Clément, Bilel, Didier, Mickael, Denis Nathalie Patrick

Elodie, Fabien, David, Julien, Kristell, Maud Lucie

Aide De Karen

Marc Laurent

Del- Phine Doloy

Delphine

François Pierre

Rou- Chon Denis

Denis Jalabert

Yves Morand

Cyrille Le Royer

Sylvie Favier

Nicolas Bernier

Eugénie Martinez

François Andrieu

Jean-Marc Fabbri

Et

pour ces très sympathiques pauses café et l'enthousiasme communicatif dont vous faites preuve ! Merci aussi à

Ce travail de thèse n'aurait pu aboutir sans l'aide et le soutien de nombreuses personnes, à commencer par mes encadrants. Un immense merci à Olivier pour m'avoir accordé sa confiance, pour sa disponibilité (en particulier pour les nombreuses relectures approfondies en cette fin de thèse), et ses encouragements dans les moments les plus délicats. J'ai particulièrement apprécié le subtil équilibre entre autonomie et encadrement que tu as imprimé dès le début de cette thèse. J'ai beaucoup appris grâce à toi sur la conduite d'un travail scientifique, et plus généralement même sur la construction d'un projet quel qu'il soit, ainsi que sur les aspects humains du travail en équipe. Cette expérience et ces compétences me seront très utiles pour la suite de ma carrière, et pour tout cela merci encore ! Je tiens à remercier Elisabeth pour sa grande disponibilité et le soutien dont tu as fait preuve chaque fois que j'en ai eu besoin. Ton soutien tout au long de cette thèse m'a vraiment permis de prendre du recul et de voir plus large que mon travail au sein de l'entreprise, de comprendre les enjeux de la thèse, et d'éviter de nombreux écueils (enfin pas tous sinon ce n'est pas drôle !). Et enfin, merci à François, qui a toujours su apporter un éclairage critique très constructif et des solutions à de nombreuses questions et problématiques rencontrées pendant cette thèse. J'ai trouvé nos discussions, scientifiques et plus générales, très enrichissantes.

Je remercie vivement tous les membres du jury. Pierre Müller et Alain Claverie pour avoir accepté d'être rapporteurs de mon travail de thèse, et merci également à Wilfried Vandervorst d'avoir présidé ce jury.

Je tiens aussi à remercier les nombreuses personnes avec qui j'ai eu la chance de travailler. Et notamment les thésards : Laurent Fauquier (je te souhaite une bonne continuation dans ta nouvelle carrière !), Aurèle Durand, Rémy Berthelon (mention spéciale au succès du LOCOS, une intégration restée … incomprise et non répétable. Un véritable one shot !), et Aurore Durel (anciennement Bonneviale, Félicitations !). Et bon courage à Damien !

Conclusion of Chapter IV Conclusion

Bibliography

Appendices

A -Strain and Stress B -Secondary Ion Mass Spectroscopy (SIMS) C -Wavelength Dispersive X-Ray Fluorescence (WDXRF) D -X-Ray Diffraction (XRD) E -Atomic Force Microscopy (AFM) F -Transmission Electron Microscopy (TEM)

Introduction

It is somewhat intriguing that the frantic computerization of our modern world mostly relies on a minute element often compared to a faucet that is either completely open and carries the binary number "1" or completely shut and carries the binary number "0". This minute element is called a transistor and is merely a tiny electrical switch made of silicon, the most abundant element of the earth crust. Since the first Intel microprocessor in 1971, transistors have moved from gigantic computers to PCs, to laptops, to smartphones and finally to any object that may interact with any object or human being -from a watch or a fridge to traffic regulation systems -, the so-called Internet of Things (IoT). This tremendous spread of electronic devices and networks into our day-to-day life has been enabled by Moore's law, a self-fulfilling prophecy of the semiconductor industry that Gordon Moore stated in 1965. According to this law, the number of transistors crammed into a chip would double every two years.

As long as Moore's law was obeyed, the semiconductor industry followed a two-sided, technological and economic, virtuous circle: the smaller transistors get, the lower their manufacturing costs and power consumptions are and the higher their performances are. The terrific machine that Moore's law is has however been running out of steam since the beginning of the 2000's. Peter Lee, a vice-president of Microsoft Research, humorously said that "There's a law about Moore's law. The number of people predicting the death of Moore's law doubles every two years"*. People now ask what might replace it, and different stories emerge depending on which horizon people are looking at.

One of them is composed from long-term perspectives and opportunities and is about going into the unknown. It starts by telling that no matter how determined we are, the ultimate limit of the "traditional" silicon transistor is the atom, and that it will put a stop at Moore's law. This story has also the merit to highlight that Moore's law is both a blessing and a curse. Moore's law has made progress predictable, but at the same time it has been a hindrance to innovation. One might see in the death of Moore's law the opportunity to shake things up and to explore new paths. Amongst these paths, there are 3D stacking of specialized chips, nanophotonics -which uses photons instead of electrons -, spintronics -which uses the spin of the electron instead of the charge -, quantum computing, graphene and other 2D materials-based computing, and even brain-inspired computing -which strives to take advantage of the 10,000 times higher brain power efficiency over silicon-based chips. Some of them have already proven to outperform today's best silicon transistors. But fabricating a fancy device in a lab is one thing, building billions of them on a single chip, all equals in terms of performances, reliability and lifetime, and at reasonable costs, is another one. This shows that there must be another story, at least for the near-future. This other story, sometimes called "More Moore", has actually begun and is about how much we can push forward Moore's law by introducing "non-conventional performance boosters". To answer this question, one needs to understand what stifles Moore's law. Firstly, downsizing transistors has become increasingly difficult in the past few years and requires huge R&D spending that only a few large companies can afford. Secondly, going to the nanometer scale brings new detrimental effects that have put power consumption and performances on quasi-plateaux for a few years. To overcome these limitations, new channel materials -strained Si in 2003 -, new gate materials -high-κ oxides in 2007 -, and even new transistor architectures -FinFET and FD-SOI in 2012 -, have been introduced. And new "boosters" are still needed to sustain transistor performances and power consumption enhancement until another story begins. This thesis strives to fit into this "More Moore" story by introducing a new material, compressively strained Silicon-Germanium, into the channel of a FD-SOI transistor architecture.

Fully Depleted Silicon-On-Insulator (FD-SOI) is a technology developed in Grenoble ecosystem by Soitec, the Laboratoire d'Electronique et Technologie de l'Information (LETI), and STMicroelectronics as an alternative to the FinFET technology promoted by Intel. The FD-SOI technology benefits from several advantages in terms of power consumption and performances over traditional bulk technologies while having limited manufacturing costs increase thanks to the similarity of most manufacturing processes with traditional ones.

Si hegemony over other materials in the semiconductor industry has been achieved thanks to two major assets: it is the most abundant element of the earth crust, which makes it cheap, and Si dioxide (SiO 2 ) is a stable oxide featuring excellent dielectric properties, which makes Si-SiO 2 an almost perfect channel-gate stack. This hegemony is however faltering. The introduction of high-κ oxides as gate materials to overcome the issue of increasing tunneling current through ultrathin SiO 2 (a few nanometers) -the major if not the only flaw of SiO 2 -has made one of the two assets of Si obsolete. In addition, the quest for higher switching speed and lower power consumption involves increasing the carrier mobility into the channel, with a first hit in 2003 with strained Si. But growing attention has been devoted to other materials, especially Ge and III-V materials, and strain engineering techniques, to go beyond strained Si performances. Compressively strained SiGe is one of the most promising candidates amongst these materials for transistors driven by hole carriers (p-type transistors). As compressively strained SiGe is not adequate for transistors driven by electrons (n-type transistors), and as both n-type and p-type transistors must be integrated on the same chip, co-integration of compressively strained SiGe in p-type transistors with another materials -Si or III-V materials -in n-type transistors is required. SiGe assets in terms of hole mobility can be combined with the benefits of the FD-SOI transistor architecture.

In summary, this thesis seeks to integrate compressively strained SiGe in p-type FD-SOI transistor channels with an industry-compatible process called condensation. The condensation technique has been chosen based on three major advantages over other techniques:

• It allows integration of compressively strained SiGe in p-type channels only while leaving n-type channels unaffected. • Fully strained and high crystal quality SiGe films directly on insulator (SiGe-On-Insulator: SGOI) can be fabricated up to 50% Ge concentration. • Only conventional manufacturing processes are used.

However, several challenges still need to be taken up for an efficient and optimized use of the condensation technique in an industrial environment: controlling kinetics of SiGe thermal oxidation (addressed in Chapter III) and fabricating high crystal quality and highly strained SGOI films by condensation (addressed in Chapter IV). These two challenges are then split into two more fundamental studies (Section III.A and Section IV.A) and two other problematics that have more direct applications to advanced FD-SOI transistors manufacturing (Section III.B and Section IV.B).

The manuscript is organized as follows:

• Chapter I: it describes the context, the theoretical and technological underpinnings of this thesis work, and ends with the definition of the two challenges addressed in Chapter III and Chapter IV.

• Chapter II: Various processes and characterization techniques and a simulation tool were used during this work. Techniques I personnaly used, either for the measurement and/or the analysis, are described in Chapter II while the other ones are presented in Appendices.

• Chapter III: The condensation technique relies on thermal oxidation of SiGe. The first part of this chapter (Section III.A) strives to bring more understanding of the mechanisms and kinetics of SiGe oxidation. Three technology-relevant oxidation processes are investigated. The second part of this chapter (Section III.B) presents two tools built thanks to the former oxidation kinetics results: a numerical simulation tool of SiGe oxidation, and a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, two key variables of the condensation process, as a function of process parameters.

• Chapter IV: High strain levels and high crystal quality of SGOI films are needed for increasing the hole mobility. We thus compare in the first part of this chapter (Section IV.A) strain levels and crystal qualities of SGOI films fabricated by two industry-compatible condensation processes. Finally, the second part of this chapter (Section IV.B) starts by describing a particular SGOI strain relaxation phenomenon occurring during transistor manufacturing after the condensation process. Two possible causes to this strain relaxation phenomenon, and related to the condensation process, are suggested and investigated.

Introduction to Chapter I

This chapter presents the context, the theoretical and technological underpinnings of this thesis work, and ends with the definition of the two challenges addressed in Chapter III and Chapter IV. Today's challenges for further enhancement of Metal Oxide Semiconductor Field Effect Transistors (MOSFETs) performances are firstly introduced. More specifically, with the purpose of boosting the ON current, the different paths for increasing the electron and hole mobility in deeply-scaled MOSFETs are reviewed. Among these paths, the benefits of using compressively strained SiGe-On-Insulator (SGOI) as channel material for the p-type MOSFET are highlighted. The properties of SiGe, and then of SiO 2 , that are relevant to this work are reviewed. A focus is given on strain in SiGe as it is essential to control strain of channel materials. The condensation technique, which allows the fabrication of SGOI films with possibly high Ge concentrations and high strain levels is then introduced. The two fundamental studies (Section III.A and Section IV.A) and the two technology-oriented problematics (Section III.B and Section IV. B) of this thesis are finally stated in order to draw up the organization of this manuscript. This section starts with a brief overview of the history of the MOSFET technology from the first concept in 1926 to today's applications. Then, a description of MOSFET working principles and basic architecture is given. An emphasis is put on MOSFET performance variables that could be enhanced by technological paths related to this thesis work. Having defined MOSFET key elements, the semiconductor industry, Moore's law and key MOSFET technological breakthroughs are introduced in order to draw the general context of this work. Lastly, current scaling issues that impede further conventional scaling are reported.

I.1.1. The Metal Oxide Semiconductor Field Effect Transistor (MOSFET)

In 1926, Lilienfeld filed a US patent describing the field effect transistor but it was not until 1947 that the transistor effect was actually observed and explained by W. Shockley, J. Bardeen and W. H. Brattain at the Bell Labs [Lilienfeld30] (see Figure I.1). The idea of Lilienfeld was to control an electric current flowing between two terminals by controlling a third potential applied between these two terminals.

The actual MOSFET technology breakthrough happened in 1960 thanks to D. Kahng and M. M. Atalla of the Bell Labs [Kahng63]. The MOSFETs were fabricated using crystalline silicon (Si) as the substrate and thermal silicon dioxide (SiO 2 ) as the gate dielectric. While germanium (Ge) was used for the first transistor by W. Shockley's team, it was not used for MOSFETs due to the non-existence of a stable germanium oxide that features good electrical performances. Complementary Metal Oxide Semiconductor (CMOS) has been the building block of Integrated Circuits (ICs) since it was patented in 1963 by F. Wanlass [Wanlass67]. A CMOS cell uses two MOSFETs working in tandem and driven by electron current and hole current for the so-called nMOSFET and pMOSFET respectively. Then, the first Intel microprocessor in 1971, along with the development of clean rooms and key process steps such as photolithography, were the beginnings of the history of the semiconductor industry.

While the building blocks of CMOS logic have remained unchanged, increasing needs for differentiated ICs have led to two main applications and their associated MOSFET design, i.e. high performance (HP) transistors, mainly used in computers and servers, and low power (LP) transistors for power sensitive markets such as mobile consumer multimedia or the growing "Internet of Things" market.

I.1.1.2. MOSFET Working Principles

A MOSFET is based on the unique properties of semiconductor materials, i.e. electrical conductivities between those of a conductor and an insulator. A semiconductor material may be either n-type when majority charge carriers are electrons, or p-type when majority charge carriers are holes. A "hole" is a pseudo-particle whose charge is the inverse of the charge of an electron, and which actually corresponds to the absence of an electron in the valence band. The contribution of holes to the current is equivalent to the contribution of all the electrons of the valence band, but study of the physics of semiconductor materials is much easier when considering holes rather than all the electrons of the valence band.

The electrical conductivity of a semiconductor material is controlled by introducing impurities called "dopants" which also determine the type of the semiconductor material. Silicon, a IV-column material in the Mendeleev periodic table, is the most used semiconductor material. Currently used dopants for silicon are boron (a III-column material) for p-type silicon and phosphorous and arsenic (both V-column materials) for n-type silicon.

The concept of the Field Effect Transistor (FET) is to control the conductivity of a semiconductor channel sandwiched between two ohmic contacts (the source and the drain) by applying a voltage to a third contact (the gate). In a MOSFET, the gate is separated from the channel by an insulating layer and the voltage applied to the gate controls the conductivity of the channel just below the insulator layer and thus impedes or allows the current to flow through the channel. Figure I.2 shows schematic representations of n-type and p-type silicon-based MOSFETs. pMOSFET and nMOSFET working principles are similar and from now on, we focus only on pMOSFET. In terms of device architecture, pMOSFETs are composed of several regions: ȃ a n-type substrate ȃ heavily p doped (p+) wells that are connected to the source and the drain: they provide holes for the current flowing through the channel ȃ the gate: a metallic material or highly-doped polycrystalline silicon ȃ the dielectric layer between the gate and the channel ȃ the channel: the n-doped silicon region under the gate Three operating modes are defined depending on the applied gate voltage:

No voltage is applied to the gate: the source-channeldrain is composed of two opposite pn junctions, which prevent the current to flow through the channel from the source to the drain. The MOSFET is in off state and the current between the source and the drain is defined as I I DS off

=

. This current arises from leakage current in pn junctions.

Off state

When a weak negative voltage is applied to the gate, majority carriers within the substrate, i.e. electrons, are expelled from the channel region: a depletion region is created. This regime is called weak inversion regime. In this regime, the current density is driven by diffusion in junctions and is exponentially dependent on the gate voltage [Sze06]:

I C W L k T q qV k T q V V k DS dep B D B G T B =       - -             - µ 0 2 1 exp exp T T       ,
where: ȃ μ 0 is the intrinsic carrier mobility in the channel ȃ C dep is the depletion zone capacitance ȃ q is the elemental charge, k B is the Boltzamm constant and T the temperature ȃ W is the gate width

The subthreshold slope S is defined as: When the applied gate voltage, V G , is decreased, holes are attracted under the gate at the dielectric-channel interface. The concentration of minority carriers (holes) is higher than the concentration of majority carriers (electrons) in this region, hence the strong "inversion" regime. The threshold voltage, V t , is defined as the applied gate voltage that corres- ponds to the switch between weak and strong inversion regimes. The depletion zone reaches a maximum expansion at the threshold voltage.

S V I k T q C C G DS B dep Ox eff = ∂ ∂ = +         log( ) ln , , 10

On state: strong inversion regime

Depletion region Inversion zone

There are two regimes depending on the drain voltage:

• The linear regime: The channel behaves like a resistance whose resistivity depends on the minority carriers concentration in the channel. The drain current is:

I W L C V V V V DS eff Ox eff G t D D = --       µ , , 1 2
where μ eff is the effective mobility in the channel. For small drain voltage, the drain current is linearly dependent on the gate voltage.

