 3 phase formation were promoted at lower temperatures. Meanwhile, the T 6 (GeSn) 5 phase formation was pushed up to 530 . In addition, the Ti(GeSn) 2 phase 204 Ti / Ge and Ti / GeSn systems evolution was observed. These differences will be further discussed in section 5.3. In order to complement the in-situ XRD results, Ti / GeSn samples, obtained under the same process conditions, were ex-situ annealed for 30 s at temperatures ranging from 250 up to 650 in a rapid thermal annealing (RTA) furnace under N 2 environment. Then, θ-2θ scans were acquired. The results obtained are presented in Fig. 5.9. From the as-deposited state and up to 350 peaks related to TiN (1 1 1), Ti (0 0 2) and Ti (1 0 1) were observed at 36.5°, 38.3°and 40.0°, respectively. At 450 three peaks that could be related to either the Ti 5 (GeSn) 3 or the Ti 6 (GeSn) 5 phases were observed at 34.4°, 37°and 40°, similar as for the in-situ results. Additionally, two peaks located at 30.5°and 32.0°were also evidenced at this temperature. These latter could only be related to β-Sn diffraction peaks due to Sn segregation. At 550 , weak peaks appeared at 41.3°and 45.0°. These were then related to the growth of the Ti 6 (GeSn) 5 phase. Then the previous peaks observed at 34.4°, 37°and 40°were related to the Ti 5 (GeSn) 3 phase. Peaks related to the Ti(GeSn) 2 phase were not observed even for the highest thermal annealing (i.e. 650 ). The Ti / Ge phase sequence during the solid-state reaction has been largely analyzed [START_REF] Thomas | Reaction of titanium with germanium and silicon-germanium alloys[END_REF][START_REF] Thomas | Nucleation and growth in the reaction of Titanium with Germanium and some Silicon-Germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF][START_REF] Han | Studies of Ti-and Ni-germanide schottky contacts on n-Ge (1 0 0) substrates[END_REF][START_REF] Janardhanam | Annealing temperature dependency of the electrical and microstructural properties of Ti and Pt contacts to n-type Ge substrates[END_REF][START_REF] Pelleg | A note on the reactions in the Ti-Ge system[END_REF]. However, certain disagreement can be encountered in literature concerning the thermal stability of the Ti-rich phases. In general, the first phase appearing was indexed as the high resistivity Ti 6 Ge 5 phase, which can be consumed at high temperatures (650-800 )

to the benefit of the TiGe 2 phase [START_REF] Thomas | Reaction of titanium with germanium and silicon-germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF]. For those studies, the growth of the Ti 5 Ge 3 phase was not observed. Other studies have verified the coexistence of the Ti 5 Ge 3 and Ti 6 Ge 5 phases [START_REF] Han | Studies of Ti-and Ni-germanide schottky contacts on n-Ge (1 0 0) substrates[END_REF][START_REF] Janardhanam | Annealing temperature dependency of the electrical and microstructural properties of Ti and Pt contacts to n-type Ge substrates[END_REF]. Yet, the Ti 6 Ge 5 phase was considered to form first at temperatures around 300 and the Ti 5 Ge 3 phase was obtained at higher temperatures (around 400-500 ). Meanwhile, the TiGe 2 was observed after annealing at 700 [START_REF] Han | Studies of Ti-and Ni-germanide schottky contacts on n-Ge (1 0 0) substrates[END_REF][START_REF] Janardhanam | Annealing temperature dependency of the electrical and microstructural properties of Ti and Pt contacts to n-type Ge substrates[END_REF].

Later, Pelleg et al. also indexed the first phase appearing as the Ti 6 Ge 5 phase and the presence of the Ti 5 Ge 3 phase was observed upon annealings at higher temperatures. Finally, the TiGe 2 was the last phase obtained [START_REF] Pelleg | A note on the reactions in the Ti-Ge system[END_REF].

It can be considered then that Ti / Ge reactions result in the growth of the Ti 6 Ge 5 phase followed by the TiGe 2 phase and Ti 5 Ge 3 is a transient phase that could be obtained under certain specific annealing conditions. From our results, we believe that the Ti 5 Ge 3 phase formed first, followed by the Ti 6 Ge 5 phase and finally the TiGe 2 phase was the last phase appearing at higher temperatures. Gaudet et al. also agreed that different phase sequences have been reported for the Ti / Ge system. The general consent is however that the TiGe 2 phase is the last phase appearing at relatively high temperatures (800 ) and is the less resistive phase [START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]. From in-situ XRD measurements of the evolution of 30 nm thick Ti layer on Ge (0 0 1), Gaudet et al. observed the Ti layer consumption to occur around 420 . This was followed by the growth of an unidentified phase which was consumed at around 550 to the benefit of the Ti 6 Ge 5 which was stable up to 900 [START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]. This sequence is in agreement with the results we obtained. From Fig. 5.2 it can be observed that the Ti peaks disappeared around 390-400 to the benefit of a phase considered as the Ti 5 Ge 3 (which could be correlated to the unidentified phase observed by Gaudet et al. [START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]). The latter was consumed next to the benefit of the Ti 6 Ge 5 phase while the TiGe 2 phase was not observed even after annealing up to 900 .

On the other hand, Wu et al. tested the deposition of Ti on Ge 0.95 Sn 0.05 layers and followed an annealing for 30 s in N 2 environment at 250, 450 and 500 . Cross-section TEM images coupled to EDS analysis confirmed the formation of a Ti x (GeSn) y intermetallic at 450 , the Ti total consumption at 500 and Sn segregation at 500 . The intermetallic phase composition was closed to the one of Ti 6 (GeSn) 5 . However, further analyses are necesary to unambiguously determine the nature of the phase [START_REF] Wu | Formation of Ohmic contacts on laser irradiated n-type 6H-SiC without thermal annealing[END_REF].

One of the rules for predicting the transformation sequence in thin film metal / semiconductor couples (mostly metal-silicides), states that the first phase to nucleate is the most stable congruently melting compound adjacent to the lowest-temperature eutectic on the bulk-equilibrium phase diagram [START_REF] Ohring | Materials Science of Thin Films: Deposition and Structure[END_REF]. From the Ti-Ge binary phase diagram, displayed in Fig. 5.10, it can be observed that three phases are thermodynamically stable Ti 5 Ge 3 (hexagonal), Ti 6 Ge 5 (orthorhombic) and TiGe 2 (orthorhombic). From these phases, the Ti 5 Ge 3 follows the rule stated before because Ti 5 Ge 3 composition of the liquid formed is the same as the composition of the solid formed during melting and is the only congruent melting in the binary phase diagram. On the other hand, an effective heat of formation (EHF) model formulated for different metal / semiconductor systems states that the first compound to form during metal / silicon or metal / germanium interaction is the congruently melting phase with the most negative effective heat of formation [START_REF] Ohring | Materials Science of Thin Films: Deposition and Structure[END_REF]. The enthalpies of formation for the different Ti-Ge compounds are listed in Table 5.2. Ti 5 Ge 3 is again the phase that suits the previous statement. Concerning the formation entalphies, Ti 5 Ge 3 and Ti 6 Ge 5 presented the lowest values, however they were quite similar. Yet, Ti 5 Ge 3 phase is the only one that presents a congruent melting in the Ti-Ge binary phase diagram. For these reasons it was considered that for both Ti / Ge and Ti / GeSn systems, the phase sequence evolved as follows the Ti-rich Ti 5 Ge(Sn) 3 phase 207 Chapter V. Alternative to Ni-based contacs Ti metallization was formed first. This phase was consumed to the benefit of the Ti-rich Ti 6 Ge(Sn) 5 phase. Later on, at higher temperatures, the Ge-rich TiGe 2 phase was obtained.

On the other hand, a great difference concerning the threshold temperatures for phase formation was observed when comparing in-situ and ex-situ XRD results for both Ti / Ge and Ti / GeSn systems. As expressed in chapter III, section 3.3, the diffusion step is important to achieve mass transport which is a time-dependent process. Additionally, temperature has a great influence on the coefficients and diffusion rates [START_REF] Callister | Materials Science and Engineering: An Introduction[END_REF]. Between insitu and ex-situ analyses, the difference in thermal budgets is obvious. For in-situ XRD analysis, one sample was heated from 50 up to 900 following temperature steps of 5 .

Once the targeted temperature was reached, the sample remained at this temperature for 7 min while the XRD pattern was acquired. Meanwhile, for RTA treated samples, each annealing temperature corresponds to a sample that was heated up to the targeted temperature for 30 s and immediately cooled down to acquired the ex-situ XRD pattern. The 7 min thermalisation promoted during the in-situ measurement might allow massdiffusion to occur which was beneficial for the phase growth.

For Ti-based germanides and stanogermanides, the diffusion step might then importantly govern the phase sequence. For the Ti / Ge system, experimental backscattering spectrum analysis of Ti / Ge / Si mobility showed that, as soon as the reaction started and the TiGe 2 phase was observed, both Si and Ge atoms were homogeneously distributed in the layer providing evidence for the high mobility of the metalloids (Ge and Si) in the Ti / Ge intermetallic compounds. [START_REF] Thomas | Some titanium germanium and silicon compounds: Reaction and properties[END_REF]. On the other hand, Aldrich et al. evidenced Ge out-diffusion and segregation during Ti / Si 1-x Ge x , which verified the high mobility of this atomic specie during the reaction [START_REF] Aldrich | Film thickness effects in the TiSi 1-x Ge x solid phase reaction[END_REF]. Then for the case of RTA samples, even if the temperature is high enough for nucleation, the mass transport (of Ge or GeSn for this case) might not be enough for the phase growth. This aspect generated then an important gap between the temperature threshold for phase formation between in-situ and ex-situ XRD analyses. Furthermore, a difference on the threshold temperatures for phase formation was evidenced when comparing Ti / Ge and Ti / GeSn systems. Indeed, for GeSn-based samples, Ti consumption, the Ti 5 Ge 3 and TiGe 2 phases growth was promoted at lower temperatures (see Table 5.1). Indeed, the presence of impurities or additional elements can play a role on the phase observed during the solid-state reaction due to the modification of interaction kinetics [START_REF] Queirolo | Diffusion in Silicides: Basic Approach and Practical Applications[END_REF]. Sn addition to Ge was expected then to modify the surface / interface energies, the diffusion kinetics which can promote the phases growth at lower temperatures.

Ti / Ge and Ti / GeSn systems morphological evolution

AFM images were acquired for samples annealed ex-situ by RTA. The objective was to compare both the morphological evolution between Ge-and GeSn-based systems, metallized with Ti, and between Ni-and Ti-based systems. Fig. 5.11 shows AFM images of the surface of samples annealed at various temperatures for Ti / Ge and Ti / GeSn samples.

The surfaces of the as-deposited samples were uniform and smooth with low root- A distinctively change on morphology was observed at 450 . This change ocurred concomitantly with the Ti 5 Ge 3 or Ti 5 (GeSn) 3 phase formation, evidenced by ex-situ XRD in Fig. 5.4 and 5.9. On the other hand, the start of agglomeration processes could also have taken place, this phenomenon needs complementary characterization techniques (as TEM images) to be verified. Indeed, RMS roughness values increased progressively with annealing temperatures, as can be observed in Fig. 5.13. RMS roughness values were sligthly higher for Ti / GeSn system. At 450 , Sn segregation occurred, see Fig. 5.9.

Sn segregation, could then explain the highest roughness values for these samples. Tibased metallization was investigated as a possible alternative to Ni-based ones. Then, a comparison on the morphological evolution between Ni / GeSn and Ti / GeSn system is necessary. The comparative AFM images are presented in Fig. 5.12.

From the images acquired in the as-deposited state for Ti / Ge, Ti / GeSn and Ni / GeSn samples, the cross-hatch pattern was observed. Then, the initial state was comparable for all the systems. As mentioned previously, Ti / Ge and Ti / GeSn surface morphology remained stable up to 350 . For Ni / GeSn layers, important agglomeration already occurred at this temperature together with Sn segregation (see Fig. 3.10). Meanwhile, Ti-based samples remained stable and at this temperature the samples still showed the cross-hatch pattern. This can be related to the fact that the intermetallic phases were not obtained at these temperatures. Later on, with the temperature increase, RMS roughness values progressively increased for all the samples as seen in Fig. 5.13. The RMS roughness values increase was more important for Ti / GeSn samples compared to Ti / Ge samples and even more important for Ni / GeSn samples. The difference Ti / Ge and Ti / GeSn systems evolution between Ti / Ge and Ti / GeSn systems, can be related to Sn segregation phenomenon, which was verified to occur at 450 for the Ti / GeSn system. At temperatures lower than 550 , RMS roughness values of the Ti / GeSn system remained lower than the ones obtained for the Ni / GeSn samples. Ti metallization could be then useful to enhanced the morphological thermal stability for GeSn-based contacts. For both kinds of samples and in the as-deposited state, R sh values were around 40 Ω/sq. For Ti / Ge samples, values remained relatively constant during the entire range of annealing temperatures analyzed (from 150 up to 650 ). A slight increase was however noticed at 450 concomitantly with the morphology changes observed by AFM in Fig. 5.11 and the Ti 5 Ge 3 phase formation evidenced by ex-situ XRD in Fig. 5.4. No further decrease of the R sh values was evidenced when the annealing temperatures increased. Indeed, the temperature dependence of the sheet resistance, of titanium germanides formed by RTA, was evidenced. The formation of Ti-rich phases normally generates an important increase of R sh while the growth of the TiGe 2 phase enables an important reduction of the resistance values, after annealing around 800-850 [START_REF] Thomas | Nucleation and growth in the reaction of Titanium with Germanium and some Silicon-Germanium alloys[END_REF][START_REF] Ashburn | Formation of Titanium and Cobalt Germanides on Si (1 0 0) using rapid thermal processing[END_REF][START_REF] Thomas | Some titanium germanium and silicon compounds: Reaction and properties[END_REF][START_REF] Hong | Formation of a C49 TiGe 2 phase during annealing a coevaporated Ti 0.33 Ge 0.67 alloy[END_REF].

For our case, no big variation was observed, this could be related to the fact that the annealing temperatures were not sufficiently high to form the low resistivity TiGe 2 phase, as verified in section 5.1 and 5.2.

For Ti / GeSn samples, the R sh observed evolution was different. Values remained stable up to 250 . An increase on R sh was noticed after annealing at 350 , which can be related to the presence of the Ti 5 (GeSn) 3 phase (see Fig. 5.9), consistently with the increase of surface roughness (see Fig. 5.13). Then a decrease, after annealing at 450 , on R sh down to 22-25 Ω/sq, was noticed. This could be related to the growth of the Ti 6 (GeSn) 5 phase. After this, R sh values remained relatively constant and the values were slightly lower compared to the Ti / Ge system.