Equation I.3

For drain voltages lower than V Dsat , the inversion layer length is reduced. and where V DD , f , C load , C G , n ON , n OFF are the applied gate voltage in the ON regime, also called the power supply voltage, the operating frequency, the total load capacitance, the gate capacitance, the number of transistor in ON state and the number of transistors in OFF state.

Achieving simultaneously power consumption and switching time reduction leads to MOSFET design requirements that may conflict with each other. Power consumption reduction is preferentially achieved by reducing P Standby , in other words V DD and I OFF . But reducing V DD induces I ON reduction and consequently switching time degradation because I V ON DD ∝ 2 . We highlight that enhancement of the effective mobility μ eff in the channel allows I ON enhancement without conflicting power consumption requirements.

Silicon-based transistors have been the foundation of the semiconductor industry since 1959, year of the first silicon-based Integrated Circuit (IC). Since then, the semiconductor industry has followed the virtuous circle schematically represented Figure I.9 and whose key enablers have been MOSFET scaling and wafer size increasing. MOSFET scaling refers to the reduction of the transistor dimensions usually defined by the half pitch between two adjacent interconnections, the so-called technology node. Moore's law, an empirical rule stated in 1965 and recalibrated in 1975 by G. Moore has well described the driving force of the semiconductor industry [Moore65]. Moore's law predicted that the number of transistors per silicon area would double every two years. Another way of stating this rule is saying that chip performances, a combination of higher transistor density and higher transistor speed, double every two years. In terms of device dimensions, Moore's law is achieved by scaling the contacted gate pitch by a scaling factor of 0.7 every two years (see Figure I.11). Until 2003, Moore's law has been synonymous with "classic" or "conventional" scaling which refers to the scaling rules described by Dennard et al. [Dennard74]. According to Dennard scaling rules, the oxide thickness, the transistor length and the transistor width (perpendicular to the channel direction) were scaled by a constant factor 1/k which in turn enhanced speed by a factor k at constant power.

But since 2003, conventional MOSFET scaling has not been sufficient to follow Moore's law because of the scaling issues described in the following section. Therefore, "non-conventional performance boosters" were commercially introduced: strained channel in 2003, high-κ metal gate stack in 2007 and new transitor architectures, FinFET and FD-SOI in 2012. These scaling issues are expected to increase further as devices are scaled down. Henceforth, Moore's law and its derivative term "More Moore" refer to slight scaling and technology advances which allow either cheaper, smaller or faster chips and not anymore necessarily all of them simultaneously.

I.1.1.5. MOSFET Scaling Issues

When the gate length is scaled down to a few tens of nanometers, the drain-junction depletion region extends into the channel region, causing a reduction in the effective gate length as schematically represented Figure I.12. Short-channel effects (SCE) result from the sharing of electrical charges between the channel on one hand and the drain and source regions on the other hand. Therefore, the distance between the source and the drain, and the applied drain voltage influence the gate control of the transistor. SCE induce two detrimental effects: drain-induced barrier lowering (DIBL) which causes the threshold voltage to decrease when the drain voltage increases, and an increase in the subthreshold slope. Both increase the leakage current and DIBL also decreases the switching speed of the transistor. Additionally, decreasing the gate dielectric thickness induces a higher tunneling current through the gate that also increases leakage currents. As a consequence, the semiconductor industry is at a point where conventional transistors cannot follow Moore's law anymore. Technological innovations are thus required to overcome these scaling issues and to sustain transistor performances enhancement. Main technological innovations, such as new MOSFET architectures like FD-SOI and FinFET, and high mobility channels are described in the following sections.

FinFET and FD-SOI are two transistor architectures whose performance advantages over conventional MOS-FETs, also called bulk MOSFETs, are first introduced along with the target applications. The second part of this section focuses on FD-SOI technology nodes that are the target applications of this work. Even if being in advanced R&D stages only, and except the FinFET architecture which has already been commercialized, multigate transistor architectures deserve to be mentionned as they may some day take over FD-SOI and FinFET architectures in most advanced CMOS devices. Multigate transistors actually refer to single gate transistors whose gate is wrapped around several sides of the channel region. FinFET, tri-gate, gate-all-around, Π-gate and Ω-gate structures are the main multigate transistor architectures [Ferain11].

FinFET Fully Depleted Transistors

Schematic profile views of a conventional bulk transistor and a FD-SOI transistor are represented Figure I.14. The FD-SOI technology features reduced parasitic currents and reduced junction capacitances between the source and the drain over bulk MOSFET technology. Other advantages of FD-SOI over bulk are better device-to-device isolation, reduced substrate coupling and lower sensibility to high-energy radiations that may cause computing errors. [START_REF] O P Paine | Oxidation of Si 1-x Ge x alloys at atmospheric and elevated pressure[END_REF].

FinFET and FD-SOI technologies do not address exactly the same markets (see Figure I.15). Due to expected higher speed, density, and costs, the 3D-architecture FinFET mainly addresses high-performance markets. While FD-SOI, being manufactured by more conventional processes, benefits from lower R&D and manufacturing costs thanks to a larger accumulated knowledge. However FD-SOI requires Silicon-On-Insulator (SOI) substrates which are more expensive than bulk ones. In the end, FD-SOI better fits the needs of power sensitive markets. From now on, we focus on FD-SOI, the transistor architecture which is the target application of this work. The FD-SOI technology addresses power sensitive markets while FinFET is more suitable for high-performance markets. From [START_REF] O P Paine | Oxidation of Si 1-x Ge x alloys at atmospheric and elevated pressure[END_REF].

FD-SOI is a technology that has been developed by the LETI, Soitec, and STMicroelectronics since the beginning of the 1990's (Figure I.16). SOI substrates specifications is a 20 nm Buried Oxide (BOx) thickness and a 10 nm SOI thickness and the final channel thickness is 6.5 nm. A Transmission Electron Microscopy (TEM) image of a FD-SOI pMOS device fabricated at Crolles site and featuring a SiGe channel, raised boron-doped SiGe source and drain, and high-κ metal gate is shown Figure I.17. This SiGe channel is a MOSFET performance booster that is described in the next section. Today's FD-SOI specifications are in the 20% to 30% range for the Ge concentration within the SiGe channel.

I.1.2.2. FD-SOI Technology Nodes

Figure I.16: R&D history of FD-SOI. From [START_REF] O P Paine | Oxidation of Si 1-x Ge x alloys at atmospheric and elevated pressure[END_REF].

I.1.3. High Mobility Channels

As seen in Section I. 1.1.3., an effective way to increase the ON current is to use channels that feature high carrier mobility. We thus start by defining what mobility is and we provide more details about the relations between mobility and current in real transistors. This section then examines electron and hole mobility of various materials that could be suitable for high-mobility channels on Si substrates. Mobility dependence on crystal orientation and strain is finally reviewed.

The mobility of a carrier in a given material describes its ability to move under an applied electric field. Mobility µ is defined for a specific crystal orientation as a function of two important materials properties, i.e. the carrier effective mass m * , and the relaxation time τ :

µ τ = * q m .
The effective mass of a carrier is used to quantify how the carrier responds to the application of external forces in a solid. Indeed, the periodic potential due to the periodic position of ions and core electrons within a crystal induces internal forces. Therefore, the behaviour of a charge carrier under an electric field in a solid is not the same as in vacuum. These internal forces may be described by the dispersion relation E k ( ), where E is the energy of the carrier, k its wave vector, and  the Planck constant. The effective mass is defined by:

m E k * - = ∂ ∂        2 2 2 1 .
The relaxation time is the average carrier transit time between two collisions and depends on various scattering mechanisms caused by any change in the perfect crystal structure. In FD-SOI MOSFETs, the channel is fully depleted so that we have to consider as scattering mechanisms the interactions of the carriers not only with the gate dielectric layer but also with the BOx layer. The scattering mechanisms are:

ȃ Broken crystal symettry induced by confinement in deeply-scaled FD-SOI MOSFET channels. ȃ Thermal vibration-induced acoustic and optical phonon scattering, also referred to as Remote Phonon Scattering (RPS), is an intrinsic mobility degradation mechanism due to coupling of phonons in dielectric layers with carriers in the channel. RPS could thus represent a fundamental limit to the channel mobility of FD-SOI MOSFETs. ȃ Remote Coulomb Scattering (RCS) is an extrinsic mechanism of scattering of carriers by trapped charges in dielectric layers. This effect may be reduced by screening them using a layer between the channel and the dielectric layer where the charges are trapped. This layer keeps away the trapped charges from the carriers. ȃ Interface roughness.

I.1.3.1. Mobility Definition

Equation I.7 Equation I.8

The relaxation time is dominated by the scattering mechanism with the shortest scattering time (Mathiessen's rule) and we usually consider Coulomb, phonon and roughness scattering mechanisms only according to the "universal mobility model" [Takagi94]. These three scattering mechanisms have different electric field and temperature dependences as shown Figure I.18. At low fields, the mobility is limited by Coulomb scattering , at medium fields by phonon scattering and finally by interface roughness at high electric fields.

I.1.3.2. Mobility Engineering for ON Current Enhancement

In Section I.1.1.2., we gave the relation between the ON current and the effective mobility of the channel material. This relation do not hold anymore for short channel transistors and under high electric fields, which are today's specifications. To discuss the impact of mobility on the ON current in real devices, we provide below a simple equation of the ON current of a MOSFET per gate width [Takagi08]:

I qN v ON S Source s  ,
where N S Source is the surface carrier concentration near the source edge and is determined by the gate voltage and the gate capacitance, and where v s is the carrier velocity near the source edge. On the other hand, the carrier ve- locity depends on the channel length. Three regimes of the velocity in MOSFETs are distinguished [Lundstrom02]:

• The stationary transport regime in long-channel MOSFETs. In this case and when velocity saturation occurs, the ON current dependence on the low field mobility becomes weaker than the linear one described by Equation I.4.

• The quasi-ballistic transport regime. In this case, nonstationary transport becomes dominant in short channels because the number of scattering events is reduced in the channel. The ON current is then redefined as:

I qN v r r ON S Source = - + inj , 1 1
where v inj and r are respectively the injection velocity near the source edge and the backscattering rate near the source edge. r is related to the mobility, therefore mobility enhancement is still an important path for MOSFET performance improvement.

• Full ballistic regime of very short-channel MOSFETs. There is no scattering event in the channel and no direct mobility dependency anymore. The ON current is:

I qN v ON S Source = inj .
In this case, MOSFET performance enhancement is achieved by improving the injection velocity which follows the following relation: v m inj / ∝ -1 2 . As shown by Equation I.7, decreasing the effective mass in- duces a mobility enhancement and therefore, mobility improvement is still an ON current enhancement path in full ballistic transport regime.

Equation I.9 Equation I.11 Equation I.10

Research has shown promising results of integration of III-V materials and germanium on silicon substrate in spite of the lattice parameter discrepancy existing between these materials [Kuhn12]. Fabrication of suitable gate stacks on these alternative materials has also been proven to be possible despite several remaining issues [START_REF] Fauquier | Benefits of XPS nanocharacterization for process development and industrial control of thin SiGe channel layers in advanced CMOS technologies[END_REF].

The CMOS technology is based on co-integration of pMOS and nMOS transistors working in tandem and featuring similar performances. Therefore, while III-V materials show electron mobilities higher than 10,000 cm 2 V - 1 s -1 , they are not the most promising materials for pMOS transistors. On the other hand, germanium has the highest hole mobility (1,900 cm 2 V -1 s -1 ) of all known semiconductor materials and is thus the most promising candidate for pMOS transistors. Germanium also features a higher theoretical electron mobility than silicon, but it has never been experimentally observed and is thus not suitable for nMOS transistors. As for SiGe, the theoretical hole mobility of unstrained SiGe is dependent on calculations parameters and must be considered with caution [Fischetti96]. Co-integrating different channel materials for pMOS and nMOS on the same chip may offer the highest performance enhancement for CMOS technology but increases process and integration complexity at the same time.

We mention that decreasing the effective mass of carriers within the channel has also drawbacks because of higher quantum confinement effects. Indeed, quantum confinement effects degrade the density of states which in turn decreases the current density. Another detrimental effect is the higher leakage current induced by the increase of the tunneling current from source to drain. Trade-offs have to be found based on target applications. STMicroelectronics chosen strategy for its FD-SOI technology is co-integration of tensely strained Si for nMOS channel and compressively strained SiGe for pMOS channel, the latter being this work's target application.

Crystal orientation of both the substrate and the channel alters band structures, and consequently mobility within the channel, because of the anisotropy of the effective mass of some materials such as Si and Ge. Conventional orientations for Si CMOS are <110> channels on (001) substrates. CMOS fabrication on ( 110) and ( 111) substrate orientations has been hindered because of poor gate oxide quality. Nevertheless, MOSFET performance enhancement has been recently shown on these orientations thanks to careful optimization of gate stack formation From [START_REF] Prabhakaran | In situ oxidation of a thin layer of Ge on Si (001): Observation of GeO to SiO 2 transition[END_REF]. I.1.3.4. Crystal Orientations Figure I.20: Schematic representations of conventional orientation (left), (001)<100> orientation (middle) and (110)<110> orientations (right) along with mobility gains.

Strain refers to the lattice deformation of a material and is either compressive or tensile. First theoretical investigation of strain-induced mobility enhancement can be traced back to the deformation potential theory formulated by J. Bardeen and W. B. Shockley in 1950 [Bardeen50] and the experimental observation of the piezoresistance effect by C. S. Smith [Smith54]. Since then, this phenomenon has been theoretically investigated by many groups, the latest extensive study based on the deformation potential theory being Fischetti and Laux's one [Fischetti96].

Strain, and subsequently stress, reduces the crystal symmetry and therefore influences band structures in two ways: ȃ Band warping (except for hydrostatic, in other words isotropic strain) induces a change in the effective mass. ȃ Band splitting changes the relative importance of intra and inter-valley scattering and therefore influences the relaxation time of charge carriers. On the experimental and technological side, strain within the channel has been implemented using various device and process technologies. First devices thought to benefit from strain-induced enhancement of electron and hole mobilities within the channel were Si/SiGe heterostructures [Manasevit82]. Lattice mismatch between Si and SiGe was suspected to induce strain which in turn increased mobility. Then, the Stanford group confirmed such a performance improvement in 1992 [Welser92] and Intel implemented strain in commercial devices in 2003 for the . For the pMOSFET, the larger lattice parameter of SiGe in the source and drain regions tends to compressively strains the smaller lattice of Si within the channel, while the nMOSFET Si channel is tensely strained by engineering the stress and thickness of a silicon nitride capping layer (see As drawn in Section I.1.1., conventional MOSFET scaling has come to an end and new technology innovations are needed to sustain the dynamics of the semiconductor industry. Firstly, new transistor architectures, i.e. multigate transistors such as FinFET and FD-SOI, can help to extend transistor scaling. STMicroelectronics has chosen to focus on the FD-SOI transistor architecture to sustain CMOS scaling, which is why FD-SOI is the target application of this work. Then, we explained in Section I.1.3. why higher mobility channels are also required to enhance transistor performances. Crystal orientation, intrinsic materials mobility and strain are three paths to increase channel mobility. Compressively strained SiGe is considered to be one of the best candidates for the pMOS channel of FD-SOI transistors.

To summarize, the target application of this thesis work, schematically represented Figure I.23, is fabrication of high Ge-content (higher than 40%) compressively strained SiGe channel for enhancing the performances of pMOS FD-SOI transistors. Different strain elements, strain levels and strain orientations can be used for mobility enhancement. This work focuses on hole mobility enhancement induced by biaxial or uniaxial strain in the SiGe-based channel. For SiGe, a dramatic increase of electron and hole mobility is expected from calculations and some experimental result have shown high factor improvements [DeSalvo14].

I.1.4. Target Application of this Thesis Work

Compressively strained SiGe channel

As drawn in the last section, co-integration of SOI and SiGe-On-Insulator (SGOI) channels in nMOSFETs and pMOSFETs for the next generations of CMOS devices is strongly needed. Several techniques to fabricate SGOI layers have been proposed (see [START_REF] Prabhakaran | In situ oxidation of a thin layer of Ge on Si (001): Observation of GeO to SiO 2 transition[END_REF] and [Takagi15] and references therein), but many of them are either not mature, costly, or incompatible with today's CMOS manufacturing processes. The condensation technique, which will be described into more details in Section I.4., shows three main advantages over the other techniques:

ȃ It allows fabrication of SGOI layers in pMOS areas only, leaving the nMOS areas unaffected. ȃ Fully strained SGOI layers a few nanometers thick with Ge concentrations up to about 50% can be fabricated. ȃ It uses conventional SOI substrates and conventional CMOS manufacturing processes, such as SiGe epitaxial growth and thermal oxidation.

However, incorporating Ge in CMOS manufacturing processes with the condensation technique is not straightforward as new issues emerge. Firstly, kinetics of SiGe thermal oxidation, upon which the condensation technique is based, must be well controlled. And secondly, the 4% lattice mismatch between Si and Ge makes SiGe layers prone to strain relaxation and extended defects generation, which are detrimental to pMOS performances. Other issues, out of the scope of this work, such as fabrication of high-quality gate stack and source and drain with low contact resistance, etch and cleaning processes, also need to be solved.

I.2. Properties of SiGe

In this section, essential material properties of SiGe are presented. Crystal structure of SiGe, phase diagrams, from which limitations in terms of annealing temperature are extracted, and band structures and bandgap of SiGe are reviewed.

Si and Ge are both elements of the IV column of the Mendeleev chart and their atomic numbers are 14 and 32 respectively. They both have four covalent bonds. SiGe with Ge content varying from 0% (pure Si) to 100% (pure Ge) crystallize in a diamond-like lattice structure whose space group is O Fd m The lattice constants of Si and Ge are 5.431 Å and 5.658 Å respectively. Therefore, there is a 4.18% lattice mismatch between Si and Ge. Si 1-x Ge x lattice parameter may be described by the empirical parabolic law proposed by Dismukes et al. [Dismukes64]:

a x nm x x SiGe ( )( ) . . . . = + + 0 5431 0 20 0 027 2

Equation I.12

The phase diagram of SiGe is shown Figure I.25. The temperature of the solidus decreases from 1412°C, the Si melting temperature, to 937°C, the Ge melting temperature, as the Ge content increases. Thermal annealing with temperatures higher than the solidus would lead to the separation of two phases, a liquid Ge-rich one and a solid Si-rich one. Consequently, a first limitation in terms of process is that the highest temperature which can be used for thermal annealing of SiGe decreases as the Ge content increases.

I.2.2. Strain and Stress

Strain is defined as the relative lattice displacement created by deformation of the crystal structure of the material. Strain is fully characterized by the following second-order tensor: Stress is defined as the force responding to strain in a unit area and is fully defined by the second-order stress tensor:

ε ε ε ε ε ε ε ε ε ε =          
σ σ σ σ σ σ σ σ σ σ =          

I.2.2.1. Definition of Strain and Stress

The section aims to give some key elements about stress and strain. A more general description of stress and strain laws and formalism is given in Appendice A. After a brief introduction to the general laws of stress and strain, we give relations about stress and strain in SiO 2 and in SiGe that are specifically used in this manuscript. This section ends with an introduction to the formalism of the mathematical treatment of the influence of stress on carrier mobility. Within the linear elastic theory, stress and strain are related by Hooke's law:

Equation

σ ε = C. or ε σ = S. ,
where C and S are the elastic moduli and compliances fourth-order tensors. C and S are inverse to each other.

For a homogeneous and isotropic material such as SiO 2 , strain and stress are related to each other by Young's modulus E and Poisson ratio ν:

ε σ ν σ σ xx xx yy zz E = - + ( ) ( ) 1 ε σ ν σ σ yy yy zz xx E = - + ( ) ( ) 1 ε σ ν σ σ zz zz xx yy E = - + ( ) ( )

Ge concentration (%)

The in-plane strain ε  and out-of-plane strain ε ⊥ in the SiGe film are:

ε   = - a a a rSiGe rSiGe ε ⊥ ⊥ = - a a a rSiGe rSiGe .
And where:

ε ν ν ε ⊥ = - 2 1  .
And:

ν ν = = - [ ] 001 0 278 0 005 . . , x
where x is the Ge fraction and a rSiGe the lattice parameter of relaxed SiGe.

Strain is sometimes confused with the in-plane deformation of the SiGe lattice parameter with respect to the Si lattice parameter. Indeed, strain of a SiGe layer is relative to the lattice parameter of relaxed SiGe and not to the Si lattice parameter. Deformations are the quantities usually measured by various strain characterization techniques such as Precession Electron Diffration (PED) and X-Ray Diffraction (XRD). Therefore, from now on, strain will always refer to the definitions given by Equations I.17 while the in-plane deformation d  and out-of-plane deformation d ⊥ are defined as:

d a a a d a a a Si Si Si Si   = - = - ⊥ ⊥ .
When a SiGe layer is biaxially strained to a (001), the in-plane deformation is equal to zero while the out-ofplane deformation for a given Ge concentration can be calculated using Dismukes law and Poisson theory:

a a x C C a a rSiGe rSiGe ⊥ = ( ) + -               1 2 1 12 11  . Figure I
.31 shows the out-of-plane deformation of a fully strained SiGe layer as a function of the Ge concentration.

Equation I.19

Equations I.17

Equation I.18 Equations I.20 Equation I.21

Mobility dependency on stress and therefore on strain is usually described by the theory of the piezoresistive effect. Piezoresistivity is an effect of stress-induced resistivity change of a material that has two contributions [Sun09]: the geometrical change of the sample, and the resistivity change induced by strain-altered band structures (and subsequently the mobility). The latter is the dominant one in semiconductor materials. Piezoresistivity and stress are related by the following equation:

- ∆ = = ∑ µ µ π τ i ik k k 1 6 ,
where π ik are the piezoresistance coefficients. The relevant stress conditions for this work are illustrated

I.2.3. Epitaxy of SiGe on Si

The lattice mismatch between SiGe and Si induces strain in both layers, which in turn alters thermodynamic properties of the growing film and therefore influences the growth morphology. We distinguish three main film morphologies:

• Pseudomorphic film:

The film is fully strained and its in-plane lattice parameter is equal to the one of Si. The out-of-plane lattice parameter is larger than the one of Si as described by Poisson's theory (see Section I.2.2.3.).

• Elastic relaxation:

As the film thickness increases, the pseudomorphic film becomes thermodynamically unstable and seeks a way to relax strain. Surface ondulations allow the accumulated energy to decrease. • Plastic relaxation: When elastic relaxation is kinetically suppressed or becomes innefective, the strain relief proceeds by plastic relaxation, which refers to the formation of extended defects such as misfit dislocations (see Figure I.37). Introduction of misfit dislocations is the most efficient way to relieve strain, but because of the high energy cost associated with the dislocation formation, introduction of dislocations becomes energetically favorable only after a certain critical film thickness. Moreover, since the dislocation formation is a thermally activated process, the onset of plastic relaxation may be delayed to higher thicknesses due to kinetic constraints. In this case, the film is said to be metastable (see Section I.2.4.2.).

I.2.4. Plastic Strain Relaxation in Compressively Strained SiGe Films

Crystals defects are defects in the three-dimensional regular arrangement of atoms. They are classified according to their extension in space, i.e., zero, one, two and three dimensions: ȃ Zero-dimensional defects, also called point defects, include vacancies, which are missing atoms in the lattice, and interstitials, which are atoms incorporated in interstitial sites of the crystal lattice. ȃ One-dimensional defects, the so-called dislocations, occur when the regular arrangement of atoms is disturbed along a line. A schematic view of a dislocation is shown Figure I.38. ȃ Two-dimensional defects are grain and phase boundary. ȃ Three-dimensional defects are precipitates and inclusions.

Figure I.38: Schematic view of an edge dislocation in a cubic crystal. A single vertical plane of atoms is missing and the dislocation line is located at the edge of the half plane.

After a brief introduction to the context and definition of this section, we give an overview of the two theoretical approaches to plastic strain relaxation: equilibrum and kinetic models. We then describe the mechanism of dislocation nucleation and associated specific crystal defects in the lattice structure of SiGe.

I.2.4.1. Introduction

This work focuses on dislocations induced by plastic relaxation in compressively strained SGOI layers formed by the condensation technique. Plastic relaxation decreases elastic strain in the layer and consequently lowers the hole mobility, and it also introduces defects that further reduce the hole mobility and may even kill transistor proper functionning in deeply-scaled transistors. Therefore, it is crucial to study the dependence of plastic relaxation of SGOI layers formed by the condensation technique on sample properties and condensation process conditions. Most of the theoretical and experimental studies are about plastic relaxation in heteroepitaxial SiGe layers on bulk Silicon. Our system, a SiGe layer sandwiched between a buried oxide layer and a thermal oxide, differs from heteroepitaxial systems because of the nature of both bottom and upper interfaces and also because of the process used to fabricate the layer.

There are two main theoretical approaches to plastic relaxation in strained layers:

I.2.4.2. The Equilibrium and Kinetic Models

ȃ Equilibrium models based on energetic considerations only were firstly proposed . A SiGe layer pseudomorphically strained with respect to Si has a certain amount of elastic strain energy that increases with the thickness of the SiGe layer. Above a critical amount of elastic strain energy, corresponding to a certain thickness called the critical thickness, introduction of plastic strain energy instead of elastic strain energy decreases the total energy of the layer and is therefore favorable. The most accepted expression of the critical thickness is the Matthews-Blakeslee one [Matthews74] (see Figure I.39). However, large discrepancies have been observed between theoretical critical thicknesses and experimental results [Bean84][Bai94]. The consensus today seems to be that the critical thickness cannot be based on energetic considerations only, and that kinetic barriers that may delay plastic strain relaxation have to be considered (the kinetic models described below).

Today's susbstrates feature very low densities of dislocations. Thus, it has been shown that technologically relevant plastic strain relaxation mechanisms are nuclation and glide of dislocations, and that dislocation multiplication from dislocations already present in the substrate and pinning can be neglected [Menéndez09]. The mechanism of dislocation nucleation, i.e. dislocation half-loop nucleation at the upper interface and its gliding and expansion within {111} planes of the SiGe layer, is illustrated Figure I.40. Inhomogeneities, such as impurities, local roughness, local higher Ge concentration due to the fluctuation of the distribution of Si and Ge atoms in the lattice, and point defects, increase the local stress and lower the activation energy of atomic displacements. Thus, sinks of dislocation half-loops are located at interfaces. In compressively strained SiGe layers, dislocation half-loops expand and glide within {111} planes because they are the denser ones and are thus favorable to atomic displacements. In order to attain equilibrium, the half-loops reach the bottom interface where they are split into three segments: two threading dislocations that extend from the bottom interface to the upper surface within {111} planes, and the misfit dislocation at the bottom interface. Dissociation of these dislocations into partial dislocations can occur as described in the following section. . Kinetic barriers to plastic strain relaxation are indeed particularly large in heteroepitaxial SiGe layers and may allow epitaxial growth of fully strained SiGe layers with thicknesses more than a decade above the equilibrium critical thickness (such as the one proposed by Matthews and Blakeslee). In this case, the system is said to be metastable and would return to the strain level defined by equilibrium models after thermal annealing. However, while equilibrium models cannot reproduce accurately most of actual epitaxial growth results, kinetic models suffer from practical limitations. Indeed, comparison with experiments and predictions are far more complexed with kinetic models because of the large number of unknown parameters. Moreover, except from the one published by Menéndez et al., they require numerical computation, making fitting to experimental data complicated.

I.2.4.3. Mechanisms of Plastic Strain Relaxation in Compressively Strained SiGe Layers

Every dislocation can be described by its dislocation vector  ξ and its burgers vector  b . The burgers vector describes the shift of one part of the crystal with respect to the other one. It is quantitatively defined by the socalled burgers circuit illustrated Figure I.41. The burgers circuit is the additional step necessary to close the circuit in the reference lattice structure relative to the same lattice structure but with a dislocation.

The above sections considered an elastic continuum while crystals are made of particular three-dimensional arrangements of atoms. Therefore, burgers vectors are usually lattice vectors. A dislocation has an intrinsic energy because of the elastic stress field around it. It was shown that the dislocation energy is roughly proportional to the square of its burgers vector. Thus, according to the Frank criterion, a dislocation whose burgers vector is  b can be split into two dislocations of burgers vectors

 b 1 and  b 2 if: b b b 2 1 2 2 2 〉 + .
As a consequence, in SiGe layers whose growth direction is <100>, the dominating slip system is {111} and the stable burgers vector is a/2<100>, which is the minimum translation in <100> orientations. This dislocation is called a perfect 60° dislocation and has an angle of 60° between its dislocation line and its burgers vector .

As illustrated Figure I.42 (a), the face-centered cubic (fcc) lattice is based on a ...ABCABCA... stacking sequence. If the correct sequence is interrupted, a stacking fault is formed. According to the Frank criterion, a 60° perfect dislocation can be dissociated into two partial dislocations whose burgers vectors are a/6<211> and a/2<1-12>, and whose angles are 90° and 30° respectively. The partial dislocations, called Shockley partial dislocations, are stacking faults gliding on {111} planes. One of the partial dislocation forms a low-energy secondnearest-neighbor stacking fault whose stacking sequence is ...ABC|BCA.... And the other one produces a high-energy nearest-neighbor stacking fault whose stacking sequence is ...ABC|CAB.... The distance between these two partial dislocations is obtained by balancing their formation energy and their repulsive energy.

Another way to produce a planar fault in the fcc lattice is to mirror the stacking sequence on one side of a {111} plane. Such a fault, called a twin, is illustrated Figure I.42 (b) and has the following stacking sequence type: ...ABC|BCACB....

A B C A B C (a) (b) ½[-110] (111)
Interdiffusion is a key mechanism of oxidation of SiGe. This section starts by giving an insight about the driving forces of diffusion as well as atomic-scale diffusion mechanisms. Then, we focus on interdiffusion of SiGe since it is the specific diffusion regime during thermal oxidation of SiGe.

I.2.5. SiGe Interdiffusion

I.2.5.1. Introduction and Continuum Theory of Diffusion

Diffusion in solids is the transport of atoms or molecules from one point to another one. Diffusion is a complex and wide subject, and this section does not intend to give an in-depth description of it. It rather encompasses mechanisms and properties that are relevant to this thesis work. In this section, we only consider diffusion in monocrystalline SiGe that do not have extended defects, such as dislocations, that can act as diffusion short-circuits. We also do not take into account other driving forces such as electrical field or temperature gradient.

The diffusion flux is proportional to the chemical gradient and the chemical gradient is proportional to the composition gradient in dilute and ideal solid solutions [Mehrer07]. Diffusion tends to equalize concentrations of elements in the solid. The most common approach to diffusion is Fick's laws. The First Fick's law is:

J = -∇ D C,
where J, D and C are the flux of particles, the diffusion coefficient or diffusivity, and the particles concentration. Diffusion is driven by Brownian motion, and is therefore a thermally activated phenomenon. Diffusion depends on thermodynamic variables such as temperature, pressure and composition, as well as on strain. Diffusivity is usually found to follow an Arrhenius law:

D D H k T B = -       0 exp , ∆
where k B is the Boltzmann constant, T is the temperature, D 0 is the pre-exponential factor, and ΔH is the activation enthalpy of the considered diffusing species. The physical interpretation of the activation energy and of the frequency factor is dependent on the considered diffusion mechanism and process, and is beyond the scope of this thesis work. Let us note that if several diffusion mechanisms are involved, for instance interstitial and vacancy mechanisms, the total diffusivity is the sum of the diffusivities associated with each diffusion mechanism.

Fick's second law is then retrieved by combining Fick's first law with the equation of continuity, which is another spelling of the law of conservation of matter. Fick's second law is:

∂ ∂ = ∇ ∇ C t D C .(
).

Most diffusion problems consists in solving Fick's second law, which is a second-order differential equation. Fick's laws are a purely continuum and phenomenological description of diffusion that do not derive from atomic-scale motion of particles. Comprehension of atomic-scale and discrete motion processes is also required for a proper use of Fick's laws, and is the object of the next section.

Interdiffusion in SiGe

We have to consider atomic diffusion mechanisms because they strongly impact diffusivity. There are two main diffusion mechanisms in monocrystalline materials: non-defect mechanisms which refer to exchanges between lattice atoms and which are not believed to play any significant role in diffusion in SiGe, and native point defect (interstitials and vacancies) mechanisms [Kube10]. Interstitials are atoms incorporated in interstitial sites in the host lattice and vacancies refer to the absence of an atom at a lattice site. These native point defect diffusion mechanisms are represented Figure I.44. In the vacancy mechanism, the diffusing atom jumps into a neighboring vacancy lattice site, while in the interstitial mechanism, an interstitial atom jumps from an interstitial site to another one. Defects such as vacancies and interstitials may be present at thermal equilibrium, with Si supporting both vacancies and self-interstitials while in Ge, the vacancy diffusion mechanism predominates. These diffusion mechanisms can simultaneously contribute to diffusion in a given environment.