As stated earlier, Ti metallization could present an alternative to Ni-based metalizations. For these reasons, it is highly important to evaluate the electrical evolution of one system compared to the other. Ti / GeSn and Ni / GeSn R sh evolution with annealing temperature is then displayed in Fig. 5.15. when contacting GeSn substrates [START_REF] Galluccio | Formation and characterization of Ni, Pt, and Ti stanogermanide contacts on Ge 0.92 Sn 0.08[END_REF]. The R sh slightly increase after 200 was related to the Ni 5 (GeSn) 3 phase formation. After 300 , the R sh decreased. The minimum R sh values were obtained at 350 . This temperature corresponds to the growth of the

Conclusions of Chapter V

This section covered a comprehensive overview concerning the use of Ti metallization to contact GeSn-based samples. In this aim, the phase sequence, the morphological and electrical properties evolution during the solid-state reaction were evaluated using the same methodology as for previous chapters.

The phase sequence during the solid-state reaction went as follows the Ti 5 (GeSn)) 3 phase was formed first. This phase was consumed to the benefit of the Ti 6 (GeSn)) 5 phase. Later on, at higher temperatures, the Ti(GeSn) 2 phase was obtained. It was encountered that the phase sequence depended importantly on the thermal budgets used. Great differences were evidenced concerning threshold temperatures for phase formation between in-situ and ex-situ analysis. The diffusion step might play an important role for the phase growth. On the other hand, great differences were also evidenced between Ge-and GeSnbased samples. Sn addition might modify surface / interface energies and promote phase formation at lower temperatures.

For Ti-samples, the Sn segregation and morphological changes were delayed to higher temperatures compared to Ni-samples. Additionally, R sh values were lower and more stable for Ti / GeSn system compared to Ni / GeSn system. The choice of this metal compared to Ni-based metallization can vary concerning integration issues in the devices and the needs to be fulfilled. However from the results obtained, Ti CMOS compatible metallization is considered as a suitable alternative to contact GeSn-based devices.

General Conclusions and Perspectives

Conclusions

The studies carried out within the framework of this PhD thesis aimed at developing and fully characterizing Ni / GeSn intermetallics. These intermetallics can be integrated in electronic or opto-electronic devices for the realization of suitable contacts for GeSn-based devices. For these aims, the evolution of Ni-based / GeSn systems during the solid-state reaction was analyzed. Indeed, modifications were made on different process levels, either concerning the GeSn layer modification, the Ni-based metallization layer or the annealing treatment. The common points for these multiple kinds of samples were the complementary characterization techniques that were used systematically and under comparable conditions to establish the appropriate parameters for contacts elaboration. The general assessment of the contacts properties was based on the study of the phase sequence, the morphological and electrical evolution with annealing temperatures. For this, multiple XRD techniques were used coupled to AFM, SEM or TEM structural analyses and the four-point probe method for electrical evaluation.

Firstly, multiple parameters of the Ni / Ge 1-x Sn x (x ranging from 0 up to 15 at.% Sn content) evolution during the solid-state reaction were inspected. The phase sequence is schematized in Fig. 6.1 (a) and went as follows a Ni-rich phase was formed first. Thanks to advanced and combined X-ray diffraction analyses (out-of-plane, in-plane and pole figures XRD), this phase was undoubtedly identified as the hexagonal ε-Ni 5 (GeSn) 3 metastable phase which is schematized in Fig. 6.1 (b). When this latter was consumed, the Ni(GeSn) orthorhombic phase was obtained. The thermal stability of the Ni-rich phase and consequently the threshold temperatures for the Ni(GeSn) formation varied with the Sn content of the GeSn layer. A trend was observed the higher the Sn content is, the higher the temperature to form the Ni(GeSn) phase. This phenomenon was related to the preferential segregation and accumulation of Sn around grain boundaries (GB), as displayed in Fig. 6.1 (c). These atoms acted then as stuffed barriers and hampered Ni diffusion which delayed the phase formation. For very high Sn content stacks (i.e. Sn > 15 at.%) the trend reversed due to full Sn segregation and migration towards the surface at low temperatures. On the other hand, the crystalline orientation of the Ni(GeSn) was altered with Sn content as well, as schematized in Fig. 6.1 (d). Indeed, a change of the preferential texture of the intermetallic was verified thanks to the combination of out-of-plane, in-plane and pole figures XRD techniques. In general, Sn addition can promote a better established epitaxial alignment between the Ni(GeSn) film and the GeSn substrate underneath. Results showed that the Ni(GeSn) phase is suitable to contact GeSn-based devices. This phase was obtained at relatively low temperatures (between 250-350 ) and the R sh values were found to be low (between 10-20 Ω/sq). However, the thermal stability of these kinds of layers was quite limited. Degradation of the intermetallics can compromise both the morphological and electrical properties. This trend can be verified in Fig. 6.2. The degradation process was observed to occur due to two phenomena. First, the grain boundary grooving followed by agglomeration and, second, Sn segregation which aggravated the stability issues. This degradation impacted directly the electrical performance leading to an important increase of R sh at around 350 and reducing the range of stability in which the contacts were reliable.

Alternative methods to enhance Ni / GeSn thermal stability were afterward studied, as schematized in Fig. 6.3. Indeed, three different solutions, that can be implemented in different steps of the process elaboration, were tested. Pre-amorphization by implantation (PAI) technique consisted of the use of different ions (i.e. Si, Ge or C) implanted to create an amorphous GeSn layer prior to Ni metallization, as displayed in Fig. 6.3 (a). The impact of the use of this technique varied from one ion to the other. Indeed, the ion size, concentration and the preferential segregation / accumulation behavior of each ion promoted variations of the phase sequence. The piled up Ge at the interfaces might have modified the interface energies, favoring the phase growth at lower temperatures. C, on the contrary, importantly delayed the Ni(GeSn) growth. This phenomenon was related to C accumulation at grain boundaries which hampered Ni diffusion. C accumulation around GB also delayed the threshold temperature for Sn segregation and for the agglomeration process. The latter, were beneficial for the Ni(GeSn) thermal stability. Indeed, the surface morphology and the R sh values remained remarkably stable up to 500 . From these results, the combination of Ge and C PAI were considered as the best combination to allow the Ni(GeSn) phase growth at low temperatures and reach high thermal stability. The addition of an alloy element to the Ni metallization layer was also evaluated. Both the addition of 10 at.% of Co or Pt were tested, as displayed in Fig. 6.3 (b). Concerning NiCo metallization, the phase sequence remained comparable to Ni metallization. The addition of Co allowed however the growth of the (NiCo) 5 (GeSn) 3 and (NiCo)(GeSn) phases at slightly lower temperatures. On the other hand, the growth of Co x Sn y or Co x Ge y compounds at high temperatures can occur. Concerning the layer stability, Sn segregation threshold was delayed up to 400 . However, morphological or electrical properties stability were not enhanced. Agglomeration occurred in the same manner as for Ni systems and electrical degredation was also verified. Meanwhile, Pt addition importantly modified the phase sequence. First, the NiPt consumption was pushed to higher temperatures. This was related to the snowplow effect and the low solubility of Pt in the intermetallic phases (verified by XRD and TEM characterization techniques).

Later, the formation of PtSn x compounds was confirmed (after around 300 ). The formation of these compounds was enabled thanks to the preferential segregation of both Pt and Sn at GB. On the other hand, Pt addition also pushed the threshold temperature for Sn segregation up to 450 . The thermal stability of the layers benefited from these phenomena. The morphological stability was enlarged thanks to the delay on grain boundary grooving and agglomeration process while the range in which R sh remained low and stable was also enhanced. For these reasons the Ni 0.9 Pt 0.1 alloy was considered as a well suited metallization to contact GeSn devices in which low values of R sh can be achieved and the morphological and electrical integrity can be maintained at higher temperatures.

A novel annealing technique was also tested, as displayed in Fig. 6.3 (c). The successful use of laser annealing to achieve the Ni(GeSn) growth was verified thanks to XRD analysis. The morphological evolution of the samples differed from samples obtained by traditional rapid thermal annealing (RTA) process. Wrinkles formed on the top surface, due to thermal induced compressive stress and a difference on the thermal expansion between the TiN layer and the melted layer underneath, were observed. On the other hand, Sn segregation was avoided thanks to this annealing technique. The increase of morphological stability and the suppression of Sn segregation positively impacted the electrical evolution. Indeed, even if the tendency was comparable to RTA samples, values remained quite low even when the surface morphology was disrupted. From these results, the combination of laser pulse at low energies (i.e. around 0.4 J/cm 2 ) and RTA at low temperatures (between 150-250 ) was considered to be a well suited thermal annealing process to obtain the Ni(GeSn) phase with enhanced electrical properties.

Based on our experimental results, the best compromise for Ni-based contact technology in terms of contact resistivity and thermal stability can be the combination of different processes. First the GeSn semiconductor can follow a PAI process combining Ge and C atoms, later Ni metallization with 10 at.% of Pt can be realized and finally thermal annealing combining laser pulse (around 0.4 J/cm 2 ) and RTA at low temperatures (around 150-250 ) can be performed, as schematized in Fig. 6.4. This combination of process can ensure high quality, stable and reliable contacts.

Furthermore, a brief overview of an alternative to Ni-based metallization consisting in the use of Ti was also presented. As for Ni-based metallization, the analyses were performed in terms of phase sequence, morphological and electrical evolution during the solid-state reaction. The phase sequence went as follows. The Ti 5 (GeSn) 3 phase was formed first. This phase was consumed to the benefit of the Ti 6 (GeSn) 5 phase.

Later on, at higher temperatures, the Ti(GeSn) 2 phase was obtained. The threshold temperatures for phase growth depended importantly on the thermal budgets used and the Sn addition to the Ge layer. For Ti-based samples, Sn segregation and morphological changes were delayed to higher temperatures compared to Ni-based samples. Additionally, R sh were lower and more stable for Ti / GeSn system compared to Ni / GeSn system. For these reasons, Ti CMOS compatible metallization was also considered as a possible alternative to contact GeSn-based devices.

Perspectives

The achievements of this PhD thesis represent a good starting point to elaborate high quality GeSn-based contacts that can be fully integrated in either electronic or optoelectronic devices. So far, a comprehensive understanding of the Ni / GeSn system and properties has been accomplished. In addition, alternatives to modify the process conditions were identified and were verified to be suited to enhance the Ni / GeSn thermal stability.

Furthermore, the results obtained open certain perspectives to continue to enrich the General Conclusions and Perspectives current knowledge on the Ni / GeSn system like -Improving the surface pre-clean parameters for GeSn layer prior to the metallization step. Surface preparation treatments are mandatory to remove particles, organic materials, metallic impurities and native oxides from the surface. Identifying the proper conditions for GeSn layers can be then useful for process elaboration.

-Refine the range of PAI doses and energies that can be applied for GeSn layer without damaging the surface.

-Identifying the exact nature of the PtSn x , Co x Ge y and Co x Sn y compounds formed during the NiPt / GeSn and NiCo / GeSn solid-state reactions, respectively.

-Modifying the strain condition of the TiN capping layer to avoid the wrinkles formation during laser annealing. This can avoid the morphological degradation and the sligth increase on R sh values.

On the other hand, the next critical steps imperatively need a deep exploration on the applicability and performance of these contacts within the actual devices. The study of the contact resistance (R c ) or resistivity (ρ c ) thanks to TLM structures can be useful to properly complement the analyses of the electrical properties of the contacts. The use in devices as LED, phtodetectors or lasers is considered as the main perspective for the following work.

In order to pursue this perspective, photodiodes and TLM structures were elaborated following the process flow shown in Fig. 6.5, for thick GeSn layers (16 at.% Sn) capped with SiGeSn layers. The contact was made on doped (n-type and p-type) Ge. The TLM structures were realized with 6 metallic pads, of 100 µm x 100 µm, spaced from 40 up to 140 µm. The metal deposition was realized with either 10 nm of Ni 0.9 Pt 0.1 alloy or with a Ti/Au couple. The samples were annealed up to 350 in a rapid thermal annealing furnace for 30 s in N 2 environement.

The values acquired either for Ni 0.9 Pt 0.1 or TiAu / samples are displayed in Table 6.1. Both kinds of samples showed an ohmic behavior. However, NiPt samples were less resistives, an important reduction was acquired concerning both R c (1.6 Ω) and ρ c (4.9x10 -6 Ω.cm 2 . These improvement on the electrical conductivity can explain the increase of the cut-off frequency (boundary frequency response at which the energy flowing through the system begins to be attenuated) observed in Fig. 6.6 for the photodiodes, when the NiPt alloy was used. The frequency was measured at 1550 nm at room temperature for different polarizations. This increase observed for NiPt-based samples indicated and enhancement on the device performance.

From these preliminary results, it was observed that NiPt allowed to reach lower resistivity values and increase the cut-off frequency. This first analysis indicated that ohmic contacts with acceptable values of contact resistivity and enhanced performance can be achieved. These results are very promising in order to pursue further improvements on the contact technology that could be integrated on GeSn-based devices.

Introduction

"Innovation", "haute performance" et le "transfert rapide de données" font partie des phrases qui peuvent décrire les progrès de l'industrie de la micro-électronique au cours des dernières années. Afin d'ameliorer les performances, différentes adaptations peuvent être faites en modifiant les étapes de production, l'architecture du dispositif et les matériaux qui constitueront les dispositifs. Ces types de développements sont réalisés pour améliorer les performances des dispositifs électroniques et opto-électroniques. Le germanium-étain (GeSn), alliage du groupe IV, correspond à un matériau intéressant afin d'être intégré dans ces deux types de dispositifs. Le GeSn peut être utilisé comme source et drain, pour stresser les MOSFET en Ge, et pour obtenir des canaux à haute mobilité pMOSFET et pTFET. D'autre part, l'ajout d'une quantité suffisante de Sn dans le réseau du Ge (environ 10 at.%) offre la possibilite d'obtenir une structure de bande interdite directe. Ce matériau pourrait alors être utilisé pour obtenir un laser du groupe IV monolithiquement intégré et compatible avec la technologie CMOS (complementary metal oxide semi-conductor) sur Si. Quelle que soit l'application, des contacts métalliques sont nécessaires pour injecter du courant électrique dans les dispositifs. Des contacts à base de Ni ont été proposés et étudiés car ils permettent la formation de l'intermétallique Ni(GeSn) à basse température avec de faibles valeurs de résistance. Néanmoins, les contacts à base de Ni présentent une mauvaise stabilité thermique liée à l'agglomération de la couche de Ni(GeSn) et à la précipitation des atomes de Sn.