Interstitial mechanism Vacancy mechanism

I.2.5.2. Atomic-Scale Diffusion Mechanisms

Interdiffusion refers to diffusion processes in alloys with composition gradients. As will be shown Section I.4.1., Si selective oxidation of SiGe induces interdiffusion in the underlying SiGe layer, hence the description of the main properties of SiGe interdiffusion in this section.

While interdiffusion is a single process described by a single interdiffusion coefficient, interdiffusion is actually driven by diffusion of Si and Ge atoms which have two different intrinsic diffusion coefficients [Merher17]. This apparent contradiction is solved when the reference frame and a vacancy flux is considered (the so-called Kirkendall effect) is considered. The interdiffusion coefficient accurately describes interdiffusion in a reference frame that is not the one of the laboratory, but the one moving with the vacancy flux. We also emphasize that the SiGe interdiffusion coefficient depends on the Ge concentration because the chemical environment of SiGe depends on the Ge concentration.

The first theoretical and mathematical description of interdiffusion in a binary alloy that takes into account the Kirkendall effect was proposed by Darken in 1948 [Mehrer07]. Without entering the details of his demonstration, the two Darken's equations relate the interdiffusion coefficient to the intrinsic and self-diffusion coefficients of the two elements of the binary alloy. Darken's equations provide a convenient framework for solving interdiffusion problems, such as the one of interdiffusion in SiGe. Darken's framework will be used for the numerical simulations of SiGe oxidation of Section III.B.2.. Generally speaking, interdiffusion is SiGe increases with the Ge concentration because of the increase of the number of weaker Si-Ge and Ge-Ge bonds compared to the number of stiffer Si-Si bonds. Compressive strain also enhances SiGe interdiffusivity [START_REF] Daillant | X-ray and neutron reflectivity: principles and applications[END_REF]. We highlight that SiGe interdiffusivity is approximately exponentially dependent on both temperature, following an Arrhenius law, and Ge content.

I.3. Properties of Thermal Oxides Grown on Si and SiGe

In this manuscript, SiO 2 always refers to amorphous thermally grown SiO 2 , also called fused silica. The usual SiO 2 phase produced in CMOS manufacturing oxidation processes is composed of corner-shared SiO 4 tetrahedra connected together in a three-dimensional network. As illustrated Figure I.47 and Figure I.48, the resulting structure features a strong short-range order with an average Si-O-Si bond angle of 145° and a bond angle range of 110° to 180°, but no long-range order. Thermally grown SiO 2 is thus considered amorphous. However, Tatsumura et al. and Munkholm et al. showed by performing X-Ray Scattering that the bulk of thermal SiO 2 retains the residual crystalline order of {111} planes of its parent Si crystal [Tatsumura04][Tatsumura05] [START_REF] Nakaharai | Characterization of 7-nm-thick strained Ge-on-insulator layer fabricated by Ge-condensation technique[END_REF]. They showed that there is no other residual crystalline order but the one originating from {111} planes and that high temperature annealing makes it disappear. They concluded that this residual crystalline order is sufficiently weak so that thermal SiO 2 should be conventionally considered amorphous.

I.3.1. Bulk Thermal SiO 2

The transition structure at the Si(100)-SiO 2 interface (about 1 nm-thick) cannot be considered abrupt in views of the oxide thicknesses studied in this work (from 1 nm to a few tens of nanometers). No atomic-scale structure model has gained widespread acceptance but several experiments and atomic-scale theoretical models have shed light on some of its properties, that are briefly introduced below: showed the existence of an approximately 1 nm-thick interfacial layer of higher density (between 2.3 g.cm -3 and 2.4 g.cm -3 ) ȃ Various atomistic models of the Si(100)-SiO 2 interface have been constructed using simulations methods [Bongiorno03].

I.4.2. FD-SOI Process Flow Integration

Composition homogenization by interdiffusion

Si selective oxidation heating system

One of the key advantages of the condensation technique for the FD-SOI process flow shown Figure I.50 is that co-integration of Si-based nMOS and compressively strained SiGe-based pMOS transistors can be performed. Indeed, a SiN hard mask is deposited on top of nMOS areas to act as an oxidation barrier. We can highlight that, at the end of the condensation process, an area at the frontier between nMOS and pMOS regions is not homogeneous because of the lateral interdiffusion triggered by the thermal anneal of the condensation process. This issue is overcome thanks to the Shallow Trench Isolation (STI) fabrication step. As illustrated Figure I.51, the STI is fabricated in the frontier area between nMOS and pMOS areas and is larger than the width of this non-homogeneous area.

I.4. The SiGe Condensation Technique

The condensation technique has been proposed for the first time in 2001 by Tezuka et al. [Tezuka01] from the Toshiba-MIRAI group as an innovative method for producing high-quality Germanium-On-Insulator (GeOI) or SiGe-On-Insulator (SGOI) buffer layers for the growth of tensely strained Si channels. In 2003, the idea to use the condensation technique to form a defect-free and compressively strained SiGe channel on insulator was published [Nakaharai03]. This technique was labelled "condensation" because the Ge content in the final SiGe layer may be higher than the one in the initial epitaxial layer. An appropriate term could have been "enrichment" instead of "condensation", because condensation also refers to the phase change from a gaseous phase to a liquid phase. It has already been shown that the condensation technique is a powerful and convenient technique for local fabrication of fully strained SGOI layers. However, several challenges still need to be taken up for an efficient and optimized use of the technique in an industrial environment. Two main challenges were identified during this thesis work and are addressed in Chapter III and Chapter IV respectively. These two problematics are then split into two more fundamental studies (Section III.A and Section IV.A) and two other problematics that have more direct applications to advanced FD-SOI CMOS manufacturing (Section III.B and Section IV.B). They are presented in the following sections.

The condensation technique is described

I.5. Problematics of this Thesis Work

The final SGOI layer composition and thickness must be well controlled for industrial CMOS manufacturing. Considering that the whole quantity of Ge is conserved during condensation (this will be shown Section IV.A.3.), the final SGOI composition depends on the amount of Si that is oxidized (this will be shown Section IV. A.1.). The total quantity of oxidized Si then depends on the thickness of the grown thermal oxide. It is therefore crucial to understand SiGe thermal oxidation kinetics when using industrial oxidation processes.

In spite of the widespread attention that has been devoted to the matter, there is an obvious lack of consensus regarding the oxidation kinetics of SiGe compared to that of Si. Indeed, a higher oxidation rate has been reported for dry oxidation of SiGe [Holland87][Kilpatrick97][Spadafora03], while some other authors have observed no difference between SiGe and Si [LeGoues89][Nayak92][Vincent07][Long12]. Several mechanisms have been proposed to explain oxidation kinetics of SiGe with respect to Si, but none of them has gained widespread acceptance. Moreover, the thickness of grown oxides range from a few nanometers to tens of nanometers depending on each publication, which makes comparison between published results difficult. Section III.A thus aims to bridge the gaps between these studies by covering various technology-relevant oxidation processes. We thus address in Section III.A the mechanisms and kinetics of thermal oxidation of SiGe by means of dry furnace oxidation, dry RTO and ISSG oxidation.

The second part of this chapter (Section III.B) presents two tools built from to the oxidation kinetics results of Section III.A. The first one is a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, a key parameter of the evolution of the SiGe composition during oxidation, as a function of process parameters. It is designed for a convenient and quick evaluation of the relative importance of the Ge pile-up mechanism with respect to SiGe interdiffusion. Accurate prediction of the oxide thickness and SiGe composition profile by numerical simulation can be useful for both a finer understanding of the dynamics of oxidation (oxide growth and evolution of the SiGe composition profile) as well as for optimization of industrial processes. This is also addressed in Section III. B. I.5.1.1. Study of the Mechanisms and Kinetics of SiGe Thermal Oxidation I.5.1.2

. Application to Industrial Processes for Advanced FD-SOI CMOS Manufacturing

I.5.1. SiGe Thermal Oxidation

As seen Section I.1.3., high Ge content and fully strained SGOI layers are required for enhancing the hole mobility of pMOSFET FD-SOI transistors. Any atom displaced from its lattice site disrupts the crystal symmetry and may degrade the hole mobility as a consequence. Extended defects are recombination centers that degrade the mobility and may even prevent the device from functioning properly. It is therefore crucial to understand how the condensation process influences the final strain level and crystal quality of SGOI layers.

Strain and dislocations formation in SGOI layers formed by condensation as well as their dependence on initial and final layer properties (Ge concentration and SiGe layer thickness) and condensation process conditions have been assessed by several groups

[Bedell04][Nakaharai06][Vincent07][Takagi10][Gunji11].
Most of these studies showed that the SGOI layer remains fully strained until a peak value, which depends on the initial SiGe concentration, SiGe thickness and condensation process conditions, is reached. The onset of strain relaxation was shown to be accompanied by dislocations formation and/or film buckling. However, the type of dislocation and the onset for for their formation vary from one publication to another. Moreover, these studies focused on condensation processes based on dry furnace oxidation, while lower thermal budget oxidation processes, such as dry Rapid Thermal Oxidation (RTO) and In-Situ Steam Generation (ISSG) oxidation, are preferred in today's semiconductor manufacturing.

We thus propose in Section IV.A to compare the strain level and the crystal quality of SGOI layers formed by condensation processes based on two industrial oxidations processes, i.e. dry RTO and ISSG oxidation.

I.5.2.2. Conservation of Strain of SGOI Films during Co-Integration of nMOS and pMOS Transistors

I.5.2.1. Strain and Crystal Quality of Full-Sheet SGOI Films Fabricated by Condensation

Strain, SGOI crystal quality and electrical quality of SGOI-BOx and SGOI-gate stack interfaces are required at the very end of chip manufacturing in order to reach high pMOS electrical performances. Indeed, further processing to fabricate both nMOS and pMOS on the same chip may alter the electrical properties of the SGOI channel. The effects of processes that immediately follow the condensation process (see Figure I.50) on the properties of the

I.5.2. Strain and Crystal Quality of SGOI Films Fabricated by Condensation

SGOI layer at the end o f chip manufacturing may be strongly correlated to the properties of the SGOI layer right after the condensation process.

As will be detailed in Section IV.B, strain relaxation of the SGOI layer after STI etch at the edge of pMOS active areas has been evidenced and shown to be detrimental to final pMOS electrical performances. It was suggested that the condensation process, which is performed a few processing steps before the STI etch, is most probably related to this phenomenon [Usuda13] [START_REF] Cooper | Strain mapping of semiconductor specimens with nm-scale resolution in a transmission electron microscope[END_REF]. Two possible causes based on properties of the SGOI layer and of its interfaces with the thermal oxide and the BOx layer were identified and are thus studied in Section IV.B.

Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 56/230

Conclusion of Chapter I

The MOSFET technology has been introduced and its limitations to further downsizing have been described. New transistor architectures and high mobility channels, two main paths for enhancement of transistor performances, were reviewed. Among these performance "boosters", the benefits of using compressively strained SGOI films as channel material for p-type MOSFETs were presented. Specificities of the condensation technique, which allows the fabrication of SGOI films with possibly high strain levels, were described. Finally, the two main challenges of this thesis, which are addressed in Chapter III and Chapter IV, were defined:

ȃ Understanding and controlling the mechanisms and kinetics of thermal oxidation of SiGe.

ȃ Fabricating highly strained and high crystal quality SGOI films by the condensation technique.

Introduction to Chapter II

Various processes and characterization techniques and a simulation tool were used to address the four problematics defined in Section I.A.. Techniques I personally used, either for the measurement and/or the analysis, are described in this chapter. Techniques used for characterizations I was only requester of and not involved with are presented in Appendices.

This chapter starts by describing the processes used to fabricate samples, with a focus on the thermal oxidation processes used in this work, i.e. dry furnace oxidation, dry Rapid Thermal Oxidation (RTO) and In-Situ Steam Generation (ISSG) oxidation.

Then, characterization techniques using equipment considered as "common" equipment are presented. Ellipsometry is the most used technique to measure thicknesses of films fabricated during this thesis work, especially thermal oxides in Section III.A.. X-Ray Reflectivity (XRR) measurements using standard tools were also performed to measure the thicknesses of various films, including thermal oxides for comparison with ellipsometry, and especially SGOI films fabricated by condensation in Chapter IV. XRR is only briefly presented as this technique is more thoroughly described later in Section II.4.1.. X-Ray Spectroscopy (XPS) allowed us to study the chemical composition of thermal oxides grown on SiGe in Section III.A.3.. Raman spectroscopy is a powerful technique for evaluation of strain, especially for very thin films (10 nm and below) when X-Ray Diffraction (XRD) is not adequate. Raman spectroscopy was thus used to measure strain levels of SGOI films fabricated by condensation in Section IV.A.. And finally, Geometrical Phase Analysis (GPA) of High-Resolution Transmission Electron Microscopy (HRTEM) images is introduced. We performed this analysis on a few SGOI samples of Section IV.A. in order to compare strain levels retrieved by GPA with the ones measured by Raman spectroscopy, and in order to reveal crystal defects such as dislocations.

Techniques that required the use of large-scale facilities are then described. It starts with XRR experiments performed on SEXTANTS beamline at Soleil synchrotron in order to measure the density of thermal oxides (Section III. A.5.). Neutron Reflectivity was used to identify whether or not there is hydrogen in SGOI films. The experiments were performed on D50 beamline at Institut Laue Langevin (ILL) nuclear facilities. Medium Energy Ion Scattering (MEIS) experiments were carried out at MEIS facilities at LETI, Grenoble, to study the depth distribution of crystal perturbations and defects of SGOI films in Section IV.A.. Finally, a simulation tool based on the Finite Element Method (FEM) was developed in COMSOL environment in order to be able to perform numerical simulations of SiGe oxidation (Section III. B.2.). The FEM and the environment of the simulation tool are thus described in this chapter.

II.1. SiGe Epitaxy Processes

For hetero-epitaxy on bulk Si, thin SiGe films were grown on 300 mm bulk {001} Si substrates in an industrial single wafer reactor (Centura from Applied Materials) using Rapid Thermal Chemical Vapor Deposition (RTCVD) techniques. All SiGe epitaxy processes used in this work are based on processes developed at STMicroelectronics before this thesis work.

Prior to deposition, Si surfaces were prepared by an in-situ high temperature bake at 1040°C for 120 seconds. Two epitaxy processes were used to grow SiGe films of two different compositions. Si 0.7 Ge 0.3 films were grown at 630°C using H 2 gas carrier and Germane (GeH 4 ) and to Dichlorosilane (DCS, SiH 2 Cl 2 ) precursors. For growth of Si 0.9 Ge 0.1 films, Silane precursor (SiH 4 ) was used instead of DCS and the growth temperature was raised to 700°C in order to reach a reasonable growth rate.

As for hetero-epitaxy on SOI substrates, another surface preparation was used because the high temperature bake would roughen or disrupt the SOI film. Firstly a standard HF-RCA clean was performed, and an in-situ dry cleaning process called SiCoNi [Labrot16] was then used. SiGe with about 30% Ge concentration were grown at 630°C using DCS and GeH 4 precursors.

These epitaxy processes were performed by Clément Pribat and Yves Campidelli from STMicroelectronics. Films composition, thickness and roughness were measured by X-Ray Diffraction (XRD) and Spectroscopic Ellipsometry techniques.

II.2. Thermal Oxidation Processes

In semiconductor manufacturing, the tolerance of chips being fabricated to high temperature and/or long oxidation durations depend on the target application. The impact of a thermal process in the chip is usually assessed by an abstract and qualitative notion called thermal budget. The thermal budget may be represented by the couple temperature-duration. The higher the oxidation temperature and duration are, the higher the thermal budget is. Establishing relative comparisons of the thermal budget of various oxidation processes is usefull for developing processes for industrial semiconductor manufacturing.

Two main families of oxidation processes are distinguished: furnace oxidation which has the highest thermal budget because of slow temperature ramp up and down (see Table II.1), and Rapid Thermal Oxidation (RTO). RTO systems use incoherent light from a honeycomb halogen lamp source to heat up a single wafer located below the lamp head (see Figure II.1). Such a system offers drastic reduction of the thermal budget compared to furnace oxidation thanks to its much lower thermal mass. Less common oxidation processes, such as low temperature oxidation under UV or laser light exposure, are not examined in this manuscript. Prior to any oxidation process, a standard HF-SPM-SC1-HCl cleaning process was performed. It resulted in the formation of an approximately 1 nm-thick chemical oxide on top of the SiGe film.

II.2.1. Introduction

II.2.2. Dry Furnace Oxidation

Furnace oxidation was performed at 900°C in a TEL FORMULA Furnace. The temperature is the one of the O 2 gas. For each SiGe oxidized film, Si monitor wafers were oxidized in the same conditions. To ensure the relevance of a comparison with RTO processes, the maximum temperature ramp allowed was used and the chamber was filled with O 2 at the end of the temperature ramp up in order to reach one atmosphere of pure oxygen at the beginning of the oxidation plateau. A N 2 purge was immediately started at the end the oxidation plateau. Oxidation during the temperature ramps was thus kept to a minimum. A typical temperature-time profile of a dry furnace oxidation process is shown

II.2.3. Rapid Thermal Oxidation (RTO)

II.2.3.1. Dry Rapid Thermal Oxidation (RTO)

For RTO processes, it has been shown that the heat redistribution length is much larger than the light absorbing depth, and therefore that the temperature is uniform across the wafer thickness and that it can be extracted by measuring the emissivity of the back side of the wafer (by use of an instrument called pyrometer) [Borisenko97]. The wafer is also constantly rotating to allow good circular uniformity. Two RTO processes were used in this study.

A typical temperature-time profile of a dry RTO process is shown Figure II.3 and has the following steps:

ȃ It starts with a first slow temperature ramp up until reaching 550°C. In this regime, the power of the lamp is handly fixed. ȃ Then, there is a plateau of a few seconds at 550°C to allow the temperature to be uniform over the whole wafer surface and the temperature monitoring system (based on a pyrometer as described above) to stabilize. ȃ During the second temperature ramp up, the power of the lamp is controlled by a closed loop system that constantly compares the set temperature-time profile and the measured one. The temperature ramp up is usually around 50°C.s -1 . ȃ When the target temperature is reached, there is a plateau in the temperature-time profile controlled by the closed loop. The oxidation time given in this manuscript is the duration of this plateau. ȃ The first temperature ramp down until 550°C is controlled by the closed loop. ȃ Below 550°C, the temperature ramp down is in open loop with a handly set lamp power. Dry RTO was performed in a single wafer rapid thermal equipment (Radiance chamber from Applied Materials). The oxidizing atmosphere was 740 Torr of pure O 2 and the temperature of the wafer was monitored by a closed loop system taking as input the measurement of the emissivity of the back side of the wafer. The set atmosphere inside the chamber was reached before the temperature ramp up. Temperature ramp up and ramp down were between 25°C/s and 75°C/s to ensure that oxidation during the temperature ramps is negligible compared to oxidation during the oxidation plateau.

Study of Thermal

ISSG oxidation was also performed into a Radiance chamber from Applied Materials. For this process, a mixture of 33% H 2 and 66% O 2 atmosphere with a total pressure of 7 Torr was used. Temperature ramps and monitoring procedure were similar to the ones of dry RTO processes. Below 600°C, the atmosphere was pure O 2 at low pressure, and H 2 was added only after the temperature exceeded 600°C, as illustrated on th temperature-time profile Figure II.4.

II.2.3.2. In-Situ Steam Generation Oxidation

II.3. Characterization Techniques

Ellipsometry is a non-destructive optical characterization technique based on the measurement of the change of light polarization upon a reflection on a surface sample. Ellipsometry can be separated into two modes: Single Wavelength Ellipsometry (SWE) and Spectroscopic Ellipsometry (SE). SWE is commonly considered as the characterization technique for measuring film thicknesses since it is non-destructive, relatively fast, possibly inline, and one of the most accurate thickness measurement techniques. However, it is highly model-dependent and input parameters must be carefully calibrated. SE allows to study more complex structures such as multilayer structures and allows access to data such as roughness or homogeneity.

The output of an ellipsometer is composed of two ellipsometric angles, Ψ and Δ, which are related to the optical indices of the sample. These angles are defined by the relation:

r r i p s = ( ) tan exp , ψ ∆
where r p and r s are the Fresnel reflection coefficients for a parallel and a perpendicular polarization respectively. The sample is described as an optical element by its Jones matrix:

J r r p s =       0 0 .
A schematic illustration of the principles of ellipsometry is shown We used SiGe optical constants published by Nolot et al. [Nolot14] and treated the composition of the SiGe layers as unknowns. When non-homogeneous SiGe layers were expected, two or three SiGe layers were used for the fitting model in order to improve the fitting quality.

These measurements were performed with the help of Francesco Abbate from STMicroeletronics.

X-Ray Reflectivity (XRR) is a thin film characterization technique based on the measurement of the reflectance of a multilayer system. The reader is referred to Section II.4.1. for a more thorough description of the technique. In-line (inside STMicroelectronics Crolles clean room) and off-line (at LETI) XRR characterization has been performed to measure the thicknesses of layers of various samples of this study. 0°-4° θ scans were acquired on a commercial Jordan Valley JVX7300 equipment for in-line measurement and on a Empyrean Panalytical equipment for off-line measurement. The X-Ray source uses the K-alpha emission line of copper, at a wavelength of 1.54056 Å. Reflectivity spectra were fitted with the commercial software of each equipment (JVXRR Simulator for the Jordan Valley and X'Pert Reflectivity for the Empyrean Panalytical). Fitting was limited to 0°-2° θ scans because the signalto-noise ratio was too low above 2°. Similarly to spectroscopic ellipsometry, two or three SiGe layers were used for the fitting model in order to improve the fitting quality when non-homogeneous SiGe layers were expected.

II.3.2. X-Ray Reflectivity (XRR)

X-Ray Photoelectron Spectroscopy (XPS) is a surface-sensitive chemical analysis technique based on the measurement of the kinetic energy of electrons ejected from the sample surface upon irradiation by X-Ray photons of known energy hν. This effect, known as the photoelectric effect and evidenced by Hertz (1887) and explained by Einstein (1935), corresponds to the ejection of an electron, called photoelectron, from its atomic orbital as a consequence of the absorption of a photon whose total energy is transferred to this electron. The incident photon energy (in the keV range) being known, the measurement of the kinetic energy of the photoelectron allows extraction of the binding energy of the core electron (see Figure II.7).

The binding energy of a core electron depends on both the element and the chemical environment of this element (such as its oxidation state). Careful analysis of XPS peak energies and shapes thus permits measurement of the composition of a sample surface and gives information about the chemical environment of each element. The depth sensitivity of the XPS technique is limited by the inelastic mean free path of an electron in a material, which is about 10 nm. XPS peaks are labelled with the element, the orbital of the ejected core electron, and when necessary information on the chemical environment of the element. For the quantitative analysis of Ge3d spectra acquired with the in-line XPS tool at STMicroelectronics Crolles clean room, a Shirley background from 27 eV to 31 eV was subtracted from the raw data. All spectra shown in this manuscript are data after subtraction of the background. Each spectrum is then decomposed into four peaks: ȃ A metallic Ge3d peak around 29.9 eV [START_REF] Muller | The electronic structure at the atomic scale of ultrathin gate oxides[END_REF]. ȃ A peak around 34 eV that we associate with the fourth oxidation state of Ge, and that we label Ge3d4+. We indeed always found a chemical shift of this peak with respect to the metallic Ge3d one of about 4.25 eV, close to values reported in the literature for the fourth oxidation state of Ge [Prabhakaran00]. We therefore fixed the chemical shift of the peak to 4.25 eV with respect to the metallic Ge3d peak. ȃ A Ge3d peak with an intermediary oxidation state was used for a proper fitting result only and referred as Ge3d 1-3+ . ȃ An O2s peak at about 26 eV, which may also include hydroxide components [START_REF] Muller | The electronic structure at the atomic scale of ultrathin gate oxides[END_REF].

II.3.3. X-Ray Photoelectron Spectroscopy (XPS)

All peak shapes were chosen to be pseudo-Voigt functions that allowed proper fit for all acquired spectra. XPS spectra were acquired on the direction normal to the sample surface with an in-line XPS Revera Vera-Flex II in STMicroelectronics Crolles clean room for the study discussed Section III.A.3. and on a direction at 45° with respect to the normal to the surface with a Versaprobe II at LETI for the study of Section IV.B.2.. Both equipements use a monochromatic source at 1486.6 eV. The measurements on the Versaprobe II were carried out by Eugénie Martinez from LETI.

Parallel Angle-Resolved XPS (pAR-XPS) is a specific XPS technique that uses the angular dependence of the photoemission spectra to provide chemical information as a function of depth. If the acquisition angle with respect to the normal of the sample surface is high, photoelectrons emitted at a depth of 10 nm would go through a much longer path (tens of nanometers) to escape the sample surface than a photoelecron emitted from the very top atomic layers. Using a high acquisition angle thus lowers even more than with a zero angle the intensity of the signal coming from buried atomic layers with respect to the surface. As a consequence, a high acquisition angle with respect to the normal of the sample surface provides information about the very top part of the sample.

In this study, pARXPS measurements were performed at 10 -9 mbar in a customized ThermoScientific Theta 300 pARXPS tool at LETI using a 1486.6 eV X-Ray source. The lens design and the two-dimensional detector allowed us to perform a parallel acquisition of the photoemission signal at various angles with respect to the direction normal to the sample surface. Eight acquisition angles from 23.75° to 76.25° were used. Acquisitions and analysis of pAR-XPS spectra were performed by Laurent Fauquier [Fauquier17].

Raman spectroscopy is a technique used for the observation of vibrational modes of atoms. The technique is based on the inelastic scattering of light, the so-called Raman scattering. The wavelength of the incoming monochromatic light, which is obtained from a laser, is therefore shifted, and the analysis of this energy shift gives information about the composition and/or strain of a material. In this study, Raman measurements were performed using a 363 nm laser beam focused on the sample using a x100 short working objective (0.9 numerical aperture). Raman spectra were acquired in backscattering geometry using a Jobin Yvon T64000 Triple Monochromator equipped with liquid N 2 cooled charge coupled device detector. The penetration depth of light at 363 nm into Si is about 15 nm, thus higher than the thicknesses of characterized SiGe layers. The presence of a thin SiO 2 layer on top of the SiGe layer does not interfere with Raman measurements as SiO 2 is transparent at 363 nm.

A Raman spectrum of a SiGe film is shown Figure II.9 where three main peaks corresponding to the Si-Si, Si-Ge and Ge-Ge vibrational modes are identified. Figure II.9: SiGe Raman spectrum and its decomposition into peaks corresponding to the Si-Si, Si-Ge and Ge-Ge vibrational modes.

II.3.4. Raman Spectroscopy

The frequency shift of each mode depends on both the composition and strain of the SiGe layer through the following relations:

ω ε ω ε i i i x x ( , ) ( ) b , = + 0 

Equation II.3

Chapter II -SiGe Thermal Oxidation Processes and SGOI Films Characterization Techniques

69/230

where i stands for Si-Si, Si-Ge, or Ge-Ge, ω i (x,ε) is the measured Raman frequency, ω 0i (x) the Raman frequency of a SiGe layer with the same composition but fully relaxed, b Si-Si = -850, b Si-Ge = -570, b Ge-G e = -450. b i were extracted from measurements of Raman spectra of fully strained SiGe layers and SiGe virtual substrates [START_REF] Schäffler | High-mobility Si and Ge structures[END_REF]. However, these empirical relations hold only for Ge concentrations below 60% because fully strained SiGe layers with Ge concentrations higher than 60% cannot be grown on Si substrates. Nevertheless, we used these relations for Ge concentrations higher than 60% because a similar tendency is expected for the Raman frequency shifts, which allows qualitative discussion above 60%.

In this study, ω 0i (x) were measured on SiGe virtual substrates prior to measurements of the SGOI samples. Figure II.10 shows the measured frequencies of the three modes of SiGe virtual substrates as a function of the Ge content, along with the fitted curves that allow extrapolation of the frequencies at any Ge concentration. Therefore, knowing the dependence of ω 0i (x) on x for each mode, and having two unknows, which are x and ε // , we could theoretically extract both x and ε // of each sample by using two of the three equations above. However, the "bell shape" of the Raman frequency of the Si-Ge mode as a function of the Ge content makes the use of this mode very inaccurate. Secondly, because the number of Ge-Ge bonds in low Ge content SiGe layers is low, the Ge-Ge mode peak signal is too low for accurate fitting. Therefore, only the Si-Si mode can be used for low Ge content SiGe layers. On the other hand, the Si-Si mode cannot be used for high Ge content SiGe layers for the same reasons.

As a consequence, one of the two unknown must be measured by another technique in order to extract the other one from Raman spectroscopy. Knowing the Ge concentration of the SGOI layers using another technique, we could extract strain by using the Si-Si mode for Ge concentrations below 60% and the Ge-Ge mode above 60%. Uncertainties in strain levels were calculated assuming the uncertainty of wavelength measurement is the monochromator resolution, i.e. 1 cm -1 .

Measurements were performed at LETI by Denis Rouchon [START_REF] Schäffler | High-mobility Si and Ge structures[END_REF]. Figure II.10: Measured Raman frequencies of Si-Si, Si-Ge and Ge-Ge modes of SiGe virtual substrates [START_REF] Schäffler | High-mobility Si and Ge structures[END_REF].

Geometric Phase Analysis (GPA) is an image analysis treatment that consists in retrieving and processing the geometric phase of High Resolution TEM (HRTEM) images [Hÿtch98]. The geometric phase of the image should not be confused with the phase shift experienced by the electrons going through the TEM lamella. Since a HRTEM image can always be described as the intersections of several systems of fringes, the method can be simply illustrated Figure II.11 on a 1D example. Figure II.11 (a) and (b) show a hypothetic HRTEM image with 1D fringes and the contrast of the image plotted along a horizontal profile respectively. As far as small deformations are considered, strain is contained in the phase of the fringe system. After retrieving the phase by numerical processing of the HRTEM image, the displacement fonction u(x) is obtained by substraction of the phase of a reference area (unstrained area) to the one of the distorted area. Strain is then calculated according the Equation II.4 by chosing two non-collinear g vectors, g α and g β : (b)

ε ε αβ βα α β β α = = ∂ ∂ + ∂ ∂         1 2 u x
The complete formulation of the GPA method is firstly based on the decomposition of the intensity of the lattice fringes of the HRTEM image into local Fourier components [Hÿtch98]. The intensity is then written as:

I I e r r g g g r ( ) = ( ) ∑ ⋅ 2π
, where g are the reciprocal lattice vectors describing the undistorted lattice. The local Fourier components are obtained by filtering in Fourier space:

I A e iP g g r r r g ( ) = ( ) ( ) ,
where the amplitude A g r ( ) describes the local contrast of the fringes and the phase P g r ( ), their position. The phase can be expressed as a function of the displacement field u r ( ) by the following relation:

P g r g u r ( ) = -⋅ ( ) 2π .
The two-dimentional displacement field is then calculated by measuring two phase images corresponding to two non-colinear reciprocal lattice vectors g α and g β : u r r a r a g g

( ) = - ( ) + ( )     1 2π α β α β P P ,
where a α and a β are the basis vectors for the lattice in real space which correspond to the two reciprocal lattice vectors g α and g β .

GPA analysis of HRTEM images of various samples was peformed using the GPA plug-in proposed by Professor Christoph T. Koch in Digital Micrograph environment. The plug-in code is based on equations of [Hÿtch98]. The following procedure was used to get deformation maps from HRTEM images: ȃ FFT of the HRTEM image (Figure II.12 (a)) to get the diffractogram of Figure II.12 (b)).

ȃ Selection of two g vectors and application of masks around these two frequencies ȃ Reverse FFT around the two selected spatial frequencies and extraction of the phase maps of Figure II.12 (c) after substraction of the phase of the chosen reference area (also shown Figure II.12 (c)). ȃ Deformation maps (Figure II.12 (d)) are obtained by combining phase maps. Shear deformation and rotations are also calculated but were always found to be uniform and close to zero so that they are not shown.

No quantitative uncertainties, deformation and spatial resolutions can be given because of the high dependence of the GPA method on the quality of the HRTEM images and on various perturbations that cannot be quantified.

Equation II.5 Equation II.6 Equation II.7 Equation II.8

Chapter II -SiGe Thermal Oxidation Processes and SGOI Films Characterization Techniques 71/230 HRTEM images were acquired at LMGP in collaboration with Laetitia Rapenne from LMGP and the GPA was performed with the help of Patricia Donnadieu from SIMaP. c) and (d) are phase maps associated with the two g vectors and showing the chosen reference area in the unstrained Si substrate. And (e) and (f) are deformation maps calculated from phase maps. Amorphous areas (the BOx layer, the thermal oxide on top and layers used for the TEM lamella preparation) are masked as the GPA is not relevant in these areas.