Cette thèse de doctorat était consacrée au développement et à la caractérisation de contacts de Ni / GeSn. L'objectif était de proposer des contacts efficaces et fiables pour des dispositifs à base de GeSn. Cette thèse a été divisée en cinq chapitres qui vont être présentées par la suite. Les principaux problèmes de ce matériau sont la faible solubilité de Sn dans le réseau Ge (< 1%) et les différences importantes de paramètres de maille (> 14%). Ces aspects peuvent rendre difficile la croissance de couches stables de GeSn, une séparation de phases et la ségrégation de Sn peuvent facilement se produire. Cependant, des couches de GeSn de haute qualité avec une teneur élevée en Sn (> 10%) ont été obtenues avec des conditions de croissance CVD à basse température [START_REF] Aubin | Low temperature epitaxy of Si, Ge, and Sn based alloys[END_REF].

Ensuite, un aspect important à garder à l'esprit, pour les dispositifs à base de GeSn, correspond au développement et à l'élaboration de contacts de haute qualité, stables et fiables. Dans le cas de couples métaux / semiconducteur (M / S) de faible épaisseur, la minimisation de l'énergie permettra la croissance de phases intermétalliques de type Brièvement, la nature de l'élément d'alliage a un impact sur la séquence de phase pendant la réaction à l'état solide. L'addition de Co a un impact positif sur la réduction de la température de formation pour les phases (NiCo) 5 (GeSn) 3 et (NiCo)(GeSn). Par contre, l'addition de Pt retarde la consommation du métal et la formation des phases vers des températures plus élevées.

F.4.3 Utilisation du recuit laser comme technique de recuit

Le but ultime du recuit d'un système semi-conducteur / métal est de favoriser son évolution vers un état thermodynamique plus stable. Ce recuit permet la diffusion des atomes et la formation des phases intermétalliques. Pour cette raison, il a été de vérifié si la technique de recuit laser est adaptée pour former correctement la phase Ni(GeSn). Des analyses DRX ex-situ ont été réalisées pour les échantillons de Ni / GeSn recuits avec une impulsion laser dans la gamme de 0.08 à 0.66 J/cm 2 . Les diagrammes de diffraction obtenus sont présentés sur la Fig. Une alternative à ce type de métallisation correspond à l'utilisation de titane (Ti) comme couche de métallisation. Ti correspond à une métallisation compatible CMOS qui peut être considérée pour les dispositifs basés sur GeSn. En effet, des valeurs ultra-faibles de résistivité de contact ont été obtenues avec l'utilisation de Ti sur des couches de GeSn avec 5 at. % Sn [START_REF] Wu | Ultra-low specific contact resistivity (1.4x10 -9 ω.cm 2 ) for metal contacts on in-situ Ga doped Ge 0.95 Sn 0.05 film[END_REF][START_REF] Wu | Formation of Ohmic contacts on laser irradiated n-type 6H-SiC without thermal annealing[END_REF][START_REF] Wu | Record low specific contact resistivity (1.2x10 -9 ω-cm 2 ) for p-type semiconductors: Incorporation of Sn into Ge and in-situ Ga doping[END_REF]. Afin de vérifier si ce métal est adapté pour des dispositifs à base de GeSn, une évaluation de la séquence de phases et de l'évolution des propriétés en fonction des températures de recuit est nécessaire.

Le système Ti / GeSn, (couche GeSn avec 10 at. % de Sn), a été analysé grâce à des techniques de diffraction DRX in-situ. Les résultats sont présentés sur la Fig. Aux basses températures, des pics caractéristiques de TiN (1 1 1) à 36.4°, Ti (0 0 2) et (1 0 1) à 38.4°et 40.1°, respectivement, Si(4 0 0) et Ge / GeSn (4 0 0) autour de 68°, sont observés. À 260 , l'intensité des pics de Ti diminue de manière significative. De plus, un léger décalage de ces pics vers des angles ayant une plus faible valeurs de 2θ est observé avec une diminution importante de l'intensité. Cela peut être lié à une consommation de Ti et à la transition de phase.

Trois pics à 34,3 °, 37,2°et 40,2°apparaissent avec la diminution d'intensité du Ti autour de 260 . La couche de Ti est consommée autour de 350 . Ces pics pourraient être liés à la croissance des phases Ti 5 (GeSn) 3 ou Ti 6 (GeSn) 5 . Ces derniers restent stables jusqu'à 530 . Au delà de cette température, les pics disparaissent au profit de la croissance d'un pic situé à 39,3°. Ce dernier ne peut être indexé que comme la phase Ti 6 (GeSn) 5 . Ainsi, les trois pics précédents, obtenus à des températures plus basses, ont été considérés comme liés à la phase Ti 5 (GeSn) 3 . Vers 800 , le pic Ti 6 (GeSn) 5 se déplace vers des angles 2 plus bas tandis que son intensité diminue. Finalement, la croissance d'un pic à 37.8°, lié à la phase Ti(GeSn) 2 , est observée. Abstract : Efforts on the semiconductor industry are constantly made to improve different parameters like the devices performance or the speed of data transference. To reach these advances, innovative alternatives can be considered : changing the production processes steps, the device architecture or the materials that will constitute the devices. Germanium-Tin (GeSn), group-IV alloy, corresponds to an interesting material to integrate in electronic or optoelectronic devices. GeSn material can be used as source and drain stressor in Ge MOSFETs (metal-oxide-semiconductor field-effect transistors), and as high mobility channels in both pMOSFETs and pT-FETs (thin field-effect transistors). On the other hand, the addition of a sufficient amount of Sn in the Ge lattice structure (about 10 at.%) offers the potential to get a direct band-gap structure. This material could be then used to achieve a monolithically integrated group-IV laser that is compatible with Si-CMOS technology (complementary metal-oxide-semiconductor). No matter the application, low-resistive, stable and reliable metallic or intermetallic contacts are key components to inject electric current in the devices. Ni / GeSn intermetallics have been considered as suited contact material for these kinds of applications. This PhD thesis was therefore dedicated to the systematic and comprehensive development and characterization of Ni / GeSn intermetallics to contact GeSn-based devices. The Ni / GeSn properties analyzed in terms of : phase sequence, morphological and electrical evolution during the solid-state reaction is presented first. As Ni / GeSn intermetallics exhibit a poor thermal stability, different alternatives such as the use of preamorphization by implantation (PAI), the addition of alloying elements (Co, Pt) and the use of laser annealing are also studied to enhance the contact thermal stability. Ultimately, another alternative concerning Ti metallization is also mentioned. Thanks to these studies, a comprehensive and systematic analysis of the Ni / GeSn system was realized. In addition, the identification of different alternatives to modify the process conditions that can enhance the Ni / GeSn system thermal stability was achieved. The results obtained represent a good starting point to elaborate high quality, stable and reliable GeSn-based contacts that can be fully integrated in electronic or opto-electronic devices. Resumen : Diferentes esfuerzos son realizados constantemente en la industria de semiconductores para mejorar diferentes par ámetros como el rendimiento de los dispositivos o la velocidad de transferencia de datos. Para alcanzar estos avances, se pueden considerar diferentes rutas innovadoras como, modificar las etapas de los procesos de producci ón, la arquitectura de los dispositivos o los materiales que constituir án a los mismos. Germanio-Esta ño (GeSn), una aleaci ón del grupo IV, corresponde a un material interesante para integrar en dispositivos electr ónicos u opto-electr ónicos. El GeSn se puede utilizar como estresor en la fuente y el drenador de MOSFET (metal-oxide-semiconductor field-effect transistors) de Ge, y para obtener canales de alta movilidad tanto en pMOSFET como en pTFET pTFETs (thin fieldeffect transistors). Por otro lado, la adici ón de una cantidad suficiente de Sn en la estructura del Ge (aproximadamente 10 at.%) ofrece el potencial de obtener un material a estructura de bandas directa. Este material podría utilizarse para lograr un l áser del grupo IV integrado monolíticamente compatible con la tecnología Si-CMOS (complementary metaloxide-semiconductor). Sin importar la aplicaci ón, los contactos met álicos o intermet álicos de baja resistencia, estables y confiables son componentes clave para inyectar la corriente el éctrica a los dispositi-vos. Ni / GeSn es considerado como el material de contacto adecuado para este tipo de aplicaciones. Por lo tanto, esta tesis se dedic ó al desarrollo y caracterizaci ón sistem ática e integral de componentes intermet álicos de Ni / GeSn para contactar dispositivos basados en GeSn. Primero se presentar án las propiedades de Ni / GeSn analizadas en t érminos de : la secuencia de fase, la evoluci ón morfol ógica y el éctrica durante la reacci ón en estado s ólido. Luego, se describir án diferentes alternativas para mejorar la estabilidad t érmica de los contactos basados en Ni. Desde el uso de la pre-amorfizaci ón por implantaci ón, aleaciones de NiPt o NiCo en lugar de Ni, hasta el uso de la impulsi ón l áser como t écnica de recocido. En última instancia, otro camino relacionado con la metalizaci ón de Ti tambi én ser á mencionada en conjunto con la verificaci ón de la resistividad de contacto. Gracias a este estudio, se realiz ó un an álisis comprensivo y sistem ático de la evoluci ón del sistema Ni / GeSn. Adicionalmente, diferentes alternativas para modificar las condiciones del proceso para aumentar la estabilidad t érmica del sistema Ni / GeSn fueron identificadas. Los resultados obtenidos representan un buen punto de de partida para elaborar contactos en GeSn que sean de alta calidad y estabilidad que puedan ser integrados en componentes electr ónicos u optoelectr ónicos.

Universit é Paris-Saclay

Universit é Paris-Saclay

Espace Technologique / Immeuble Discovery Route de l'Orme aux Merisiers RD 128 / 91190 Saint-Aubin, France

  ing method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.2 Ni / Ge and Ni / GeSn phase sequence . . . . . . . . . . . . . . . . 4.3.3 Ni / Ge and Ni / GeSn morphological evolution . . . . . . . . . . . 4.3.4 Ni / Ge and Ni / GeSn electrical evolution . . . . . . . . . . . . . . 4.4 Conclusion of Chapter IV . . . . . . . . . . . . . . . . . . . . . . . . . . . V Alternative to Ni-based contacts: Ti metallization 5.1 Ti / Ge system phase sequence . . . . . . . . . . . . . . . . . . . . . . . . 5.2 Ti / GeSn system phase sequence . . . . . . . . . . . . . . . . . . . . . . . 5.3 Discussion: Ti / Ge and Ti / GeSn systems comparison . . . . . . . . . . . 5.4 Ti / Ge and Ti / GeSn systems morphological evolution . . . . . . . . . . 5.5 Ti / Ge and Ti / GeSn systems electrical evolution . . . . . . . . . . . . . 5.6 Conclusions of Chapter V . . . . . . . . . . . . . . . . . . . . . . . . . . . General Conclusions and Perspectives A NiGe phase powder piffraction PDF chart from the ICDD data base B Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.85 Sn 0.15 systems pole figures C In-situ XRD patterns for different PAI samples

Figure 0 . 1 :

 01 Figure 0.1: Recent advances and increase in Ge 1-x Sn x publications per year from Scopus abstract and citation database from Elsevier.

Figure 0 . 2 :

 02 Figure 0.2: Schematic representation of the chapters division in this PhD thesis.

Figure 1 . 1 :

 11 Figure 1.1: Schematic representation of chapter I section division.

Figure 1 . 2 :

 12 Figure 1.2: Ge-Sn binary phase diagram. (a) Full diagram and (b) focus on the low-Sn concentrations zone [4]

Figure 1 . 3 :

 13 Figure 1.3: Timeline of the historical Ge 1-x Sn x benchmarks adapted from [6, 7].

Figure 1 . 4 :

 14 Figure 1.4: Surface AFM images (scan size 5 µm x 5 µm) of the surfaces of thin Ge 1-x Sn x layers grown

Figure 1 . 5 :

 15 Figure 1.5: Surface AFM images (scan size 5 µm x 5 µm) of the surfaces of thick Ge 1-x Sn x layers [5].

Figure 1 . 6 :

 16 Figure 1.6: Calculated band structures of Ge and α-Sn. Ge is an indirect bandgap semiconductors with positive band-gap 0.7 eV, while semiconducting α-Sn with a diamond lattice structure exhibits a negative and direct band-gap [6].

Figure 1 . 7 :

 17 Figure 1.7: GeSn indirect to direct band-gap transition, adapted from [6, 23].

Figure 1 . 8 :

 18 Figure 1.8: Schematic representation of the formation of a M / S intermetallic.

Contents

2. 1 2 . 2 2 . 3 2 . 4

 1222324 Structural characterization by X-ray diffraction . . . . . . . . . . . . . 2.1.1 Bragg-Brentano θ-2θ diffraction configuration . . . . . . . . . 2.1.2 In-plane reciprocal space map (RSM) . . . . . . . . . . . . . . 2.1.3 In-plane Pole Figures (IP-PF) . . . . . . . . . . . . . . . . . . Morphological characterization . . . . . . . . . . . . . . . . . . . . . . 2.2.1 Atomic Force Microscopy (AFM) Tapping mode . . . . . . . . 2.2.2 Transmission Electron Microscopy (TEM) . . . . . . . . . . . Electrical characterization . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1 Four-point probe electrical measurements . . . . . . . . . . . . Conclusions of Chapter II . . . . . . . . . . . . . . . . . . . . . . . . .

  2.1.

Figure 2 . 1 :

 21 Figure 2.1: Schematic representation of chapter II section division.

Figure 2 . 2 :

 22 Figure 2.2: Schematic representation of the diffraction of an incident X-Ray beam on atomic planes.

Figure 2 . 3 :

 23 Figure 2.3: Schematic representation of the Bragg-Brentano θ-2θ Diffraction Setup.

Figure 2 . 4 :

 24 Figure 2.4: Schematic representation of out-of-plane and in-plane XRD configurations.

Figure 2 . 5 :

 25 Figure 2.5: Example of ex-situ θ-2θ XRD diffraction patterns acquired.

Figure 2 . 6 :

 26 Figure 2.6: Example of the patterns acquired during in-situ θ-2θ XRD diffraction

Figure 2 . 7 :

 27 Figure 2.7: In-plane XRD set up for sample rotation [49].

Figure 2 . 8 :

 28 Figure 2.8: Example of an in-plane RSM patterns (a) flat view, (b) integration over a range of φ and (c) equivalent integrated in-plane θ-2θ acquired after integration in φ.

Figure 2 . 9 :

 29 Figure 2.9: Pole figure components [50].

Figure 2 . 10 :

 210 Figure 2.10: Movement of the goniometer during the acquisition of an in-plane pole figure [50].

Figure 2 . 11 :

 211 Figure 2.11: Example of the acquired pole figures and the analysis of the results achieved with the CaRine 3.1 software.

Figure 2 . 12 :

 212 Figure 2.12: Schematic representation of the AFM principle.

Figure 2 . 13 :

 213 Figure 2.13: Example of an acquired AFM image from a Ni / GeSn after annealing at 450 .

2 . 14 .