BOx

S O I S iG e

BOx

BOx

In-plane deformation

BOx

Out-of-plane deformation

Deformation scale (e) (f)

Compressive

Tensile Unstrained X-Ray Reflectivity (XRR) is a technique that probes the profile of optical index along the normal of a thin film system. It can also probe the roughness at the interface between two layers of different optical indices. In the specular geometry illustrated Figure II.13 (a), a monochromatic X-Ray beam is directed towards the surface of the sample with an incident angle θ, and the reflected intensity is measured at a scattering angle of 2θ with respect to the incident beam. The reflected intensity is measured as a function of the incident angle θ in a so-called θ-2θ scan, and plotted sometimes as a function of the modulus of the wave vector transfer q (see Figure II.13 (a)). As the reflected intensity rapidly decreases with θ, the typical θ scan range is 0°-5° for the most commonly used experimental set-up (such as the one presented Section II.3.2. that uses the K-alpha emission line of copper as the X-Ray source). However, for XRR experiments based on low-energy X-Rays produced with synchrotron radiation, angular scans up to a few tens of degrees can be performed.

As illustrated Figure II.13 (b), the reflected intensity is modulated by constructive and destructive interferences between waves scattered at different interfaces of a multilayer system. Thicknesses and optical indices of a multilayer system can be extracted from fitting the experimental reflectance with numerical models. As will be seen into the next section, the optical index of a material at X-Ray energies directly depends on the electron density.

II.4.1. Resonant Soft X-Ray Reflectivity (R-SoXR)

II.4.1.1. Introduction

II.4.1.2. Interaction of X-Rays with Matter in Reflectivity Experiments

Electromagnetic radiations, such as X-Rays, interact with electrons mainly but hardly with atomic nuclei. More precisely, the interaction is between the electric field of the radiation and the electron. Two main types of interactions can be distinguished: ȃ Elastic scattering, also called Thomson scattering, in which the incident wave is scattered by the charged particle, here an electron, without any energy loss. ȃ Absorption: all the energy is transferred to an electron which is excited to an empty upper state and the photon vanishes. The excited electron may be expelled from the atom (photoelectric absorption), or the atom releases the energy through the emission of another photon (fluorescence). Important transitions for X-Rays are the ones involving the electrons on the inner shells such as the K and L ones.

At the atomic level, these two interaction processes with an atom k are described by the complex atomic factor f k which is defined by the following equation [Daillant08]:

A A sca i nc = - ( ) r r P f k 0 φ ,
where A sca is the scattered electric field by an atom k, A inc the incident electric field, r 0 the classical electron radius, Φ the scattering angle, P(Φ) the polarization factor, and r the distance from the atom to the observation point.

Giving an exact and complete description of all these processes, only possible in the frame of the quantum

θ 2θ q k k 0 θ θ θ θ' n 1 n 2 n 3 (

II.4. Characterization Techniques using Large-Scale Facilities

Equation II.9

theory, is out of scope of this work. However, key results of this theory are used in the optical formalism described below. The interaction of X-Rays with matter in reflectivity experiments is presented according to an optical formalism in which the properties of materials are described by the refractive index. The use of the optical formalism is made possible because materials can be considered as a homogeneous continuum for X-Ray radiations even if they are actually not at the atomic scale. The index of refraction is commonly expressed as:

n i = -- 1 δ β,
where δ and β account for scattering and absorption of the material respectively. For X-Rays and far from any absorption edge, the index of refraction is very close to 1 (slightly lower). One can show that:

δ λ π ρ λ π λ = = + ( ) ∑ r r Z f V e e e k k m k 2 2 2 2 ' β λ π λ = ( ) ∑ r f V e k m k 2 2 '' f Z f if k k k k = + + ' ' ' ,
where r e is the classical electron radius of the electron, λ is the wavelength of the incident X-Ray wave and ρ e is the electron density, Z k is the number of electrons of the atom k in the unit cell, V m is the volume of the unit cell, and f k ' and f k '' are the real and imaginary parts of the anomalous factor at the specific wavelength λ.

Let us give some insight about the meaning of each factor of the index of refraction shown Equations II.11. In a classical description of Thomson scattering of X-Rays by the electrons of an atom k, and when assuming that the frequency of X-Rays is large compared to the atomic frequency, one can show that the atomic scattering factor f k of the atom k is simply its atomic number Z k . This contribution accounts for the Z k term in Equations II.11 and is much larger that f k ' and f k '' as long as the incident wave frequency is far from the any absorption edge of the element k. However, the frequency of the incident X-Ray wave may not be greater than the atomic frequencies for the inner electron shells such as the K and L ones. The correction factors f k ' and f k '' are introduced in order to take into account this effect. This correction is called dispersion correction.

II.4.1.3. X-Ray Reflection by Planar Multilayers

When the X-Ray wave propagates in a heterogeneous medium with regions of different electron densities, hence of different X-Ray optical indices, calculation of the reflectance is performed by applying the boundary conditions of the electric and magnetic fields at each interface. The fact that multiple reflections have to be considered for multilayer systems requires a matrix formalism. The reader is referred to [Daillant08] for the complete calculation based on the theory of reflection.

An example of the theoretical X-Ray reflectivity curve of a homogeneous SiGe layer on a Si substrate is shown Figure II.14. As the optical index is slightly less than 1 for X-Rays, total reflection is observed below a critical angle θ c that depends on the electron density of the top layer only. The oscillations, also called Kiessig fringes, are the results of interferences between the waves reflected at the surface and at the interface between the two layers. We can show that the period of these fringes depends on the layer thickness only, while the amplitude of the fringes depends on the optical contrast, that is on the electron density contrast between the two layers. The critical angle and the fringe period are indicated on Figure II.14.

One last important effect to consider in X-Ray reflectivity is the interface roughness. Roughness tends to decrease the total reflected intensity in the specular reflectivity geometry by scattering a part of the incident radiation towards non-specular directions. Roughness can be taken into account by introducing a transition layer.

Equation II.10 Equations II.11

Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 74/230 Figure II.14: XRR spectra of a SiGe layer on a Si substrate simulated at the wavelength of the K-alpha emission line of copper.

II.4.1.4. Resonant Soft X-Ray Reflectivity (R-SoXR)

Resonant X-Ray Reflectivity (R-SoXR) is called "resonant" because the energy of the incident X-ray radiation coincides and resonates with the absorption edge of an element of the multilayer system. By choosing specific X-ray energies around an absorption edge, R-SoXR can be used to tune and enhance the optical contrast between layers [Nayak15]. Indeed, large changes of reflectivity at an interface, and as a result of interferences in multilayer systems, are induced by the considerable variation of the complex index of refraction n around an absorption edge. These large changes arise from the correction factors f' and f''. An accurate determination of scattering factors of each element of the multilayer system is thus required for extraction of electron densities from R-SoXR spectra. As will be detailed into Section III.A.5.1.1., the imaginary part f'' of the scattering factor can be retrieved by absorption measurements and the real part f' is then calculated from the imaginary part using Kramers-Kronig transformation. Synchrotron radiation is needed as R-SoXR experiments require the use of several energies around an absorption edge.

II.4.1.5. Experimental

The XRR experiments were performed at SEXTANTS beamline at Soleil Synchrotron. The SEXTANTS beamline, which covers the 50-1700 eV energy range, is dedicated to the investigation of electronic and magnetic properties of materials using soft X-Ray scattering techniques. The experiments were performed using a dedicated diffractometer working under high vacuum (in the 10 -9 mbar) in order to lower soft X-Ray absorption by air: the RESOXS reflectometer [Jaouen04]. The RESOXS reflectometer features resolution and reproducibility smaller than 0.005°, largely sufficient for soft X-Ray reflectivity experiments where the θ range is large (up to 60° in these experiments). The intensity of the reflected beam is measured by a silicon photodiode.

Prior to reflectivity experiments, absorption of several samples was measured in Total Electron Yield (T.E.Y.) and Total Fluorescence Yield (T.F.Y.) modes in order to extract O scattering factors around the K absorption edge (The extraction procedure is detailed in Section III.5.1.1.). T.E.Y. refers to the measurement of the current generated in the sample by absorption of soft X-Ray photons while T.F.Y. refers to measurement of X-Ray fluorescence emission by a detector close to the sample. The escape depth for the fluoresced X-Rays is larger than the one of electrons generated by absorption. It is difficult to give precise values of the probed depth of each mode because it depends both on the energy of the incoming photon and on the probed material. The T.E.Y. mode usually probes between 1 and 5 nm while the T.F.Y. probes between 10 nm and 50 nm. The polarization of the incoming beam for XRR measurements was circular.

These experiments were carried out in collaboration with Patrice Gergaud from LETI and Nicolas Jaouen from Soleil synchrotron.

Chapter . Fitting Procedure XRR spectra were fitted using the Dyna software. Dyna, which stands for "Dynamic" X-Ray Reflectivity, is a program that simulates conventional, resonant and magnetic X-Ray reflectivity in the Matlab environment. The algorithm is based on Parrat formalism [Elzo12]. Densities, thicknesses and roughnesses of each defined layer of the stack can be either fixed or free parameters. Details and assumptions specific to the experiments performed during this thesis are presented in Section III. A.5.1.3..

II.4.2. Neutron Reflectivity

II.4.2.1. Introduction

In the past decades, neutron scattering was mainly used to determine positions and motions of atoms in condensed matter (1994 Nobel Prize to Shull and Brockhouse) as well as to study polymers and biological systems thanks to a high sensitivity to isotopes labeling (especially 1 H versus 2 H). In the past few years, neutron scattering started to exhibit usefulness for the characterization of materials in the semiconductor industry

[Diebold99][Se- gura15].
Neutron Reflectivity is a technique similar to X-Ray Reflectivity (XRR) in the sense that neutrons can also be reflected at interfaces between two different materials. The formalism presented in Section II.4.1. can be transposed to neutron reflectivity provided that scattering factors are not the ones seen by X-Rays but the ones seen by neutrons. We thus emphasize in Section II.4.2. aspects that are specific to neutrons.

The combination of neutron reflectivity with XRR allows the measurement of hydrogen content in multilayer stacks, the purpose of this technique in this work, while XRR alone is not sensitive enough for such a light element [Segura15].

II.4.2.2. Interactions of Neutrons with Matter in Reflectivity Experiments

Scattering of neutrons by nuclei comes mainly from the strong interaction. Far from any nuclear resonance, which is the case of the experiments carried out in this work, absorption is negligible so that neutron scattering is considered elastic. The reader is referred to [Daillant08] for details of neutron-matter interactions.

While XRR is usually described in an optical formalism based on either the scattering factor or the optical index, neutron reflectivity formalism usually uses the nuclear scattering length. For neutrons, Equation II.9 is written differently:

A A b e r sc in ikr = - - ,
where A in , A sc , b and r are the incoming wave amplitude, scattered amplitude, scattering length and distance from the scattering element. As a convention, the minus sign is in b definition in order for b to be positive for most elements. The scattering length characterizes the scattering power of an element. Contrary to X-Rays, the neutrons scattering length depends on the atomic number in a non-monotonous way, as illustrated Figure II.15. We emphasize that 1 H has a negative scattering length, which makes neutron reflectivity very sensitive to its presence.

Equation II.12

Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 76/230 Then, the index of refraction n seen by neutrons is:

n b  1 2 2 - λ ρ π ,
where λ, ρ and b are the neutron wavelength, atomic density and scattering length.

II.4.2.3. Experimental

Neutrons reflectivity experiments were carried out at Institut Laue Langevin (ILL) facilities, Grenoble, on D50 beamline, and within the Institut de Recherche Technologique (IRT) nanolec framework. Neutrons reflectivity experiments were performed by Jaime Segura, Robert Cubbit and Philip Gutfreund, all from ILL. As will be detailed in Section IV. B.3., the mass density must be measured by another technique to allow the determination of hydrogen concentration by neutron reflectivity (which is the purpose the experiments described in details in Section IV. B.3.). We measured the mass density of the same samples by XRR with Patrice Gergaud from LETI on a Empyrean Panalytical equipment (see Section II.3.2.).

Equation II.13

Medium Energy Ion Scattering (MEIS) is a technique derived from the Rutherford Backscattering Spectroscopy (RBS) technique used to measure the composition and to determine structural properties of thin films . While RBS uses ion energies in the MeV range, MEIS is based on ions accelerated at so-called medium energies, between 50 keV and 500 keV. MEIS benefits from the use of an electrostatic detector, which allows much better energy resolution than the silicon diodes of RBS systems, and which in turn allows better depth resolution (a few Å for MEIS against about 15 nm for RBS). MEIS is much less commonly used than RBS with a dozen tools worldwide only, and long acquisition durations (~10h) are usually required to get reasonable statistics. MEIS equipment can thus be considered as "quasi large-scale facilities".

Basically, a parallel H + or He + ion beam is accelerated between 50 keV and 500 keV towards the sample surface. Most of the medium-energy ions are implanted into the sample at depths much greater than the probed surface. However, a small fraction of these ions travel close enough to atoms of the samples so that they are scattered by the Coulomb interaction with nucleus. Then, another fraction of these scattered ions see their trajectories deflected backward and may consequently escape the surface of the sample.

Chemical composition profiles as well as some structural properties are extracted from the measurement of the energy and of the angle of ions backscattered by the atoms of the sample (see Figure II.16). Ions that encounter only one scattering process are considered for the MEIS analysis as a second scattering process before the ion

II.4.3.1. Introduction

II.4.3. Medium Energy Ion Scattering (MEIS)

escapes the sample surface is highly unlikely. The energy and angle of backscattered ions are governed by two types of ion-atom interactions, which are described in the next two sections. Specific angle configurations between the ion beam and the crystal orientation of the sample, called channeling conditions, are used to probe the crystal quality and are described in Section IV. A.5.1.. Ions that travel through a solid lose kinetic energy because of inelastic Coulomb interactions with the electrons of atoms of the sample. Ions that are backscattered by atoms thus lose energy during the inward travel until they are backscattered and during the backward travel until they escape the sample surface. At fixed incidence and collection angles, the deeper the atom that backscattered the ion is, the higher the lost kinetic energy is when the ion escapes the sample surface. The kinetic energy of the backscattered ions thus carry information about the depth at which they are backscattered.

The average kinetic energy of ions as a function of depth can be calculated based on the stopping power of each element of the sample, which is defined as:

ε ρ E dE dx ( ) = 1 ,
where ρ (g.cm -3 ) is the mass density of the material and x the depth (cm). The stopping power of compounds are calculated according to the Bragg-Kleemann rule:

ε ε C i i i N E w E ( ) = ( ) = ∑ 1 ,
where w i is the mass fraction of element i.

Figure II.17 shows the stopping powers of Si, Ge and O which are the three elements of the samples of this study, as a function of the kinetic energy of ions.

II.4.3.2. Electronic Energy Loss of Medium Energy Ions in Solids

Figure II.16: Schematic illustration of the principles of the MEIS technique. E 0 refers to the energy of incident ions, E to the energy of backscattered ions, and θ i and θ s to the incident and scattering angles respectively.

Studied layer

θ i θ s E 0 E H + /He + Figure II
.17: Calculated stopping powers of O, Si and Ge as a function of ion energy. Calculated from [START_REF] Kahng | Electric field controlled semiconductor device[END_REF].

Equation II.14 Equation II.15

The random nature of inelastic Coulomb energy loss induces some dispersion commonly referred to as straggling. Models that take into account the straggling effect for the MEIS energy range were developed . Gaussian distributions around the average energy at the certain depth are usually used as a consequence of the straggling effect.

II.4.3.3. Elastic Nuclear Collision of Medium Energy Ions with Nuclei in Solids

Ions that travel close enough to the sample nuclei undergo elastic nuclear collision (also called Coulomb or Rutherford scattering) with the nuclei. This interaction is said to be elastic because no energy is lost in the spatial system of the mass center of the ion+nuclei system. However, in the spatial system of the laboratory, some impulse is transferred from the ion to the nucleus during the collision process, which results in the ion losing some kinetic energy. The ratio of the ion final and incident energy is called kinematic factor and can be calculated by considering energy and momentum conservation in the ion+nuclei mass center system [Busch00]:

K M M M M M M S i n S i n S i n θ θ θ ( )= -( ) + ( ) + ( )           1 1 2 2 2 sin c os ,
where θ s is the scattering angle of the ion in the laboratory frame, M i the ion mass and M n the nuclei mass. The kinematic factor does not depend on the ion energy. Figure II.18 shows that the energy that is transferred from the ion to the nuclei increases with the scattering angle, hence the difficulty to separate the elements with the MEIS technique at small scattering angles. We also highlight that the heavier the nuclei is, the lower the transferred energy is at a fixed scattering angle.

Quantification of each element can be performed by considering the scattering cross section of each element, which represents the probability of collision. The scattering cross section, also called Rutherford cross section, in the laboratory frame for a problem of collision between two elements is:

σ θ πε θ θ E Z Z E e M M S i n i n S S , sin cos ( )=       -       +       4 4 4 1 2 0 2 2    -       2 4 2 1 sin s in , θ θ S i n S M M
where Z i and Z n are the atomic numbers of the ion and of the target nuclei respectively. For the MEIS energy range, there are some corrections taking into account the partial screening of the nuclei charge by the electrons of the target atom [START_REF] Kahng | Electric field controlled semiconductor device[END_REF]. We highlight that the higher the atomic numbers of the ion and the target nuclei are, the higher the cross section is. For a given ion type and a fixed energy, Ge has consequently a higher scattering cross section than Si.

Equation II.16

Equation II.17

Chapter II -SiGe Thermal Oxidation Processes and SGOI Films Characterization Techniques 79/230 II.4.3.4. Channeling Channeling is a mechanism that takes place when the ion beam is aligned along a crystallographic direction of a single crystal. Ions that are not scattered by the first atomic layer see their trajectories deflected towards the space between rows of atoms. As illustrated Figure II.19, this mechanism induces the formation of a shadow cone which drastically reduces the probability of an ion to be scattered by deeper atoms. Ions that are deflected into the space between rows of atoms have oscillating trajectories because of small angle scattering processes occurring when ions get close to atoms.

As a consequence, except from the first atomic layer, channeling reduces the number of scattering events with a reduction factor called χ min . The χ min factor is the percentage of atoms displaced from theoretical lattice sites. Vibrational amplitudes of atoms, surface rearrangements, or de-channeling of the ion beam through a top amorphous layer such as a native oxide layer make the channeling process not perfect. χ min ranges from about 4% for a high-quality Si crystal to 100% for totally amorphous materials. However, when there is an atom located at an interstitial site of a crystal, scattering on this atom occurs as described in the former section. Structural defects such as dislocations thus decrease the channeling efficiency. In other words, structural defects increase the χ min factor to intermediary values. Crystal defects profiles as a function of depth can be extracted from the comparison of MEIS spectra acquired in non-channeling conditions, also called random conditions, with MEIS spectra acquired in channeling conditions.

II.4.3.5. Experimental

The MEIS study was performed at the MEIS facilities of Minatec, Grenoble, with a 100 keV He + ion beam. The ion beam diameter was about 1 mm. The detector was an electrostatic detector with a 40° angle detection amplitude around the central angle. For instance, when the detector is set in a 90° scattering angle configuration, it detects ions that are backscattered with scattering angles between 70° and 110°. Narrow angle bands were selected and energy scans in this band were summed to get reasonable statistics. However, no more than a few consecutive scans at the same sample location could be used because ion bombardment tends to damage the sample. The ion beam location on the sample was thus regularly changed in order not to include any biased data into the final summed spectra. Several random and channeling configurations were used and will be detailed in Section IV.A.5.1.. Channeling conditions were fixed by finding the minimum intensity of the Ge peak.

Measurements were performed by François Pierre from LETI. As there is no fitting algorithm included in the program, simulated spectra were manually adjusted to the experimental spectra following the procedure described below: ȃ The incident angle was always the set value.

II.4.3.6. Simulation of MEIS Spectra

ȃ An energy shift less than 0.6 keV and a detection angle shift less than 0.5° were applied to each experimental spectra. The energy shift comes from slight variations of the ion source with time while the angle shift arises from limitations of mechanical precision. When both the Si surface peak and O surface peak could be identified, the actual energy and detection angle could be exactly calculated. Indeed, both elements are present in the first atomic layer of the sample, so that the energy of backscattered ions depend on two factors only, the nature of the element and the scattering angle. When the O surface peak could not be identified, the energy and detection angle shifts were adjusted between realistic values (0.6 keV and 0.5°) in order to get the best fit.

ȃ In random configuration, an iterative manual adjustment of all layers parameters was performed in keeping with realistic possible variations with XRR and ellipsometry results:

• Ions that are backscattered by the Ge atoms located at the very top of the SiGe layer lose energy through two processes: elastic Coulomb scattering induces an energy loss that is known once the ion incident energy and scattering angle are calibrated (see former step of the procedure), and inelastic loss through the inward and outward travels through the top oxide only. When fixing the top oxide composition to SiO 2 , its thickness could be determined by considering the energy of ions backscattered by Ge atoms at the top of the SiGe layer (which corresponds to the highest energy of the Ge peak).

• As will be detailed into Section IV.A.5.1.2., the Ge peak is isolated on the spectra while the Si peaks of the oxide, the SiGe, the BOx layer and the substrate overlap with each other. Moreover, Ge has a higher scattering cross section than Si and O, which results into higher statistics for the Ge peak even for low Ge concentrations. As a consequence, adjustment of the SiGe layer thickness and composition is more robust and easier when considering the Ge peak. Matching of the experimental and simulated Si peaks was checked for consistency considerations.

ȃ In channeling conditions:

• The composition, thickness and atomic density of each layer were fixed to the values determined by the analysis of random spectra. χ min factors were adjusted by considering the Ge peak for the reasons mentioned above, and consistency with the Si peak was also checked.

II.5.1. Introduction to the Finite Element Method (FEM)

The Finite Element Method (FEM) is described in depth in [START_REF] Yang | Performance dependence of CMOS on silicon substrate orientation for ultrathin oxynitride and HfO 2 gate dielectrics[END_REF]. Most space and time-dependent problems are described by laws of physics that are expressed by Partial Differential Equations (PEDs). In most cases, PDEs cannot be solved analytically. The most general definition of the FEM is a numerical method for solving PDEs. The FEM is based on the division, or discretization, of the system into a finite number of elements. Then, PEDs are approximated by linear equations for each discretized element. Solutions of these linearized PDEs are computed at each node of the system, which are the intersections of the discretized elements. The solutions, called discrete solutions, are extrapolated between the nodes and an approximate continuous solution is finally obtained to the non-discretized initial problem by merging the solutions for each element. The convergence to a unique solution is guaranteed by the definition of initial and boundary conditions well adapted to the problem and by a well-refined discretization of the system. The denser the discretization, also referred to as the mesh, the closer the approximate solution is to the actual solution. However, a denser mesh induces longer computing durations, so that a trade-off has to be found for each problem. The FEM was used to simulate thermal oxidation of SiGe.

II.5. Simulation by the Finite Element Method (FEM)

Oxidation kinetics of SiGe and Ge redistribution during SiGe oxidation by means of dry RTO and ISSG oxidation were studied by the FEM in COMSOL Multiphysics environment. The model was derived from the model built by Pierre Morin [Morin15]. The key elements of the environment of the model are:

II.5.2. Simulation of Thermal Oxidation of SiGe

ȃ Modeling environment: Two physical modules were used:

• Transport of Ge: the model of SiGe interdiffusion is based on Darken's equations.

• Deformed geometry: the deformed geometry module was used to model the growth of the oxide.

ȃ Geometry and Materials: a 1D geometry with moving interfaces was used.

ȃ Specific conditions: Initial conditions of the simulations are the initial layer compositions and thicknesses.

ȃ Mesh: A fine mesh was defined in Si and SiGe layers while a much wider mesh was chosen for the oxide because there is no Ge diffusion in the oxide and thus no need for calculation.

The details of the model will be presented in Section III. B.2.2..

Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 82/230

Conclusion of Chapter II

The processes to fabricate samples of this study, techniques that were used to characterize them, and the environment of a developed simulation tool were described in this chapter.

More precisely, the three oxidation processes used in this study were presented, i.e. dry furnace oxidation, dry RTO and ISSG oxidation. Then, characterization techniques using equipment considered as "common" equipment were reviewed: ellipsometry and XRR for thickness measurement, XPS for studying the chemical composition of thermal oxides grown on SiGe, Raman spectroscopy for measurement of strain of SGOI films, and GPA of HRTEM images for measurement of strain levels as well as for revealing crystal defects. Techniques that required the use of large-scale facilities were also reviewed: XRR experiments performed on SEXTANTS beamline at Soleil synchrotron in order to measure the density of thermal oxides, neutron reflectivity to identify whether or not there is hydrogen in SGOI films, and finally MEIS experiments to study the depth distribution of crystal perturbations and defects of SGOI films. This chapter ended with a description of the FEM and the environment of the developed simulation tool.

« La liberté est une chose toute bête, une maladie dont l'hygiène sociale la plus

III.B. Application to Industrial Processes for Advanced FD-SOI CMOS Manufacturing

III.B.1. Condensation Abacus: Oxide thickness and Ge Diffusion Length

III.B.2. Simulation of SiGe Thermal Oxidation

Conclusion of Chapter III

Introduction to Chapter III

We showed in Section I.5.1.1. that understanding and controlling SiGe oxidation mechanisms and kinetics is crucial for an efficient use of the condensation technique in an industrial environment.

The first part of this chapter (Section III.A) strives to bring more understanding of the mechanisms and kinetics of SiGe oxidation.

After a literature review of SiGe oxidation in terms of mechanisms, thermodynamics and kinetics, the composition of the oxide is investigated, with a focus on the formation of a small quantity of Ge oxides. As these Ge oxides may affect the condensation technique and/or further manufacturing processes, we identify the mechanisms and conditions that lead to their formation.

Then, an in-depth study of the kinetics of SiGe oxidations by means of three technology-relevant processes is carried out. Oxidation kinetics is correlated with the evolution of the SiGe composition during oxidation. The evolution of the SiGe composition during oxidation is discussed within the framework of a competition between the Ge pile-up mechanism that tends to increase the Ge concentration below the oxide and SiGe interdiffusion which tends to smooth the SiGe composition. The influence of the Ge concentration at the oxidation interface on the oxide density measured by X-Ray Reflectivity (XRR) and as a result on oxidation kinetics is then studied. The dependence of O 2 diffusivity on the oxide density is finally discussed.

The second part of this chapter (Section III.B) presents two tools built from to the former oxidation kinetics results. The first one is a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, a key parameter of the evolution of the SiGe composition during oxidation, as a function of process parameters. It is designed for a convenient and quick evaluation of the relative importance of the Ge pile-up mechanism with respect to SiGe interdiffusion. The second one is a numerical simulation tool of oxidation of SiGe which can be used both for a finer understanding of the dynamics of SiGe oxidation and for accurate prediction of output parameters of the condensation process (such as the oxide thickness and the SiGe composition profile in the SGOI film).

III.A.1. Literature Review

Oxidation comes from the archaic French word "oxide" (changed later into "oxyde") coined by L.-B. Guyton de Morveau (1737 -1816) andA . Lavoisier (1743. Lavoisier ( -1794)), while the archaic French word "oxide" derives from the Ancient Greek "όξύς" (sharp, acid, thrill). Oxidation refers to a loss of electrons in bonding between atoms, or, more generally, to an increase of the oxidation state of an element. It is considered that, when determining the oxidation state of Si in silicon oxide, the electrons involved in bonding with O atoms are attributed to O. Each bond between a Si atom and a O atom gives a higher oxidation state to Si. Si has four covalent bonds with O atoms in stoichiometric SiO 2 , which results in the oxidation state of Si in SiO 2 to be +4. The case of Ge is similar to Si.

The literature review of Section III.A.1. starts with thermodynamic considerations about the SiGe oxidation reaction in order to determine the expected oxide composition. Then, the Deal and Grove model of kinetics of Si oxidation is presented [Deal65]. Its limitations and attemps to revise it to better fit experimental results are summed up. Finally, SiGe oxidation kinetics with respect to Si is discussed and shown to remain a matter of debate.

III.A.1.1. Introduction

III.A.1.2. Thermodynamic Considerations

Early studies of oxidation of SiGe were performed in the 1970's [Balk71][Margalit72], but it was not until the late 1980's that models describing thermodynamics, physical mechanisms and kinetics were proposed. Experiments of oxidation of SiGe have been performed in steam, dry oxygen, at various temperatures, pressures and on various SiGe films.

The following thermodynamic considerations are made at atmospheric pressure since considered thermodynamic activities change only slightly with pressure [Paine91], and at 1000 K, a temperature close to those used in this work. We do not consider silicon and germanium sub-oxides (SiO x and GeO x ) which have much lower formation enthalpies. Therefore, we consider the ternary system Si-Ge-O and the following standard Gibbs energies of formation at 1000 K from [Barin77]:

III.A. Study of Mechanisms and Kinetics of SiGe Thermal Oxidation

In order to determine what is the stable oxide in the Si-Ge-O ternary system, we consider the following reaction:

GeO Si SiO Ge 2 2 + = +
The Gibbs energy of this reaction is -353.81 kJ.mol -1 at 1000 K, therefore the SiO 2 -Ge phase dominates and Ge is thermodynamically stable with SiO 2 .

Si 1-x Ge x is considered as an ideal solid solution whose Gibbs energy of formation is given as a function of x:

∆G Si Ge RT x x x x o x x 1 1 1 - = - ( ) - ( )+     ln ln ,
where R is the gas constant and T the temperature in Kelvin. In order to determine whether GeO 2 reduction by Si in Si 1-x Ge x is favorable, we consider the following reaction:

GeO Si S iO Ge Si Ge Si Ge x x x x 2 2 1 1 + = + - - ( )
Equation III.1

III.A.1.3. Kinetics of Oxidation of Si and SiGe

III.A.1.3.1. The Deal and Grove Model for Oxidation of Si

Oxidation of pure Si has been thoroughly investigated since the 1960's for various applications such as gate oxide formation, masks against implant or diffusion, surface passivation, dielectric isolation between adjacent devices, etc… The two following chemical reactions describe the thermal oxidation of pure Si when a silicon surface is exposed to an oxygen or a vapor water atmosphere:

ȃ "Dry" oxidation in O 2 ambient: Si O SiO s g s ( ) ( ) ( ) + → 2 2 ȃ "Wet" oxidation in H 2 O ambient: Si H O SiO H s g s g ( ) ( ) ( ) ( ) + → + 2 2 2 2 2
It has been shown that oxidation involves diffusion of the oxidizing species, either O 2 or H 2 O, through the thermal oxide to the Si-SiO 2 interface where the oxidation reaction occurs [Ligenza60]. Taking into account Si and SiO 2 densities and molecular weights, we can show that growth of a SiO 2 layer of thickness t ox corresponds to the "consumption" of a Si layer of thickness 0.44t ox .

In 1965, Deal and Grove published a seminal work on kinetics of Si thermal oxidation which is in good agreement with most published oxidation data [Deal65]. Their model describes both dry and wet oxidation chemical reactions for temperatures between 700°C and 1300°C, partial pressures between 0.2 atm and 1 atm, and oxide thicknesses in the 30 nm to 2000 nm range. The Deal and Grove model is based on the following three-step process and is illustrated Figure III.1: ȃ The oxidizing species is transported from the outer gas to the surface of the oxide. ȃ The oxidizing species diffuses from the surface of the oxide through the oxide layer and towards the oxide-Si interface. ȃ The oxidizing species reacts with Si atoms at the oxide-Si interface to form SiO 2 .

The equilibrium constant of this reaction is:

K a a a a x x SiO SiO Ge SiGe GeO GeO Si SiGe = × × = - 2 2 2 2 1 ,
where a i j is the activity of i in j. This reaction is favorable if:

∆ ∆ G G RT K x x RT G G G G R R o SiO G e G eO Si < + ( )< -       < - + - - ( ) 0 0 1 1 2 2 ln ln .
The previous equation holds at 1000 K if 1 322 10 19 -≥ × - x .

, meaning that unless the Si content is lower that 3 22 10 19 . × -, GeO 2 will be reduced by Si to form SiO 2 . 3.22.10 -19 being extremely small, oxidation of SiGe is conside- red thermodynamically Si-selective in the whole range of Ge concentration.

Kinetics of furnace oxidation of Si has received much attention since the 1960's because it is a crucial process of CMOS manufacturing. We thus start with a description of the pioneering work of Deal and Grove that laid the foundations of all studies of oxidation of Si and SiGe that followed. Its limitations are then emphasized, and revised models are reviewed to conclude with today's most accepted models. Finally, the specificities of oxidation of SiGe with respect to Si are discussed. C * , C 0 , C i are respectively the equilibrium concentration of the oxidizing species in the gas phase, the concentration of the oxidizing species at the outer surface of the oxide, and the concentration of the oxidizing species at the Si-SiO 2 interface.

Equation III.2 Equation III.3

The authors made the following assumptions in their model:

ȃ Steady-state regime. ȃ Growth in the direction perpendicular to the surface of the sample. ȃ Constant temperature and pressure in the oxidizing gas. ȃ They assumed that the flow of oxidizing species from the gas to the vicinity of the outer surface follows a law analogous to Newton's law of cooling:

F h C C 1 0 = - * (
). ȃ C * is related to the partial pressure of the oxidizing species in the gas by Henry's law: C Kp * = . Henry's law holds for ideal dilute solutions and in the absence of dissociation or reactions of the oxidizing species at the outer surface.