 214 Coupling both techniques can be highly advantageous to analyze the chemical composition and the morphological stability of the layers. Indeed, these techniques allowed Wirths et al. to (i) verify the continuity and crystallinity of Ni / GeSn intermetallic layers, (ii) follow the Ni and Sn atoms distributions and (iii) verify the phase formed at Chapter II. Characterization techniques for intermetallic contacts

Figure 2 . 14 :

 214 Figure 2.14: Example of TEM coupled with EDS image obtained for a Ni / Ge 0.85 Sn 0.15 sample annealed up to 250 , in which the Ge distribution was identified.

Figure 2 . 15 :

 215 Figure 2.15: Schematic representation of the four-point probe method.

  3.1.

Figure 3 . 1 :

 31 Figure 3.1: Schematic representation of chapter III section division.

Figure 3 . 2 :

 32 Figure 3.2: Ni-Ge binary phase diagram [63].

  3.3.

Figure 3 . 3 :

 33 Figure 3.3: Morphological assesment of Ni / GeSn layers (a) Poly-crystalline identification with HR-TEM images taken on Ni(Ge 0.94 Sn 0.06 ) and Ni(Ge 0.88 Sn 0.12 ) [18], (b) SEM images of different Ni / GeSn and Ni / GeSn samples with different Sn concentrations and annealing temperatures to verify thermal stability [21] and (c) AFM images of Ni / GeSn samples confirming surface degradation when changing the Sn concentration and annealing temperature [20].

Figure 3 . 4 :

 34 Figure 3.4: Schematic representation of R sh evolution with thermal annealing [9, 11, 12, 14, 15, 17, 18, 21, 58].

Figure 3 . 5 :

 35 Figure 3.5: Stability ranges for the minimum resistance values for different Ni / GeSn systems assessed by previous works [9, 11, 15, 21, 58].

  3.6).

Figure 3 . 6 :

 36 Figure 3.6: Ge and GeSn layers charaterization by ω-2θ scans (top row) and AFM images (scan size 5 µm x 5 µm, bottom row).

Figure 3 . 7 :

 37 Figure 3.7: Schematic representation of the steps to obtain the Ni / GeSn analyzed stacks.

3. 3 . 1

 31 The Ni / Ge 0.90 Sn 0.10 phase sequence As described in chapter II, section 2.1.1, the Ni / Ge 0.90 Sn 0.10 phase formation sequence and crystalline evolution were monitored by in-situ XRD. θ-2θ scans over the 20-90°range were performed. Analyses were taken from 50 up to 600 . The Ni / Ge 0.9 Sn 0.1 in-situ XRD pattern is shown in Fig.3.8. It displays a contour map, where the y-axis shows the temperature and the x-axis, the diffraction angle.

Figure 3 . 8 :

 38 Figure 3.8: In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni and a Ge 0.90 Sn 0.10 layer.

Figure 3 . 9 :

 39 Figure 3.9: In-situ X-ray diffraction measurement of the reaction between Ni and Ge 0.90 Sn 0.10 layers focus on the low temperature range 50 -180 .

Figure 3 . 10 :

 310 Figure 3.10: Ex-situ XRD diagrams measured during the solid-state reaction between Ni and Ge 0.90 Sn 0.10 at the as-deposited and annealed by RTA.

Figure 3 . 11 :

 311 Figure 3.11: Summary of the phase formation sequences obtained, thanks to in-situ and ex-situ XRD analyses for a 10 nm Ni deposited on Ge 0.90 Sn 0.10 layers.

Figure 3 . 12 :

 312 Figure 3.12: Surface AFM images (scan size 5 µm x 5 µm) of Ni / Ge 0.90 Sn 0.10 samples annealed at various temperatures by RTA.

Figure 3 . 13 :

 313 Figure 3.13: Schematic representation of the degradation process of a Ni / GeSn thin film by means of grain boundary grooving and agglomeration phenomena.

Figure 3 . 14 :

 314 Figure 3.14: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for the Ni / Ge 0.90 Sn 0.10 system.

Figure 3 . 15 :

 315 Figure 3.15: TEM images in bright field mode acquired for the Ni / Ge 0.90 Sn 0.10 system in the as-deposited state and after anneals at 350 and 400

Figure 3 . 16 :

 316 Figure 3.16: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the Ni / Ge 0.90 Sn 0.10 systems. From the as-deposited state and up to 300 , R sh values remained constant at around 25-28 Ω/sq. This temperature range corresponded to the Ni-rich phase domain, as evidenced in section 3.3.1. At 350 , minimum values of R sh were reached around 9.1 Ω/sq. At this temperature the growth of the Ni-monostanogermanide phase was verified, see Fig. 3.10.

Figure 3 . 17 :

 317 Figure 3.17: In-situ XRD patterns measured during the solid-state reaction of 10 nm of Ni and a Ge 0.9 Sn 0.1 layer with HF pre-clean.

Figure 3 . 18 :

 318 Figure 3.18: Summary of the in-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni and Ge 0.9 Sn 0.1 layers with different surface treatments. Dashed zone corresponds to the range of temperatures at which the phases were observed simultaneously.

Figure 3 . 19 :

 319 Figure 3.19: Energy area fit deconvolution for a GeSn substrate as-received for (a) Ge 2p and (b) Sn 3d.)

Figure 3 . 20 :

 320 Figure 3.20: In-situ XRD patterns measured during the solid-state reaction of 10 nm of Ni and Ge 0.90 Sn 0.10 without pre-clean and with Ar plasma treatment. Focus comparison of the crystalline out-of-plane orientation for the Ni-rich and Ni(GeSn) phase.

Figure 3 . 21 :

 321 Figure 3.21: Surface AFM images (scan size 5 µm x 5 µm) of Ni / Ge 0.90 Sn 0.10 samples with different surface preparation, annealed at various temperatures by RTA.

Figure 3 . 22 :

 322 Figure 3.22: Evolution of the root-mean-square as a function of the rapid thermal annealing temperature for the Ni / Ge 0.90 Sn 0.10 systems with different surface pre-clean preparations.

Figure 3 . 23 :

 323 Figure 3.23: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the Ni / Ge 0.90 Sn 0.10 systems with different surface pre-clean preparations.

3. 5 . 1

 51 Study of the Ni / Ge 0.90 Sn 0.10 system at low temperature In order to clarify the nature of the Ni-rich phase and to discuss the presence of the 2θ angle shift observed at 120-130 (see Fig.3.8 and 3.9) and the increase of the Ni-rich phase peak intensities, complementary in-plane RSM analyses were conducted. Samples were annealed respecting the annealing conditions and the thermal budgets (same ramp and soak steps) applied for the in-situ XRD measurement in Fig. 3.8 and explained in chapter II, section 2.1.1. In this case, the annealing was interrupted at 180 . Results obtained by in-plane RSM measurements for samples in the as-deposited state and annealed at 180 are respectively shown in Fig. 3.24 (a) and (b). The corresponding patterns integrated over the full φ range are presented in Fig. 3.25.

Figure 3 . 24 :

 324 Figure 3.24: In-plane RSM for Ni / Ge 0.90 Sn 0.10 samples in (a) the as-deposited state and (b) annealed up to 180 .

Figure 3 . 25 :

 325 Figure 3.25: In-plane RSM patterns integrated over the full φ range for Ni / Ge 0.90 Sn 0.10 samples in the as-deposited state and annealed at 180 .

Figure 3 . 26 :

 326 Figure 3.26: Evolution of the intensity as a function of temperature for the Ni peak and Ni-rich phase XRD peak located at about 32.0°in the Ni / Ge 0.90 Sn 0.10 system.

Figure 3 . 27 :

 327 Figure 3.27: Evolution of the position and FWHM as a function of temperature for the Ni-rich phase XRD peak located at about 32.0°in the Ni / Ge 0.90 Sn 0.10 system.

  Pole figures were acquired at a 2θ angle of 31.9°for the sample in the as-deposited state and the one annealed at 180 , the results are shown in Fig. 3.29. Additionally, a pole figure for 2θ = 45.3°at 180 was also recorded. The obtained pole figures are shown in Fig. 3.30.

Figure 3 . 28 :

 328 Figure 3.28: Stereographic projection indicating the angular positions of the main low-index poles of Ge / GeSn with (0 0 1) orientation of the cubic crystal.

Figure 3 . 30 :

 330 Figure 3.30: In-plane pole figures recorded at 2θ = 45.3°for Ni / Ge 0.9 Sn 0.1 samples annealed at 180 .

Figure 3 . 31 :

 331 Figure 3.31: Simulations of the stereoprojection of (3 1 2) on the (3 1 2) planes for a monoclinic ε'-Ni-rich phase.

Figure 3 . 32 :

 332 Figure 3.32: Summary of the orientations given in Table 3.4.

Figure 3 . 33 :

 333 Figure 3.33: Summary of the phase formation sequences obtained, thanks to in-situ XRD analyses for a 10 nm Ni deposited on Ge 1-x Sn x layers. Striped area corresponds to the range of temperature from which Ni and Ni-rich phase were observed simultaneously.

  AFM images of the Ni / Ge, Ni / Ge 0.94 Sn 0.06 , Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 systems are presented in Fig. 3.34 for comparison.

Figure 3 . 34 :

 334 Figure 3.34: Surface AFM images (scan size 5 µm x 5 µm) of Ni / Ge 1-x Sn x samples annealed at various temperatures by RTA.

  drastic for Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 systems. Finally, at 550 dots were observed in the surface for the Ni / Ge 0.94 Sn 0.06 , Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 systems and can be related to Sn segregation. Chapter III. Study of the Ni / GeSn system After 350 , a strong and progressive increase of the surface roughness occurred which was more pronounced when the Sn content increased, see Fig. 3.35. This stronger RMS roughness increase for high Sn content samples is a result of the drastic morphological changes observed for the Ni / Ge 0.90 Sn 0.10 and Ni / Ge 0.85 Sn 0.15 systems in Fig. 3.34.

Figure 3 . 35 :

 335 Figure 3.35: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for the Ni / Ge 1-x Sn x systems.

Figure 3 . 36 :

 336 Figure 3.36: TEM images in bright field mode acquired for Ni / Ge 0.90 Sn 0.10 (top row) at the as-deposited state and after anneals at 350 and 400 , and Ni / Ge 0.85 Sn 0.15 (bottom row) after anneals at 250 and 350

Figure 3 . 37 :

 337 Figure 3.37: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the Ni / Ge 1-x Sn x systems.

Fig. 3 .

 3 Fig. 3.38 presents a focus on the low temperature area of the Ni / Ge 1-x Sn x in-situ XRD patterns in the 20-40°angular region, which yields information on the evolution of the Ni-rich phase in samples with various Sn contents. From the as-deposited state, the Nirich peak was observed at 31.9°for the Ni / Ge sample and around 31.6°for Ni / GeSn samples (see Fig.3.39).

Figure 3 . 38 :

 338 Figure 3.38: Low temperature in-situ XRD patterns (in the 20-40°angular range) acquired during the solid-state reaction between Ni and Ge 1-x Sn x layers with Sn content ranging from 0 to 10 at.%. Focus on the Ni-rich phase peak.

Figure 3 . 39 :

 339 Figure 3.39: Ni-rich phase peak position evolution for different Ni / Ge 1-x Sn x systems.

Figure 3 . 40 :

 340 Figure 3.40: Low temperature in-situ XRD pattern (in the 20-55°angular range) acquired during the solid-state reaction between Ni and a Ge 0.85 Sn 0.15 layer.

  3.10. These samples were analyzed by cross-sectional TEM coupled with EDS and EELS maps. The objective was to map the distribution of the elements during the solid-state reaction, with a special focus on Sn incorporation, distribution and segregation. Fig.3.41 shows the cartography obtained for Ni / Ge 0.90 Sn 0.10 samples in the as-deposited state and after anneals at 200, 350, 400 and 550 .

Figure 3 . 41 :

 341 Figure 3.41: Cross-section TEM images coupled with EDS and EELS analysis of Ni / Ge 0.90 Sn 0.10 samples in the as-deposited state and after RTA anneals at 200, 350, 400 and 550 .

Figure 3 . 42 :

 342 Figure 3.42: EDS profiles of the Ni / Ge 0.90 Sn 0.10 sample annealed at (a) 200 and (b) 350 by RTA for 30 s.

Figure 3 . 43 :

 343 Figure 3.43: Schematics of Sn behavior during Ni-GeSn solid-state reaction.

Figure 3 . 44 :

 344 Figure 3.44: Schematics representation of atoms disposition for (a) Ni 5 (GeSn) 3 phase and (b) Ni(GeSn) phase.

  Fig 3.45.

Figure 3 . 45 :

 345 Figure 3.45: Schematics representation of (a) Ni diffusion through grain boundaries and (b) Sn accumulation at GB hampering Ni diffusion.

  3.8 and 3.38).

Figure 3 . 46 :

 346 Figure 3.46: In-situ XRD analysis of a heating-cooling cycle from 50 up to 600 and from 600 down to 55 , for (a) Ni / Ge 0.90 Sn 0.10 and (b) 175 for Ni / Ge 0.85 Sn 0.15 samples.

Figure 3 . 47 :

 347 Figure 3.47: Cross-section TEM images coupled with EDS analysis of Ni / Ge 0.85 Sn 0.15 samples after RTA at 250 and 350 .

Figure 3 . 48 :

 348 Figure 3.48: EDS profiles of the Ni / Ge 0.85 Sn 0.15 sample annealed at 250 by RTA for 30 s.

Figure 3 . 49 :

 349 Figure 3.49: Schematic representation of the pole figures that can be obtained from a film with different types of textures [101].

Figure 3 . 50 :

 350 Figure 3.50: Pole figures acquired for Ni on Ge (0 0 1) substrate for (a) (1 1 1) [52] and (b) (2 1 1)[START_REF] Gaudet | Reaction of thin Ni films with Ge: Phase formation and texture[END_REF] poles.

3

 3

  .50, corresponded to a combination of five different epitaxial components and three very weak axiotaxy components for which either the NiGe (1 2 1), (2 2 0), or (2 1 1) plane was aligned with a Ge {1 1 0} plane. Finally, studies of Stiphout et al. in thin films Ni-germanides, showed mostly epitaxial components together with weaker axiotaxial lines of the NiGe (2 1 1) and NiGe (2 0 2) planes aligned with the Ge (2 2 0) planes

Figure 3 .

 3 Figure 3.51: θ-2θ of different Ni / Ge 1-x Sn x samples annealed in-situ up to 600 and acquired after cooled down to room temperature.

Figure 3 . 52 :

 352 Figure 3.52: Integrated in-plane RSM patterns for Ni / Ge 1-x Sn x samples annealed up to 600 and acquired after cooled down to room temperature.

  1 1) with Ge/GeSn (2 2 0) was observed.Chapter III. Study of the Ni / GeSn system In order to fully understand Ni-germanides texture and the impact of Sn addition, in-plane pole figures were acquired for Ni / Ge and Ni / Ge 0.90 Sn 0.10 systems for different Bragg peaks at 2θ = 30.6°, 34.9°, 45.5°and 54.1°. Samples were in-situ annealed up to 600 and the analysis were realized when cooling down to room temperature. The results are presented in Fig.3.54.

Figure 3 . 53 :

 353 Figure 3.53: Full in-plane RSM patterns for Ni / Ge and Ni / Ge 0.90 Sn 0.10 samples annealed up to 600 . Patterns were acquired once the samples were cooling down to room temperature.

Figure 3 . 54 :

 354 Figure 3.54: In-plane pole figures for Ni / Ge and Ni / Ge 0.90 Sn 0.10 samples annealed up to 600 . The pole figures were measured after cooling the samples down to room temperature.

Table 3 . 6 :

 36 Summary of the orientations observed for the NiGe phase. Each symbol corresponds to an orientation in Fig.

Figure 3 . 55 :

 355 Figure 3.55: Summary of the main results described in chapter III.

Figure 4 . 1 :

 41 Figure 4.1: Schematic representation of chapter IV section division.

Figure 4 . 2 :

 42 Figure 4.2: In-situ XRD patterns measured during the solid-state reaction of 10 nm of Ni and Ge 0.90 Sn 0.10 with C PAI.

Figure 4 . 3 :

 43 Figure 4.3: Summary of the phase formation sequences obtained during the solid-state reaction between 10 nm of Ni and Ge 0.90 Sn 0.10 layers without PAI and with various PAI (Ge, Si, C and Ge + C). Stripped zones represent the range of temperatures in which the phases were present simultaneously.

Table 4 . 1 :

 41 Critic threshold temperatures during the Ni / GeSn phase sequence for the reference samples and PAI samples. Red color highlighted an increase on temperature compared to the reference while blue color highlighted a decrease of temperature system, a Ni-rich phase appeared since the as-deposited state. Ni consumption was observed at 120 . For all the PAI samples, the Ni-rich phase was observed since 95 . Si or Ge PAI did not modify the temperature for Ni consumption (120 ). However the intensity of this peak was much weaker than for the reference sample (see Fig. 4.4). Meanwhile, for C PAI and C + Ge PAI systems Ni consumption occurred around 95-105 . This can indicate that Ni diffusion happened easily in these samples. Pre-amorphization by ions implantation (PAI) Tsui et al. observed fast Ni dissolution and diffusion along the ion implantation induced defective zone during silicide formation with PAI process [107]. Additionally, Ozcan et al., Chen et al. and Turcotte-Tremblay et al. evidenced the growth of the phases at lower temperatures in amorphous or nearly amorphous substrates

Figure 4 . 4 :

 44 Figure 4.4: Comparison of the intensity evolution with temperature of the Ni XRD peak located around 44.9°for the systems without PAI and with Si, Ge, C and Ge + C PAI.

4. 1 . 1 . 3

 113 Impact of C PAI Ozcan et al. showed, by in-situ XRD, that C PAI at room temperature delayed NiSi formation by approximately 15 [130]. Luo et al. evidenced adittionally a delay on the NiSi 2 phase formation

Figure 4 . 5 :

 45 Figure 4.5: In-situ XRD patterns measured during the solid-state reaction of 10 nm of Ni and Ge 0.90 Sn 0.10 without PAI and with C PAI. Focus comparison of the crystalline out-of-plane orientation for the Ni-rich and Ni(GeSn) phase.

Fig. 4 .

 4 Fig. 4.6).

Figure 4 . 6 :

 46 Figure 4.6: Schematics representation of Si, Ge and C ions implanted on GeSn layer during the phase sequence.

Figure 4 . 7 :

 47 Figure 4.7: Cross-section TEM images coupled to EDS analysis of Ni / GeSn samples with Ge PAI (top row) and Ge + C PAI (bottom row), in the as-deposited state.

Figure 4 . 8 :

 48 Figure 4.8: Sn segregation threshold extracted from XRD analyses of samples annealed by RTA.

Figure 4 . 9 :

 49 Figure 4.9: AFM images (scan size 5 µm x 5 µm) of the surface of samples annealed at various temperatures for the Ni / GeSn systems witout PAI and with the different PAI treatments realized.

Figure 4 . 10 :

 410 Figure 4.10: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for the Ni / Ge 0.90 Sn 0.10 reference system and systems with different PAI treatments.

  ples treated with Ge and Ge + C PAI and annealed at 350 for 30 s in N 2 environment. The results obtained are shown in Fig. 4.11.

Figure 4 . 11 :

 411 Figure 4.11: Cross-section TEM images coupled with EDS analysis of Ni / GeSn samples with Ge PAI (top row) and Ge + C PAI (bottom row), annealed at 350 .

Figure 4 . 12 :

 412 Figure 4.12: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for Ni / GeSn without and with various PAI treatments.

1 . 1 .

 11 The results are shown in Fig.4.13 as a contour map. Characteristic peaks of TiN (1 1 1) at 36.4°, Ni or NiCo (1 1 1) at 44.9°, Si (4 0 0) and Ge/GeSn (4 0 0) around 68°were observed at low temperatures for both systems.

Figure 4 . 13 :

 413 Figure 4.13: In-situ XRD patterns measured during the solid-state reaction between (a) 10 nm of Ni on a Ge layer and (b) 10 nm of NiCo on a Ge layer.

Figure 4 . 14 :

 414 Figure 4.14: In-situ XRD patterns measured during the solid-state reaction between (a) Ni / Ge and (b) NiCo / Ge systems. Focus on the mono-germanide destabilization at high temperature.

  Detavernier et al. and Ozcan et al. observed a similar phenomenon when studying the NiSi formation by in-situ XRD diffraction[START_REF] Ozcan | Effects of temperature dependent pre-amorphization implantation on NiPt silicide formation and thermal stability on Si(1 0 0)[END_REF][START_REF] Detavernier | Thermal expansion of the isostructural PtSi and NiSi: Negative expansion coefficient in NiSi and stress effects in thin films[END_REF]. Detavernier et al. showed that the NiSi phase, with an orthorhombic structure type MnP, displayed the rare occurrence of contracting during heating at least along the axis with the smallest unit cell dimension. On the other hand, although the axis with the smallest cell contracted upon heating, the total volume of the unit cell actually increased for increasing temperatures. This, generated a considerable anisotropy in the thermal expansion behavior of the crystalline film. This anisotropic thermal expansion was also observed for NiGe by Gaudet et al.[START_REF] Gaudet | Thin film reaction of transition metals with germanium[END_REF]. NiGe phase shares the same crystalline orthorhombic crystalline structure with the following lattice parameters a = 0.545 nm , b = 0.3427 nm and c = 0.5885 nm. Then, for our case, b axis is expected to contract during annealing at high temperatures (after 450 ) and shift certain peaks related to specific orientations towards higher 2θ angles.

Figure 4 . 15 :

 415 Figure 4.15: In-plane RSM for (a) Ni / Ge and (b) NiCo / Ge systems. Patterns were acquired after anneal at 700 and cooling down to room temperature.

  4.15 while the respective integrated patterns are shown in Fig. 4.16.

Figure 4 . 16 :

 416 Figure 4.16: Integrated in-plane RSM in φ for both Ni / Ge and NiCo / Ge systems.

Figure 4 . 17 :

 417 Figure 4.17: Ex-situ XRD patterns measured after RTA for NiCo / Ge system.From the as-deposited state and up to 200 , TiN (1 1 1) and Ni (1 1 1) peaks were observed (at 2θ = 36.5°and 2θ = 44.4°respectively). At 250 , a peak appeared at 46.7°w hich can be related to the growth of the (NiCo) 5 Ge 3 phase. Later, after 350 , this NiCo-rich phase was consumed to the growth of the (NiCo)Ge phase. This phase remaind

Figure 4 . 18 :

 418 Figure 4.18: Surface AFM images (scan size 5 µm x 5 µm) of samples annealed at various temperatures. Ni / Ge (top row), Ni 0.9 Co 0.1 / Ge (bottom row).

Fig. 4 . 19 .

 419 Fig.4.[START_REF] Tong | Ni(Ge 1-x Sn x ) ohmic contact formation on n-type Ge 1-x Sn x using selenium or sulfur implant and segregation[END_REF]. At 500 a decrease on RMS roughness was observed for the NiCo / Ge system. Two explanations were considered for this behavior, (i) the incidence of an artifact or (ii) the formation of Co x Ge y transient phases (as explained in section 4.2.1.1) that could improve the surface states and promote a less rough surface.

Figure 4 . 19 :

 419 Figure 4.19: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for the Ni / Ge and NiCo / Ge systems.

Figure 4 . 20 :

 420 Figure 4.20: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature comparing Ni / Ge and NiCo / Ge systems.

Figure 4 . 21 :

 421 Figure 4.21: In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni 0.9 Co 0.1 and a Ge 0.90 Sn 0.10 layer.

  .

  3.

Table 4 . 3 :

 43 Critic threshold temperatures during the Ni(Co) / Ge(Sn) phase sequences System Ni(Co) Ni(Co)-rich Ni(Co)Ge(

Figure 4 . 22 :

 422 Figure 4.22: Patterns extracted from the in-situ XRD analysis of the Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 system.

  6). Thus, to highlight Co impact on the Sn segregation threshold, XRD analyses were systematically performed on samples annealed by RTA for 30 s in N 2 environment. The ex-situ XRD patterns are presented in Fig. 4.23.

Figure 4 . 23 :

 423 Figure 4.23: Ex-situ XRD patterns measured after RTA for the Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 system.

Figure 4 . 24 :

 424 Figure 4.24: Sn segregation threshold extracted from XRD analyses of samples annealed by RTA.

Figure 4 . 25 :

 425 Figure 4.25: In-plane RSM of the Ni 0.9 Co 0.1 / Ge 0.90 Sn 0.10 system. Patterns were acquired on samples annealed up to 700 and after cooling down to room temperature.

Figure 4 . 28 :

 428 Figure 4.28: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for the different NiCo / Ge 1-x Sn x systems.

Figure 4 . 29 :

 429 Figure 4.29: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the NiCo / Ge, Ni / Ge 0.90 Sn 0.10 and NiCo / Ge 0.90 Sn 0.10 systems.

Figure 4 . 30 :

 430 Figure 4.30: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for the NiCo / Ge, NiCo / Ge 0.94 Sn 0.06 , NiCo / Ge 0.90 Sn 0.10 and NiCo / Ge 0.85 Sn 0.15 systems.

Figure 4 . 31 :

 431 Figure 4.31: In-situ XRD patterns measured during the solid-state reaction between 10 nm of Ni 0.9 Pt 0.1 and a Ge 0.90 Sn 0.10 layer.

Figure 4 . 32 :

 432 Figure 4.32: Evolution of the angular position and FWHM as a function of temperature for the Ni 0.9 Pt 0.1 XRD peak located at about 43.9°in the Ni 0.9 Pt 0.1 / Ge 0.9 Sn 0.1 system.

Figure 4 . 33 :

 433 Figure 4.33: In-situ X-ray diffraction measurement of the reaction between NiPt and Ge 0.90 Sn 0.10 Focus on the (NiPt)(GeSn) peak shift (a) Zoom on the in-situ X-ray diffraction pattern in the 37-55°r egion, (b) Evolution of the angular position for the mono-stanogermanide peak located around 53.5°.

  4.34), the peak observed at 39.2°disappeared and a new peak appeared at 42.6°. The identification of these three peaks will be further explained in section 4.2.3.4.

Figure 4 . 34 :

 434 Figure 4.34: Evolution of the intensity of XRD peaks located at 39.2 and 42.0°in the Ni 0.9 Pt 0.1 /Ge 0.90 Sn 0.10 system with annealing temperature.

Figure 4 . 35 :

 435 Figure 4.35: Ex-situ XRD diagrams measured during the solid-state reaction between Ni 0.9 Pt 0.1 and a Ge 0.90 Sn 0.10 layer in the as-deposited and annealed by RTA.

Figure 4 . 36 :

 436 Figure 4.36: Low temperature in-situ XRD patterns (in the 20-35°angular range) acquired during the solid-state reaction between Ni 0.9 Pt 0.1 and Ge 1-x Sn x layers with Sn content ranging from 0 up to 15 at.%. Focus on the NiPt-rich phase peak.

4 . 32 )

 432 which corresponded to around 3% of dilatation and cannot be due to stress variations. This shift was then linked to a particular redistribution of Pt. Indeed, low incorporation of Pt in the Ni 2 Si phase and accumulation around the Ni-silicide / Si interfaces was already reported by Finstad et al. when investigating Si / Ni / Pt reaction [167] and by Lee et al. when studying the Ni 0.95 Pt 0.05 / Si reaction [150]. Cojocaru-Miredin et al. evidenced a redistribution of Pt during the silicide formation which generated Pt enrichment at both the Ni 1-x Pt x / Ni 2 Si and Ni 2 Si / NiSi interfaces.

  4.31 and Fig. 4.35. In order to verify Pt solubility in these phases TEM images coupled to EDS analyses were made for NiPt / Ge 0.90 Sn 0.10 samples annealed at 200 and 350 , for which the presence of the Ni(Pt)-rich phase and the Ni(Pt)-monostanogermanide phases were verified, respectively (see Fig. 4.35). The acquired maps are presented in Fig. 4.37.

Figure 4 . 37 :

 437 Figure 4.37: Cross-section TEM images coupled with EDS analysis of Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 samples after RTA at 200 (top row) and 350 (bottom row).

  3.41). In zone 3 the homogenous distribution of Sn in the Ni(GeSn) layer was verified. Meanwhile, Pt was almost entirely accumulated in the Ni(GeSn) top zone 4, in which a slight increase of Sn was also evidenced. Finally, zone 5 shows the TiN protective layer. The Pt snowplow effect and the low solubility of this element on both the Ni-rich and Ni-monostanogermanide 157 Chapter IV. Alternatives to improve the Ni / GeSn thermal stability phases were then verified.

Figure 4 . 38 :

 438 Figure 4.38: EDS profiles of Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 samples after RTA at 200 and 350 .

Figure 4 . 39 :

 439 Figure 4.39: In-plane RSM integrated patterns for Ni 0.9 Pt 0.1 /Ge 0.90 Sn 0.10 samples annealed at 400 and 600 .

Figure 4 . 40 :

 440 Figure 4.40: Cross-section TEM images coupled to EELS analysis of a Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 sample after RTA at 550 .

Fig. 4 .

 4 Fig. 4.40 shows the cartography obtained for a Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 sample after annealing at 550 . The TiN layer was rough and discontinuous due to the layers agglomeration wich was verified by AFM images in Fig. 4.42. The lost of continuity enabled a complete agglomeration of the Ni(GeSn) layer. Grains were mainly made of Ni and Ge atoms. Meanwhile, Sn was predominantly localized at the surface and at a grain boundaries together with Pt forming a PtSn x compound. However, the precision of the EDS analysis was not enough to obtain a proper identification of the stoichiometry of this compound.

Figure 4 . 41 :

 441 Figure 4.41: Sn segregation threshold extracted from XRD analyses of samples annealed by RTA, for Ni and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems.

Figure 4 . 42 :

 442 Figure 4.42: Surface AFM images (scan size 5 µm x 5 µm) of Ni 0.9 Pt 0.1 / Ge 1-x Sn x samples annealed at various temperatures by RTA.

  4.45 summarizes the impact of Co or Pt addition on the Ni consumption temperature and the Ni-rich and the mono-stanogermanide phases formation temperatures.

Figure 4 . 45 :

 445 Figure 4.45: Summary of the phase formation sequences obtained, thanks to in situ XRD analyses for Ni / GeSn, NiCo / GeSn and NiPt / GeSn systems.

5

 5 Ge 3 or NiGe should decrease the nucleation barriers. The phase sequences of Ni, Co and Pt germanides are depicted in Fig. 4.46.

Figure 4 . 46 :

 446 Figure 4.46: Phase sequences of the reaction of 30 nm-thick Ni, Co or Pt films on Ge (0 0 1). Adapted from [66].

Figure 4 . 47 :

 447 Figure 4.47: Sn segregation threshold extracted from XRD analyses of samples annealed by RTA, for Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems.

Figure 4 . 48 :

 448 Figure 4.48: Surface AFM images (scan size 5 µm x 5 µm) of Ni (top row), Ni 0.9 Co 0.1 (middle row) and Ni 0.9 Pt 0.1 (bottom row) / Ge 0.90 Sn 0.10 samples annealed at various temperatures by RTA.

Figure 4 . 49 :

 449 Figure 4.49: TEM images in bright field mode acquire for Ni / Ge 0.90 Sn 0.10 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 samples after anneals at 400 .For the Ni / Ge 0.90 Sn 0.10 sample, agglomeration was well verified at 400 . Grains consisting of Ni and Ge atoms were only observed completely agglomerated while Sn segregation occurred (refer to Fig.3.41). Meanwhile, the NiPt / Ge 0.90 Sn 0.10 layer remained smooth and continous. For these kinds of intermetallics, agglomeration starts with the grain boundary grooving and results in agglomeration, grains separation and island formation. The driven force for a film to agglomerate is the total surface / interface energy of the intermetallics and its substrate[START_REF] Gambino | Silicides and ohmic contacts[END_REF]. The impact of Pt on Ni stanogermanide is thus consistent with results obtained for Ni(Pt) / Si and Ni(Pt) / Ge systems[START_REF] Mangelinck | Enhancement of thermal stability of NiSi films on (1 0 0) Si and (1 1 1) Si by Pt addition[END_REF][START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF][START_REF] Zhang | Improvement of thermal stability of Ni germanide using a NiPt(1%) alloy on Ge-on-Si substrate for nanoscale Ge MOSFETs[END_REF][START_REF] Kang | Thermally robust Ni germanide technology using cosputtering of Ni and Pt for high-performance nanoscale Ge MOSFETs[END_REF]. NiGe and PtGe have the same orthorhombic crystalline structure. Thus, Pt could be soluble in the monogermanide phase, i.e. NiGe. Even if Pt solubility was observed to be low, some diffusion in the upper part of the Ni(GeSn) intermetallic was evidenced in Fig.4.37. Pt-Ge bond energy (400-500 kJ/mol) is higher than the NiGe bond energy (290 kJ/mol)[START_REF] Jin | Effect of Pt on agglomeration and Ge out diffusion in Ni(Pt) germanosilicide[END_REF]. As a result, Pt will reduce the breakdown rates of chemical bonds which will reduce the grain boundary grooving development and could explain the morphological stability improvement at high temperatures. This assumption on the grain boundary grooving kinetics is endorsed by AFM pictures. For the Ni / GeSn system, agglomeration is too fast to observe grain boundary grooving. On the other hand, due to Pt addition, the groove development being slower, the phenomenon is observed at 350

Figure 4 . 50 :

 450 Figure 4.50: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems.

This section 4 . 2

 42 issued to two publications, (i) in the Journal of Applied Physics, 124, 085305, (2018) and (ii) in the journal of Materials Science in Semiconductor Processing, 108, 104890, (2020).

4. 3

 3 UV-nanosecond laser annealing: Novel technique to anneal Ni / GeSn layersIt was observed in previous sections that either the use of pre-amorphization by ions implantation (PAI) prior to metallization or the use of an alloy element (Pt or Co) can be beneficial to improve certain aspects of the Ni / GeSn intermetallics. Indeed, the use of C PAI remarkably increased the Ni(GeSn) thermal stability and delayed Sn segregation up to 500-550 . Meanwhile, Co addition to the Ni metallization layer delayed Sn segregation up to 400 and Pt addition delayed both, agglomeration and Sn segregation up to 450 . However, another alternative can be considered in order to improve the thermal stability of the Ni / GeSn intermetallics by means of laser annealing.

Figure 4 . 51 :

 451 Figure 4.51: Numerical simulations concerning critical energy thresholds (a) Melting threshold for TiN / Ni / Ge samples, (b) Morphological variation onset simulations corresponding to the Ge melt threshold.

Figure 4 . 52 :

 452 Figure 4.52: Scheme of the annealed samples (a) Ni / Ge and Ni / GeSn stacks, (b) mapping of the laser annealing shots with different conditions on a 200 mm wafer.

  4.54.

Figure 4 . 53 :

 453 Figure 4.53: Time-resolved reflectometry (TRR) for different energy densities during laser annealing of (a) Ni / Ge and (b) Ni / GeSn samples.

Figure 4 . 54 :

 454 Figure 4.54: TRR intensity average for Ni / Ge and Ni / GeSn samples, before and after the laser annealing treatment.

Figure 4 . 55 :

 455 Figure 4.55: Evolution of the reflection % at different wavelengths for Ni / Ge and Ni / GeSn samples.

Figure 4 . 57 :

 457 Figure 4.57: Ex-situ XRD patterns for a Ni / GeSn sample laser annealed by 10 laser pulses for 0.32, 0.48 and 0.64 J/cm 2 energy densities.

Figure 4 . 58 :

 458 Figure 4.58: Ex-situ XRD patterns for Ni / GeSn samples comparing the phase sequence on (a) RTA samples and (b) laser annealed samples by 1 laser pulse.

Figure 4 . 59 :

 459 Figure 4.59: Ex-situ XRD patterns for samples laser annealed couples with RTA at 200 for (a) Ni / Ge and (b) Ni / GeSn systems.

Figure 4 . 61 :

 461 Figure 4.61: Root-mean-square surface roughness evolution comparing Ni / Ge and Ni / GeSn systems annealed by 1 laser pulse.

Figure 4 . 62 :

 462 Figure 4.62: Schematic representation of wrinkles formation on the TiN capping layer.

  4.63.

Figure 4 . 63 :

 463 Figure 4.63: Surface SEM images (scan size 30 µm x 30 µm) of Ni / GeSn samples laser annealed with 1 laser pulse (top row), 3 laser pulses (middle row) and 10 laser pulses (bottom row).

Figure 4 . 64 :

 464 Figure 4.64: Surface AFM images (scan size 5 µm x 5 µm) of Ni / GeSn samples. RTA only (top row), laser annealing (bottom row).

Figure 4 . 65 :

 465 Figure 4.65: Root-mean-square surface roughness evolution for Ni / GeSn system annealed by RTA or by 1 shot laser pulse.

Figure 4 . 66 :

 466 Figure 4.66: Surface AFM images (scan size 5 µm x 5 µm) of Ni / GeSn samples. RTA only (top row), laser annealing (middle row) and laser annealing couple to RTA at 200 (bottom row).

Figure 4 . 67 :

 467 Figure 4.67: Comparison of sheet resistance evolution for Ni / Ge and Ni / GeSn systems with 1 pulseshot in a range from 0.08 up to 0.66 J/cm 2 of energy densities.

Figure 4 . 68 :

 468 Figure 4.68: Evolution of the sheet resistance as a function of the laser energy densities for the Ni / GeSn system when either 1, 3 or 10 pulses when realized in a range from 0.08 up to 0.66 J/cm 2 .

Figure 4 . 70 :

 470 Figure 4.70: Evolution of the FWHM values collected for the the Ni-monostanogermanide peak around 53.5°from various XRD measurements for samples annealed by RTA or laser annealing.

Figure 4 . 71 :

 471 Figure 4.71: Comparison of sheet resistance evolution for Ni / Ge and Ni / GeSn samples, when annealed by 1 pulse laser shot or 1 pulse laser shot coupled to RTA at 200 .

Figure 4 . 72 :

 472 Figure 4.72: Scheme of the techniques analyzed in this chapter to increase the Ni / GeSn thermal stability.

1 5 . 6

 156 Ti / Ge system phase sequence . . . . . . . . . . . . . . . . . . . . . . 5.2 Ti / GeSn system phase sequence . . . . . . . . . . . . . . . . . . . . 5.3 Discussion: Ti / Ge and Ti / GeSn systems comparison . . . . . . . . 5.4 Ti / Ge and Ti / GeSn systems morphological evolution . . . . . . . . 5.5 Ti / Ge and Ti / GeSn systems electrical evolution . . . . . . . . . . Conclusions of Chapter V . . . . . . . . . . . . . . . . . . . . . . . . .

  5.1.

Figure 5 . 1 :

 51 Figure 5.1: Schematic representation of chapter IV section division.

Figure 5 . 2 :

 52 Figure 5.2: In-situ XRD patterns measured during the solid-state reaction between 20 nm of Ti and a Ge layer.

Figure 5 . 3 :

 53 Figure 5.3: Evolution of the position and intensity as a function of temperature for the Ti (0 0 2) peak located around 38.4°for the Ti / Ge system.

Figure 5 . 4 :

 54 Figure 5.4: Ex-situ XRD patterns measured during the solid-state reaction between 20 nm of Ti and a Ge layer, after RTA during 30 s in N 2 environment.

  Chapter V. Alternative to Ni-based contacs Ti metallization analysis (see Fig.5.2 and 5.3) and can be related to a phase transition. This shift can be better observed in Fig.5.5.

Figure 5 . 5 :

 55 Figure 5.5: Ex-situ XRD patterns measured during the solid-state reaction between 20 nm of Ti and a Ge layer, after RTA during 30 s on N 2 environment at 400 and 450 .

Figure 5 . 6 :

 56 Figure 5.6: In-situ XRD patterns measured during the solid-state reaction between 20 nm of Ti and a Ge 0.90 Sn 0.10 layer.

  5.1.

Figure 5 . 7 :

 57 Figure 5.7: Evolution of the position and intensity as a function of temperature for the Ti (0 0 2) peak located around 38.4°for the Ti / GeSn system.

Figure 5 . 8 :

 58 Figure 5.8: Summary of the phase formation sequences obtained, thanks to in situ XRD analyses for a 20 nm of Ti deposited on Ge and Ge 0.90 Sn 0.10 layers.

Figure 5 . 9 :

 59 Figure 5.9: Ex-situ XRD patterns measured during the solid-state reaction between 20 nm of Ti and a Ge 0.90 Sn 0.10 layer, after RTA during 30 s in N 2 environment.

205

  Chapter V. Alternative to Ni-based contacs Ti metallization 5.3 Discussion: Ti / Ge and Ti / GeSn systems comparison

Figure 5 . 10 :

 510 Figure 5.10: Ti-Ge binary phase diagram [217].

Figure 5 . 11 :

 511 Figure 5.11: Surface AFM images (scan size 5 µm x 5 µm) of Ti / Ge (top row) and Ti / GeSn (bottom row), from samples annealed at various temperatures by RTA.

Figure 5 . 12 :

 512 Figure 5.12: Surface AFM images (scan size 5 µm x 5 µm) of Ti / Ge (top row), Ti / GeSn (middle row) and Ni / GeSn (bottom row), from samples annealed at various temperatures by RTA.

Figure 5 . 13 :

 513 Figure 5.13: Evolution of the root-mean-square roughness values as a function of the rapid thermal annealing temperature for Ti / Ge, Ti / GeSn and Ni / GeSn systems.

5. 5

 5 Ti / Ge and Ti / GeSn systems electrical evolutionThe sheet resistance as a function of RTA temperatures was acquired for different samples to compare the Ti / Ge and GeSn-based systems evolution and Ni and Ti-based samples. Fig.5.14 shows the results obtained for Ti-based samples.

Figure 5 . 14 :

 514 Figure 5.14: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for Ti / Ge and Ti / GeSn systems.

Figure 5 . 15 :

 515 Figure 5.15: Evolution of the sheet resistance as a function of the rapid thermal annealing temperature for Ti / GeSn and Ni / GeSn systems.

Figure 6 . 1 :

 61 Figure 6.1: Schematic representation of the main results for Ni / GeSn samples (a) phase sequence during the solid-state reaction, (b) atomic disposition of the hexagonal Ni 5 (GeSn) 3 phase, (c) segregated Sn acting as stuffed barrier and (d) difference in texture for NiGe and Ni(GeSn) phases.

Figure 6 . 2 :

 62 Figure 6.2: R sh evolution for the Ni / Ge 0.90 Sn 0.10 sample with annealing temperatures correlating the morphological evolution of the layers.

Figure 6 . 3 :

 63 Figure 6.3: Schematic representation of the main alternatives studied to improve the Ni / GeSn thermal stability (a) pre-amorphization by implantation (PAI), (b) addition of Co or Pt as an alloy element and (c) use of laser pulses as thermal annealing.

Figure 6 . 4 :

 64 Figure 6.4: Schematic representation of the process steps that can be consider to ensure high quality, stable and reliable GeSn-based contacts.

Figure 6 . 5 :

 65 Figure 6.5: Schematic representation of (a) the process flow to obtain the TLM and photodiodes samples, (b) the TLM geometrical parameters and (c) image of the measured photodiodes.

Figure 6 . 6 :

 66 Figure 6.6: Cut-off frequency for NiPt and Ti/Au metallized photodiodes.

Figure A. 1 :

 1 Figure A.1: NiGe phase powder diffraction PDF chart (04-017-6384) from the ICDD data base.

F. 1

 1 Résumé du Chapitre I. Le GeSn: de la croissance aux applications L'alliage GeSn correspond à un matériau du groupe IV qui présente une bande interdite directe qui est avantageuse pour différentes applications pour l'industrie des semiconducteurs. Effectivement, l'introduction de Sn dans la maille de Ge permet d'obtenir un matériaux à bande interdite directe comme observé dsur la Fig. F.1.

Figure F. 1 :F. 3

 13 Figure F.1: GeSn transition de bande interdite de bande indirecte à bande directe.

Figure F. 2 :F. 4 . 2

 242 Figure F.2: Diagrammes DRX in-situ mesurés pendant la réaction à l'état solide entre 10 nm de Ni et un couche de Ge 0.90 Sn 0.10 .

  F.5.

Figure F. 5 :

 5 Figure F.5: DRX ex-situ pour les échantillons recuits par 1 impulsion laser pour le système Ni / GeSn.

  F.6.

Figure F. 6 :

 6 Figure F.6: Diagrammes DRX in-situ mesurés lors de la réaction à l'état solide entre 20 nm de Ti et une couche de Ge 0.90 Sn 0.10 .

  Titre : D éveloppement et caract érisation de contacts sur GeSn Mots cl és : Ni, GeSn, r éaction a l' état solide, XRD, AFM, R sh R ésum é : Des efforts dans l'industrie des semiconducteurs sont constamment d éploy és afin d'am éliorer diff érents param ètres comme les performances des dispositifs ou la vitesse de transfert des donn ées. Pour atteindre ces avanc ées, des voies innovantes peuvent être consid ér ées telles que la modification des étapes des proc éd és de production, l'architecture des dispositifs ou les mat ériaux qui constitueront ces dispositifs. Le germanium-étain (GeSn), alliage du groupe IV, est un mat ériau int éressant à int égrer dans des dispositifs électroniques ou optoélectroniques. Le GeSn peut être utilis é comme stresseur des source et drain dans les MOSFET Ge (metal-oxide-semiconductor field-effect transistors) et pour obtenir des canaux à haute mobilit é dans les pMOSFET et les pTFET (thin field-effect transistors). D'autre part, l'ajout d'une quantit é suffisante de Sn dans le r éseau du Ge (environ 10 at.%) offre la possibilit é d'obtenir une structure à bande interdite directe. Ce mat ériau pourrait alors être utilis é pour r éaliser un laser du groupe IV int égr é de fac ¸on monolithique et compatible avec la technologie Si-CMOS (complementary metal-oxide-semiconductor). Quelle que soit l'application, des contacts m étalliques ou interm étalliques à faible r ésistance, stables et fiables sont des composants cl és pour injecter le courant électrique dans les dispositifs. Les interm étalliques Ni / GeSn ont ét é consid ér és comme un mat ériau de contact adapt é. Cette th èse était donc consacr ée au l' étude syst ématique compl ète et à la caract érisation d'interm étalliques Ni / GeSn pour contacter des dispositifs bas és sur GeSn. Les propri ét és des couches de type Ni / GeSn analys ées en termes de la s équence de phases, de l' évolution morphologique et électrique au cours de la r éaction à l' état solide, sont pr ésent ées en premier. Ensuite, diff érentes alternatives seront d écrites pour am éliorer la stabilit é thermique des contacts à base de Ni : l'utilisation de la pr é-amorphisation par implantation (PAI), l'addition d'un él ément d'alliage (Co, Pt) et l'utilisation de l'impulsion laser comme technique de recuit. Une autre alternative concernant la m étallisation de GeSn à base de Ti est également mentionn ée. Gr âce à ces études, une analyse compl ète et syst ématique du syst ème Ni / GeSn a ét é r éalis ée. De plus, l'identification de diff érentes alternatives pour modifier les conditions du processus qui peuvent am éliorer la stabilit é thermique du syst ème Ni / GeSn a ét é r éalis ée. Les r ésultats obtenus repr ésentent un bon point de d épart pour élaborer des contacts de haute qualit é, stables et fiables sur GeSn, qui peuvent être enti èrement int égr és dans des dispositifs électroniques ou opto électroniques. alliages GeSn Title : Development and characterization of GeSn-based contacts Keywords : Ni, GeSn, Solid-state reaction, XRD, AFM, R sh

  Espace Technologique / Immeuble Discovery Route de l'Orme aux Merisiers RD 128 / 91190 Saint-Aubin, France contacts on GeSn alloys Título : Desarrollo y caracterizaci ón de contactos sobre GeSn Palabras claves : Ni, GeSn, reacci ón en estado s ólido, XRD, AFM, R sh

Table 1 . 1 :

 11 Growing parameters of Ge 1-x Sn x layers

Table 1 . 2 :

 12 Non exhaustive recapitulative table of the metallization parameters for GeSn layers (N/A stands for Not Available)

	Reference	Surface	Deposition	Metal	Thermal
		treatment	technique	layer	treatment
		HFH 2 O 150 +	Co-sputtering in		350-550
	[9]	deionized water	a magnetron	6 nm Ni/Pt	30 s
			sputtering system		N 2
			Electron beam		250-400
	[12]	N/A	evaporation	30 nm Ni	30 s
			Electron beam		350
	[10]	N/A	evaporation	20 nm Ni	30 s
					N 2
					200-450
	[11]	N/A	PVD	20 nm Ni	30 s
		HFH 2 O 1100 +			350-550
	[13]	deionized water	Sputtering	8 nm Ni	30 s
					N 2
			Electron beam		400-600
	[14]	N/A	evaporation in	10 nm Ni	60-240 min
			UHV		N 2
					360-550
	[15]	HF	Evaporation tool	15 nm Ni	30 s
					N 2
	[16]	N/A	Sputtering	50 nm Ni	20-850 0.5 /s
					He
			Electron beam		300-500
	[17]	N/A	evaporation	30 nm Ni	30 s
					N 2
					250-400
	[18]	HF (1 % of water)	N/A	10 nm Ni	30 s
					N 2 H 2 /91
					350
	[19]	N/A	N/A	10 nm Ni	30 s
					N 2
		Diluted HF +	Electron beam		350-550

  Study of the Ni / Ge 0.90 Sn 0.10 system at low temperature . . Nature of the Ni-rich phase . . . . . . . . . . . . . . . . . . . Impact of Sn content on the Ni / GeSn system evolution . . . . . . . Sn behavior during the Ni / GeSn solid-state reaction . . . . . 3.6.5 Incorporation of Sn in the intermetallics . . . . . . . . . . . . 3.6.6 Impact of Sn on the Ni / GeSn phase sequence during solidstate reaction . . . . . . . . . . . . . . . . . . . . . . . . . . . Specific case: the Ni / Ge 0.85 Sn 0.15 system . . . . . . . . . . . Study of the out-of-plane and in-plane orientations of the NiGe and Ni(GeSn) phases: Results . . . . . . . . . . . . . . . . . . 3.7.2 Discussion: NiGe and Ni(GeSn) texture . . . . . . . . . . . . 3.8 Conclusions of Chapter III . . . . . . . . . . . . . . . . . . . . . . . .

	3.5.2 3.6 3.6.7 3.7 Ni(GeSn) crystallographic texture: Impact of Sn content . . . . . . . 3.7.1

Contents 3.1 State of the art: Ni / GeSn properties and evolution with annealing temperatures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1.1 Ni / GeSn phase sequence . . . . . . . . . . . . . . . . . . . . 3.1.2 Ni / GeSn morphological properties and stability . . . . . . . 3.1.3 Ni / GeSn electrical properties and stability . . . . . . . . . . 3.2 Experimental procedure . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1 Growth and characterization of GeSn layers . . . . . . . . . . 3.2.2 Ni / GeSn samples preparation . . . . . . . . . . . . . . . . . 3.3 General characterization of the Ni / Ge 0.90 Sn 0.10 system . . . . . . . . 3.3.1 The Ni / Ge 0.90 Sn 0.10 phase sequence . . . . . . . . . . . . . . 3.3.2 The Ni / Ge 0.90 Sn 0.10 morphological evolution . . . . . . . . . 3.3.3 The Ni / Ge 0.90 Sn 0.10 electrical evolution . . . . . . . . . . . . 3.4 Effect of surface preparation on the Ni / GeSn properties evolution during the solid-state reaction . . . . . . . . . . . . . . . . . . . . . . 3.4.1 Effect of surface preparation on the Ni / GeSn phase sequence 3.4.2 Effect of surface preparation on the morphological evolution . 3.4.3 Effect of surface preparation on the electrical evolution . . . . 3.5 Ni-rich phase crystallographic structure determination . . . . . . . . . 3.5.1 3.6.1 Sn content variation on the phase sequence, the morphological and the electrical properties evolution . . . . . . . . . . . . . . 3.6.2 Focus on the Ni-rich phase . . . . . . . . . . . . . . . . . . . . 3.6.3 Depth profile analysis cross-sectional TEM together with EELS and EDS mappings . . . . . . . . . . . . . . . . . . . . . . . . 3.6.4

Table 3 .

 3

	Phase	Composition	System	Space Group
		at.% Ge
	Ni	0 to 12	Cubic	Fm3m
	β-Ni 3 Ge	22.5 to 25	Cubic	Pm3m
	γ-Ni 3 Ge	25.6	...	...
	δ-Ni 5 Ge 2 28	Hexagonal	P6 3 /mmc
	Ni 2 Ge	33.5	Orthorhombic Pnma
	ε'-Ni 5 Ge 3 ∼37	Monoclinic	C2
	ε-Ni 5 Ge 3 33.8 to 38	Hexagonal	P6 3 /mmc
	Ni 19 Ge 12 38 to 41	Monoclinic	C2
	Ni 3 Ge 2	41 to 43	Hexagonal	P6 3 /mmc
	NiGe	50	Orthorhombic Pnma
	Ge	100	Cubic	Fd 3 m

1: Ni-Ge Crystal Structure Data (reproduced and adapted from

[START_REF] Nash | The Ge-Ni (germanium-nickel) system[END_REF]

)

Table 3 . 2 :

 32 Critic threshold temperatures during the Ni / GeSn phase sequence for the reference samples and samples treated by different surface treatments. Red color highlighted an increase on temperature compared to the reference while blue color highlighted a decrease of temperature

	System	Ni	Ni-rich	Ni(GeSn)	Ni(GeSn)
		consumption	growth	growth	destabilization
	Reference

Table 3 .

 3

	System	Angle 23.7°bulk surface Angle 76.2°near surface
		Ge 2p Ox% Sn 3d Ox% Ge 2p Ox% Sn 3d Ox%
	Reference	77.0	56.7	92.8	90.2
	Ar	6.5	6.1	20.5	39.8
	He	23.2	24.7	25.3	51.3

3: Summary of the oxides proportions measured for 23.7 ± 3.7°and 76.2 ± 3.7°angles collection corresponding to the bulk and the near surface responses from the photoelectrons

  Nishimura et al. found an epitaxial Ni 2 (GeSn) formation in samples with an Sn content of 3.6%, while different studies reported that the Ni-rich phase corresponded to the Ni 5 (GeSn) 3 phase[START_REF] Liu | Reduction of formation temperature of nickel mono-stanogermanide [Ni(GeSn)] by the incorporation of tin[END_REF][START_REF] Demeulemeester | Sn diffusion during Ni germanide growth on Ge 1-x Sn x[END_REF][START_REF] Wirths | Ni(SiGeSn) metal contact formation on low bandgap strained (Si)Ge(Sn) semiconductors[END_REF][START_REF] Wirths | Ternary and quaternary Ni(Si)Ge(Sn) contact formation for highly strained Ge p-and n-MOSFETs[END_REF]. In addition, Schulte-Braucks et al. confirmed the presence of Ni 5 (GeSn) 3 peaks with the appearance of supplementary peaks that could be attributed to Ni 2 (GeSn) or Ni 3 (GeSn) 2 phases. As part of our systematic investigation of Ni / GeSn solid-state reaction, we performed a detailed analysis of the nature of the Ni-rich phase using in-plane reciprocal space map (RSM) and in-plane pole figures. The complementary in-plane RSM and the in-plane pole figures were carried out following the parameters described in chapter II, sections 2.1.2 and 2.1.3.

Table 3 . 7 :

 37 Summary of the orientations observed for the Ni(GeSn) phase. Each symbol corresponds to an orientation in Fig.3.54

Table 4 . 2 :

 42 FWHM values collected for the Ni-monostanogermanide peak around 53.5°from various in-situ XRD measurements for the reference system and systems with different PAI

		T ()		FWHM (°) of the Ni(GeSn) peak
	[-0.8ex]		Reference Si PAI Ge PAI C PAI Ge + C PAI
		350	0.23	0.28	0.28	0.39	0.62
		365	0.23	0.23	0.31	0.39	0.39
	In-situ	390	0.31	0.28	0.38	0.62	0.85
	XRD	420	0.23	0.31	0.31	0.62	0.54
		435	0.23	0.23	0.23	0.47	0.57
		450	0.23	0.31	0.23	0.54	0.47
		500	0.31	0.28	0.38	0.38	0.57
	plantation of C ion into the layer, which was clearly beneficial for delaying agglomeration
	and enhancing the morphological stability of the layers.

Table 4 . 4 :

 44 Critic threshold temperatures during the Ni(Pt) / Ge(Sn) phase sequences

	System	NiPt	NiPt-rich	(NiPt)(GeSn)
		consumption consumption	shift
	NiPt / Ge	140	165	-
	NiPt / Ge 0.94 Sn 0.06	140	210	405
	NiPt / Ge 0.90 Sn 0.10	140	240	350
	NiPt / Ge 0.85 Sn 0.15	140	245	305

Table 4 .

 4 5.

Table 4 . 5 :

 45 Enthalpies of formation of the Pt-Sn[START_REF] Anres | Enthalpy of formation of the (PtSn) system[END_REF], Ni-Sn[START_REF] Nash | The Ni-Sn (Nickel-Tin) System[END_REF] and Pt-Ge[START_REF] Lbibb | Calorimetric determination of the enthalpies of formation of the solid compounds of the Pt-Ge system[END_REF] systems

	Phase x P t (atomic fraction)	f orm H m (kJ/mol)
	Pt 3 Sn	0.75	-50.2
	PtSn	0.5	-58.6
	Pt 2 Sn 3	0.4	-54.4
	PtSn 2	0.33	-52.3
	PtSn 3	0.25	Metastable
	PtSn 4	0.2	-27.2
	Pt 3 Ge 2	0.6	-42.4
	PtGe	0.5	-45.0
	Pt 2 Ge 3	0.4	-41.9
	PtGe 2	0.33	-32.4
	Ni 3 Sn 4	0.43	-25.3
	Ni 3 Sn 2	0.6	-31.8
	Ni 3 Sn	0.75	-26.3
	NiSi	0.5	Metastable

Table 4 . 6 :

 46 FWHM values collected for the Ni-monostanogermanide peak around 53.5°from in-situ XRD measurements for Ni, Ni 0.9 Co 0.1 and Ni 0.9 Pt 0.1 / Ge 0.90 Sn 0.10 systems

		T () FWHM (°) of the Ni(GeSn) peak
			Ni / Ni 0.9 Co 0.1	Ni 0.9 Pt 0.1
		250	0.39	0.47	0.47
		260	0.31	0.31	0.47
	In-situ	270	0.31	0.31	0.39
	XRD	280	0.31	0.39	0.39
		290	0.31	0.39	0.39
		300	0.31	0.39	0.39
		340	0.23	0.31	0.39
		350	0.23	0.31	0.47
		400	0.31	0.31	0.31
		450	0.23	0.23	0.31
		500	0.31	0.28	0.47

Table 4 . 7 :

 47 FWHM values collected for the Ni-monostanogermanide peak around 53.5°from various XRD measurements for samples annealed by RTA or laser annealing

	RTA ()	Peak position (°) FWHM (°)
	350	32.4	0.15
	400	32.4	0.21
	450	32.4	0.18
	Energy densities (J/cm 2 ) Peak position (°) FWHM (°)
	0.38	32.3	0.16
	0.42	32.4	0.21
	0.48	32.6	0.47
	0.50	32.6	0.47
	0.52	32.5	0.42
	0.60	32.4	0.37
	0.66	32.4	0.37

Table 5 . 1 :

 51 Critical threshold temperatures during in-situ XRD analyses of the Ti / Ge(Sn) phase sequences Ge(Sn) 3 Ti 6 Ge(Sn) 5 TiGe(Sn) 2

	System	Ti T 5 shift consumption formation Ti	formation formation
	Ti / Ge	350 390	390	470	-
	Ti / GeSn 260 350	260	530	835

Table 5 . 2 :

 52 Enthalpies of formation of the different Ti-Ge compounds

	Phase	H f (kJ/mol)
	Ti 5 Ge 3	-69.8 [218]
	Ti 6 Ge 5	-70.1 [218]
	TiGe 2	-39.2 [219]

Table 6 .

 6 1: Contact resitance and resistivity acquired from TLM method for the different samples analyzed

	System	Type	Contact	Contact
		of contact resistance (Ω) resistivity (Ω.cm 2 )
	NiPt	Ohmic	1.6	4.9x10 -6
	Ti/Au	Ohmic	6.3	3.9x10 -4

Le processus d'élaboration d'un dispositif, dans l'industrie des semiconducteurs, peut exiger de nombreuses étapes, dans lesquelles la qualité des matériaux intégrés devrait être haute. Afin de garantir que tous les composants répondent aux exigences standard, les propriétés et les performances doivent être évaluées et mesurées. La façon d'atteindre ces objectif est d'utiliser différentes techniques de caractérisations qui peuvent permettre de comprendre les propriétés des matériaux. Un contact métal / semiconducteur (M / S) est un composant fondamental de nombreux dispositifs électroniques et opto-électroniques. En effet, cette jonction M / S permettra l'échange de charges électroniques entre le semiconducteur et un circuit externe. Pour ces raisons, l'étude, la caractérisation et l'amélioration des contacts ont fait l'objet de recherches actives depuis des décennies[START_REF] Brillson | Contacts to Semiconductors: Fundamentals and Technology[END_REF]. Du point de vue des dispositifs, la fiabilité, la stabilité et les performances électriques élevées du contact métallique sont essentielles pour la micro / nano et l'opto-électronique. Une approche globale nécessite une compréhension complète de l'interface, au niveau physique et chimique, et son évolution lors de la modification de certains paramètres du procédé (température, élément d'alliage, traitements de surface, etc.). Pour atteindre

evolution was observed as for the NiCo / Ge system. The roughness evolution with RTA was also acquired for NiCo / Ge 0.94 Sn 0.06 and NiCo / Ge 0.85 Sn 0.15 systems and is presented in Fig. 4.28 for comparison. From Fig. 4.28 it can be observed that, in general, the morphological evolution is comparable. The RMS roughness values remained low and stable up to a threshold temperature at which they increased due to the layers degradation. Differences were the grow of the mono-germanide phase occurred at lower temperatures (see Fig. 4.36).

The Ni-germanide surface slightly agglomerated after 550 , the root-mean-square values remained quite low, however (see Fig. 4.43). For the NiPt / Ge 0.94 Sn 0.06 and / Ge 0.90 Sn 0.10 systems the cross-hatch pattern was preserved up to 400 . At this temperature, signs of grain boundary grooving were observed and further agglomeration was verified at 500-550 . Meanwhile for the NiPt / Ge 0.85 Sn 0.15 sample, the cross-hatch pattern was preserved up to 350 while agglomeration occurred at 400 . Finally at 550 agglomeration was verified for all the samples.

In order to better visualize the effect of Sn content variation of root-mean-square (RMS) roughness values were extracted form the AFM images, the evolution is presented in Fig. 4.43. Sn segregation for NiPt / GeSn systems. Indeed, a relationship between Sn content and degradation was noticed for Ni / GeSn systems in chapter III. The RMS roughness values increase was more important as the Sn content further increased. This reinforced the hypothesis that the main reason for surface morphology degradation corresponds to the agglomeration process which will be aggravated by the Sn segregation.

NiPt / GeSn electrical evolution with temperature

The R sh evolution with annealing temperature was acquired with a four-point probe meter. The objective was to verify the impact of the phase sequence and morphology evolution with temperature on the electrical properties, and verify the process window at which values remained low, stable and in acceptable limits for integration. Fig. 4.44 shows the evolution of the sheet resistance, for NiPt systems with different Sn contents and after various RTA from 150 up to 550 . From Fig. 4.56 it can be observed that for both kinds of samples (Ni / Ge and Ni / GeSn), phase sequence evolution was comparable. At low energy densities (ED), Ni (1 1 1) at 44.7°together with TiN (1 1 1) at 36.4°peaks were present. After a threshold energy a peak around 46.5°related to the Ni-rich phase, (Ni 5 Ge 3 or Ni 5 (GeSn) 3 ), appeared while the Ni layer was consumed. Later on, the Ni-rich phase was consumed to the benefit of the NiGe or Ni(GeSn) phases. However, the threshold energies for phase formation varied from one system to the other. Indeed, the Ni 5 Ge 3 phase was observed at 0.36 up to 0.46 J/cm 2 while the Ni 5 (GeSn) 3 phase appeared since 0.32 up to 0.42 J/cm 2 . Then, the NiGe phase grew at 0.46 J/cm 2 compared to the Ni(GeSn) one, which appeared at 0.38 J/cm 2 , both peaks were positioned around 32.4°. The growth of the intermetallic phases can explain the slight decrease on reflectometry observed around ED > 0.4 J/cm 2 for Ni / Ge system and around 0.36 J/cm 2 (see Fig. 4.53 and 4.54).

The differences concerning the threshold energies for phase formation can be thus related to the differences in reflection/absorption of the beam energy from each sample. Ni / GeSn samples would reach the melting threshold at lower energy densities than Ni / Ge system, which allowed the phase growth to occur at lower ED. The effect of realizing multi-pulses, on the phase sequence, for the Ni / GeSn system was also verified. XRD patterns from Ni / GeSn samples treated with 10 laser pulses of 0.32, 0.48 and 0.64 J/cm 2 is shown in Fig. 4.57.

followed by RTA at 200 was verified to be efficient to nucleate the Ni-rich phase with low thermal budgets for the Ni / Ge and Ni / GeSn systems. The growth of the monogermanide or mono-stanogermanide phase was however not shifted towards lower energy densities. Indeed, the growth of the Ni(Ge) and Ni(GeSn) phases where observed after 0.46 J/cm 2 and 0.42 J/cm 2 , respectively.

Ni / Ge and Ni / GeSn morphological evolution

In order to verify the surface stability (as realized previously in different sections) and follow the morphological evolution with the energy density of the laser shot, atomic force microscopy (AFM) images with a Bruker FastScan equipment were acquired. Whatever the system studied and up to a certain energy density, (i.e. ED < 0.38 J/cm 2 for Ge-based and ED < 0.32 J/cm 2 for GeSn-based systems), the samples exhibited a cross-hatch pattern along the 1 1 0 crystallographic direction, as observed in chapter III, section 3.2.1. Then, low energy pulses did not modify the original surface morphology. After 0.38 and 0.34 J/cm 2 , for Ge-based and GeSn-based systems, respectively, the surface morphology changed. Wrinkles were observed in the surface together with an increase on the root-mean-square (RMS) roughness values(see Fig. RMS roughness values increased at lower ED for the Ni / GeSn system compared to the Ni / Ge system. This behavior was linked to the fact that the energy absorption was greater for Ni / GeSn stacks (see Fig. 4.55). Melting can thus occur at lower energies concomitantly with the growth of the intermetallic phases (see Fig. 4.56). However, once the threshold energy for surface morphology changes was passed, for both samples, the surface RMS roughness increase was comparable.

Wrinkles formation was related to a variation of strain in the capping TiN layer when the underneath layer Ni / Ge or Ni / GeSn was melt. In our case, the TiN capping layer presents a stress of -6 GPa. Then, when the Ni / Ge or Ni / GeSn layer underneath reached the metling threshold the TiN released the accumulated stress through deformation by wrinkles formation. and later decreased thanks to the growth of the mono-stanogermanide (Ni(GeSn)) phase. Minimum values around 0.48-0.50 J/cm 2 were encountered for all the samples after. Then, R sh values importantly increased. This increase was related to the variations in surface roughness due to wrinkling. The increase was more important for 3-pulses and 10-pulses samples compared to the use of a single pulse. This behavior can be related to the higher wrinkles amplitude when the number of pulses increased, as observed in Fig. 4.63. The electrical evolution of samples annealed by laser pulses was compared to RTA samples in Fig. 4.69. R sh remained stable and around 30 Ω/sq until the Ni-rich phase was obtained (around 250 for RTA samples, 0.32-0.34 J/cm 2 for laser annealing). Later on, R sh values decreased thanks to the growth of the Ni(GeSn) phase, at 350 for RTA samples and around 0.38 J/cm 2 for laser annealed samples. The minimum values for both kinds of thermal treatments were also comparable, for RTA samples the minimum Ni-mono-stanogermanide (Ni(GeSn)) phase. It happened concomitantly with the surface morphology changes observed by AFM images in Fig. 5.12 and related to agglomeration and Sn segregation. After agglomeration and Sn segregation occurred, the R sh importantly and progressively increased for the Ni / GeSn system. Meanwhile, the Ti / GeSn system showed stable values at 450 and beyond. Even if the RMS roughness increased (see Fig. 5.13), the R sh values remained stable. Then a higher electrical stability for Ti-based samples was evidenced even if the morphology varied.

NiGe phase powder piffraction PDF chart from the ICDD data base

Ni / Ge 0.94 Sn 0.06 and Ni / Ge 0.85 Sn 0.15 systems pole figures

Appendix C

In-situ XRD patterns for different PAI samples Les échantillons tels que déposés montrent un motif de cross-hatch le long des directions cristallographiques 1 1 0 . Ce motif résulte de la propagation de dislocations à 60°s ur le plan (1 1 1). Cette morphologie de surface provient directement de l'épitaxie de la couche de Ge / GeSn. Cette morphologie de surface est conservée jusqu'à 300 . À 350 , la morphologie a changé radicalement à cause de l'agglomération de la couche de Ni(GeSn). Ce phénomène est aggravé à cause de la ségrégation de Sn. À des températures plus élevées, la morphologie continue d'évoluer avec une augmentation de la rugosité. F.4 Résumé du Chapitre IV. Amélioration de la stabilité thermique de la phase Ni(GeSn)

Dans ce chapitre, différentes options visant à retarder l'agglomération et la ségrégation de Sn tout en garantissant de faibles valeurs de R sh seront présentées.

F.4.1 Pré-amorphisation par implantation ionique

Une première alternative correspond à la modification de la couche de GeSn avant la métallisation du Ni par implantation ionique, plus précisément, pré-amorphisation par implantation (PAI). Il s'agit d'une technique hors équilibre qui permet d'amorphiser la couche semiconductrice. La séquence de phase est comparable à celle obtenue pour le système Ni / GeSn de référence. Des différences ont cependant été constatées concernant les températures de transformation des phases (voir Tableau F.1).

Globalement, la même séquence de phases est observée pour les différents systèmes analysés. En effet, la phase hexagonale Ni 5 (GeSn) 3 est obtenue en premier, suivie de la phase Ni(GeSn). Cependant, la nature de l'ion implanté a un impact sur la séquence de phases pendant la réaction à l'état solide. La PAI C retarde considérablement la formation de la phase Ni(GeSn) vers des températures plus élevées. Ce phénomène est lié à l'accumulation de C aux joints de grains qui peut empêcher la diffusion du Ni. À l'autre extrême, la PAI Ge réduit légèrement la température de croissance du Ni(GeSn). On s'attend à ce que les ions Ge s'accumulent aux interfaces. Cette accumulation pourrait alors modifier l'énergie d'interface et permettre la nucléation du Ni(GeSn) à des températures plus basses.

Conclusions générales

Les études menées dans le cadre de cette thèse visaient à développer et caractériser de manière approfondie les intermétalliques Ni / GeSn. L'objectif était de proposer des contacts adaptés pour des dispositifs à base de GeSn. Ainsi, l'évolution de plusieurs systèmes à base de Ni / GeSn au cours de la réaction à l'état solide a été analysée. La séquence de phases se déroule comme suit. Une phase riche en Ni est formée en premier. Cette phase est identifiée comme la phase métastable hexagonale ε-Ni 5 (GeSn) 3 . Lorsque cette dernière a été consommée, la phase orthorhombique Ni(GeSn) est obtenue. Les résultats ont montré que la phase Ni(GeSn) peut être adaptée pour contacter des dispositifs à base de GeSn. Cette phase est obtenue à des températures relativement basses (entre 250-350 ) et les R sh se sont révélées faibles. Cependant, la stabilité thermique de ces types de couches est assez limitée. La dégradation des intermétalliques peut compromettre les propriétés morphologiques et électriques. Le processus de dégradation se produit en raison de deux phénomènes: l'agglomération et la ségrégation de Sn.

Des méthodes alternatives pour améliorer la stabilité thermique du système Ni / GeSn ont ensuite été étudiées. Trois solutions différentes, qui peuvent intervenir à différentes étapes de l'élaboration du dispositif, ont été testées. La technique de pré-amorphisation par implantation (PAI), l'addition de Pt ou de Co à la couche de Ni et l'utilisation du laser comment méthode de recuit. Ces méthodes sont des solutions alternatives pouvant améliorer efficacement la stabilité thermique des couches de Ni / GeSn. Finalement, un aperçu d'une alternative à la métallisation à base de Ni qui consiste a utiliser le Ti a également été présentée. Ce métal a également été considéré comme une alternative appropriée pour contacter des dispositifs basés sur GeSn.

Les résultats de cette thèse de doctorat représentent un bon point de départ pour élaborer des contacts GeSn de haute qualité qui peuvent être intégrés dans des dispositifs électroniques ou opto-électroniques. Une vaste compréhension du système et des propriétés des intermétalliques Ni / GeSn a été acquise. De plus, des alternatives pour modifier les conditions d'élaboration des contacts ont été identifiées et ont été vérifiées comme étant adaptées pour améliorer la stabilité thermique du système Ni / GeSn.