ȃ D eff if the diffusion coefficient of the oxidizing species in the oxide layer. It is assumed to be constant, and since steady-state regime is assumed, diffusion of the oxidizing species follows Fick's laws:

F D C C x eff i 2 0 0 = - .
ȃ The oxidation reaction between the oxidizing species and Si occurs at the abrupt oxide-Si interface only and the flow corresponding to this reaction is assumed to be governed by the first-order relation:

F kC i 3 =
where k is the reaction constant.

Considering these asumptions, the oxide thickness versus oxidation time is found to follow the following linear-parabolic law:

x t A t A B ( ) / / , 2 1 4 1 2 1 2 = + +       - τ where x t
( ) is the oxide thickness, t the oxidation time, τ accounts for the duration needed to form an initial 20 to 30 nm-thick oxide layer, and A and B are constants related to the physical parameters of the model. The authors examined two limiting forms of Equation III.5:

ȃ For short oxidation times: t A B  2 4 , x t B A t ( ) .  + ( ) τ
The growth regime follows a linear law, with B A being referred to as the linear rate constant, and is surface-reaction limited.

ȃ For long oxidation times:

t A B  2 4 , x t Bt ( ) 2  .
The general linear-parabolic law is reduced to a parabolic law, with a parabolic rate constant B. In this re- gime, the oxidation rate is limited by diffusion of the oxidizing species through the thermal oxide.

Equation III.4 Equation III.5 Equation III.6

The Deal and Grove model accurately describes wet oxidation for all oxide thicknesses and dry oxidation for oxides thicker than 30 nm. However, the main limitation of the Deal and Grove model appears to be dry oxidation for thin oxide formation (thickness below 30 nm). As illustrated Figure III.2, a faster oxidation rate compared with the Deal and Grove model is observed for oxides thinner than 30 nm. As today's needs are to control oxidation kinetics by means of dry oxidation for oxides thinner than 30 nm, revised models trying to accurately describe Si oxidation below 30 nm have been proposed and are introduced into the next section. [START_REF] Ghani | A 90nm high volume manufacturing logic technology featuring novel 45nm gate length strained silicon CMOS transistors[END_REF] have shown that no oxygen atom is incorporated into the bulk of the oxide and that the oxide grows at the Si-oxide interface. These experiments have thus clearly confirmed the Deal and Grove sequential oxidation process described in the former section. However, kinetics of these mechanisms, especially in the thin oxide regime, have been put into question. Attempts at explaining Si oxidation kinetics can be separated into two main groups, depending on the limiting mechanism of oxidation in the thin oxide regime, either the interfacial reaction of O 2 with Si, or more recently, diffusion of O 2 through the oxide.

In keeping with Deal and Grove, Massoud et al. assumed an interfacial reaction-limited oxidation in the thin oxide regime and discussed various mechanisms to explain the higher oxidation rate [Massoud85][Massoud85']. They will not be discussed here as they have been dismissed later.

The hypothesis according to which the interfacial reaction is the limiting mechanism for thin oxides has been questioned by various investigations [Doremus84][Doremus89] [START_REF] Bongiorno | Reaction of the Oxygen Molecule at the Si(100)-SiO 2 Interface During Silicon Oxidation[END_REF]]. Some experiments and first-principles calculations show that the interfacial reaction is almost instantaneous and thus cannot be a limiting mechanism of oxidation [Bongiorno04'][Tsetseris06][Akiyama08]. These studies support that even for thin oxides, the limiting mechanism of oxidation is diffusion of O 2 through the growing oxide. To our knwoledge, Doremus was among the first ones to claim that O 2 diffusion is always the limiting mechanism and that the linear regime experimentally observed by Deal and Grove for thin oxides could be explained by modulation of O 2 diffusivity by strain in the thermal oxide. Awaji et al. and Kosowsky et al. showed by X-Ray Reflectivity (XRR) experiments the presence of a higher density oxide layer about one nanometer thick at the Si-oxide interface [START_REF] Afanas' Ev | Electronic properties of Ge dangling bond centers at Si 1-x Ge x /SiO 2 interfaces[END_REF][Kosowsky97]. The existence of this higher density interfacial oxide layer was also supported by first-principles calculations of the structure of the Si-SiO 2 interface conducted by Bongiorno et al. [Bongiorno03']. The same authors also showed by calculations that O 2 diffusion through a higher density oxide is decreased. Two disruptive models of Si oxidation based on the results of these investigations have been published

[Watanabe06][Cui09].
As a conclusion, most recent investigations of Si dry furnace oxidation claim that the limiting mechanism is diffusion of O 2 through the growing oxide even for thin oxides. According to these studies, O 2 diffusion, and as a consequence oxidation kinetics, would be influenced by the oxide density.

While there is extensive data on dry furnace oxidation in the literature discussed above, there is a lack of data for RTO of Si. Other mechanisms, such as drift of ionic oxygen species, may have to be considered for RTO [Young87][Young87'][Kazor95]. While some models have been proposed for RTO of Si, none of them was widely

III.A.1.3.3. Kinetics of SiGe Oxidation

An enhancement of the wet oxidation rate has been observed for SiGe relative to Si [START_REF] Labrot | Low thermal budget for Si and SiGe surface preparation for FD-SOI technology[END_REF], while both faster [Spadafora05] and identical [Long12] dry oxidation rates for SiGe versus Si have been reported, illustrating an obvious lack of consensus. Moreover, an intrinsic difficulty that arises when studying the dynamics of oxidation of SiGe films is the evolution of the composition of the film during the oxidation process. Indeed, depending on the process parameters, Ge tends to accumulate at the oxidizing interface at various rates. Most authors only considered the initial composition of the film when reporting SiGe oxidation rates and not the Ge concentration at the oxidation interface. It makes the analysis of the influence of Ge on the oxidation rate difficult [Long12]. In addition, the range of oxide thicknesses varies widely from one study to another, so that comparisons between published experimental data is complicated.

Several mechanisms have been proposed to explain SiGe oxidation kinetics and can be separated into two main categories depending on the assumed limiting mechanism, either the interfacial reaction or diffusion of the oxidizing species: ȃ Interfacial reaction limited oxidation:

• The lower Si-Ge binding energy compared to the Si-Si binding energy would increase the interfacial reaction rate and therefore the oxidation rate [START_REF] Houghton | Strain relaxation kinetics in Si 1-x Ge x /Si heterostructures[END_REF]. Such a mechanism would induce a higher oxidation rate in both dry and wet oxidation conditions.

• GeO formation followed by immediate reduction by Si at the oxide-SiGe interface was proposed as a transient reaction occuring in parallel to direct Si oxidation [Prabkakaran93]. This catalytic reaction would increase the oxidation rate in both wet and dry oxidation environments.

• Oxidation of pure Si induces stress at the Si-SiO 2 interface because a unit volume of Si is converted into 2.25 volumes of SiO 2 . This stress is partially relieved by generation and injection of Si atoms in interstitial positions into the oxide and into the Si substrate. The majority of these interstitials are injected toward the oxide and is oxidized by incoming O 2 , while the minor part causes interstitials supersaturation in the Si film [START_REF] Dong | On the role and modeling of compressive strain in Si-Ge interdiffusion for SiGe heterostructures[END_REF].

It has been observed that interstitials supersaturation is suppressed for oxidation of SiGe . As for the causes of the lower interstitials injection toward the substrate in the case of oxidation of SiGe, LeGoues et al. proposed that the larger lattice parameter of SiGe accounts for a part of the stress generated at the oxidizing interface. As a result, the interstitials injection required to relieve the interfacial stress would be lower. It was also suggested that the presence of SiGe at the oxidizing interface changes the balance between interstitials injected toward the oxide and those injected toward the substrate [Na-politani05]. Indeed, interstitials have higher formation energies in SiGe than in Si, which results in a lower interstitials concentration in SiGe.

The model schematically illustrated Figure III.3 and based on this mechanism was proposed to explain kinetics of oxidation of SiGe [START_REF] Labrot | Low thermal budget for Si and SiGe surface preparation for FD-SOI technology[END_REF]. Due to the high oxidation rate under wet oxidation conditions, interstitials do not have time to diffuse away from the oxidizing interface and accumulate in a thin layer below the oxide. This accumulation would slow down the oxidation rate of pure Si compared to SiGe, which in fact results in an enhancement of the wet oxidation rate of SiGe with respect to Si. As for dry oxidation conditions, there is no interstitials accumulation at the oxidizing interface for oxidation of both Si and SiGe because interstitials have time to diffuse away in both cases. Indeed, the oxidation rate is much lower for dry oxidation conditions than for wet oxidation conditions. Assuming that the interfacial reaction is the limiting step, LeGoues et al. argued that the oxidation rate of SiGe should not be higher than the one of Si in dry oxidation conditions. accepted, especially because of the important dependency of RTO mechanisms on differences from one RTO apparatus to another (such as the lamp spectrum and the distance between the lamp and the wafer) [Borisenko97].

ȃ Diffusion limited oxidation:

• Shimura et al. proposed that the higher oxidation rate of SiGe in wet conditions is caused by the lower Si-Ge binding energy because the limiting mechanism is the interfacial reaction in wet conditions [Shimura10]. While in dry environment, there is no or hardly any difference between Si and SiGe according to the authors. Based on Watanabe et al. model [Watanabe06] of oxidation of Si, they proposed that O 2 diffusion through the thermal oxide layer is the rate limiting mechanism of oxidation of SiGe in dry ambient, whatever the oxide thickness is. And since they showed that the atomic structures of thermal oxides grown on Si and on SiGe are nearly identical, they claimed that the oxidation rate of SiGe should not be higher than the one of Si.

We emphasize that all proposed explanations, except the one of Shimura et al., assumed an interfacial reaction-limited oxidation in the thin oxide regime in dry atmosphere. Indeed, these explanations were proposed before the recent investigations about oxidation of Si reviewed in Section III. A.1.3.2. argued that the limiting mechanism is diffusion of O 2 through the oxide. In keeping with conclusions drawn in Section III. A.1.3.2., the higher oxidation rate of SiGe should arise from different O 2 diffusivities through the thermal oxide depending on the oxidized material, Si or SiGe. Finally, we also note that no specific model or explanation has been proposed for RTO of SiGe while it has been shown that mechanisms, and as a result kinetics, of furnace oxidation and RTO of Si may be different In summary, we showed that thermal oxidation of SiGe is thermodynamically Si-selective in the whole range of Ge concentration. We will study by X-Ray Spectroscopy (XPS) the composition of the oxides grown on SiGe in Section III. A.3. in order to check if thermal oxidation of SiGe actually keeps with thermodynamic predictions.

It is widely accepted that thermal oxidation of Si follows a three-step sequence:

ȃ The oxidizing species is transported from the outer gas to the surface of the oxide. ȃ The oxidizing species diffuses from the surface of the oxide through the oxide layer and towards the oxide-Si interface. ȃ The oxidizing species reacts with Si atoms at the oxide-Si interface to form SiO 2 .

III.A.1.4. Conclusion of Section III.A.1.

III.A.2. Experimental

III.A.2.1. Oxidation Processes and Characterization Techniques

Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 films were epitaxially grown on bulk Si substrates. Film thicknesses were about 20 nm. We mention that we observed slight differences of composition (less than 1% for the Ge concentration) and of film thicknesses depending on the moment when these layers were grown. It induces slight differences in terms of SIMS profiles as will be seen later in this manuscript, but without invalidating the results and the analysis methodology.

Prior to any oxidation process, a standard HF-SPM-SC1-HCl cleaning process was performed. It resulted in the formation of an approximately 1 nm-thick oxide on top of the SiGe film.

Dry furnace oxidation at 900°C, dry RTO at 900°C, 1000°C and 1100°C, and finally ISSG Oxidation at 850°C, 950°C and 1050°C were performed on SiGe and Si layers. From three to seven oxidation durations were chosen for each couple of process type and temperature. All oxidation durations refer to the duration of the oxidation plateau and were chosen to target oxide thicknesses in the 30-150 Å range, which are typical oxide thicknesses at the end of the condensation process.

The composition of grown oxides as a function of depth was firstly characterized on four representative samples by parallel Angle-Resolved X-Ray Spectroscopy (pAR-XPS). Ge oxide formation during SiGe oxidation was investigated on all samples by XPS data acquired on the in-line XPS equipment. While the in-line XPS tool cannot provide depth-resolved information, it is optimized for fast and easy XPS data acquisition, and is therefore suitable for basic characterization of such a high number of samples.

Oxide thicknesses were measured by Spectroscopic Ellipsometry (SE) between 240 nm and 780 nm using the SiGe optical constants published by Nolot et al. [Nolot14]. A three-layer model composed of the oxide and two SiGe layers was used for fitting spectra. Thicknesses of the three layers and index of refraction of SiGe layers were treated as unknowns while the index of refraction of the thermal oxide and of the Si substrate were fixed to known values. The oxide thicknesses were also measured by XRR to validate ellipsometry results. Indeed, XRR is much less model-dependent than ellipsometry for the measurement of oxide thickness on SiGe. XRR also benefits from a high contrast between SiO 2 and SiGe thanks to large density difference between these two materials. XRR is thus a robust and reference technique for measurement of the oxide thickness on SiGe. We chose to use a three-layer model for ellipsometry because it matches pretty well XRR results for all samples, and using an ellipsometry model with more SiGe layers does not bring any significant improvement. However, XRR is not suitable for measurement of oxide on Si because of the low density contrast between Si and SiO 2 [Diebold99]. While XRR is suitable for oxide on SiGe only, ellipsometry is suitable for oxide measurement on both Si and SiGe, and was thus chosen for consistency of results. The relative difference between ellipsometry and XRR for the oxide thickness was found to be always in the 3 to 4 Å range. The uncertainty on the measured oxide thickness was set to ± 2 Å. All uncertainties given in this chapter come from uncertainty calculations taking as input this oxide thickness uncertainty. The calculated uncertainties show how confident we can reasonably be about the results and the conclusions of this work.

SIMS was performed to measure the Ge concentration as a function of depth after oxidation in order to correlate oxidation kinetics with Ge redistribution.

Oxidation kinetics will also be discussed in view of the density of thermal oxides. To this end, the density of grown oxides was measured by XRR on the SEXTANTS Beamline at Soleil Synchrotron. Details of the experimental We showed that most recent investigations claim that the limiting mechanism is diffusion of O 2 through the growing thermal oxide. These studies also argue that O 2 diffusion, and as a result oxidation kinetics, is influenced by the oxide density.

As for SiGe, there is no consensus about kinetics of dry oxidation of SiGe with respect to Si. We highlighted that the fact that the Ge concentration in the SiGe film, and especially at the SiGe-oxide interface, evolves during oxidation makes analysis of SiGe oxidation kinetics more difficult. In addition, oxidation kinetics of both Si and SiGe by means of RTO has not received much attention while RTO is commonly used today. Kinetics of oxidation of SiGe by means of various processes (dry furnace oxidation, dry RTO and ISSG oxidation) will be covered in Section III.A.4. and Section III.A.5..

III.A.2.2. Characterization of Strain Conservation During Oxidation

The crystallinity and strain of SiGe films after epitaxy and after oxidation have been evaluated by X-Ray Diffraction (XRD). Figure III.4 shows the experimental and fitted rocking curves along the (004) direction of a Si 0.7 Ge 0.3 layer after epitaxy and after dry RTO at 1100°C for 40 seconds. Curves after RTO were vertically shifted for clarity considerations. After epitaxy, the rocking curve was fitted with a single SiGe layer, which shows that the epitaxial layer is monocrystalline and homogenous. After RTO, a three-layer model was used for SiGe. As will be shown in the following sections, the composition of the SiGe layer after RTO is not homogeneous anymore, which explains why the main diffraction peak is broader and why a three-layer model was needed for fitting. Interference fringes are still observed and prove that the SiGe layer is still a monocrystal and that the broadening of the peak is not due to a polycrystalline structure, but from composition non-homogeneity. However, XRD could not be performed on samples oxidized by means of dry furnace oxidation and ISSG oxidation.

III.A.2.3. Control of Within Wafer Uniformity of the Oxide Thickness

The oxide thickness was also measured along the wafer radius to ensure that the oxidation processes are uniform enough within a wafer. Figure III.5 shows the oxide thickness measured by ellipsometry along the wafer radius for two oxidized samples, one by means of dry RTO and one by means of ISSG oxidation. All samples were found to feature the same profile of oxide thickness along the wafer radius. It is found that the oxide thickness along the radius of the wafer is comprised in a few Å range. XRR measurement was also performed and found to agree well with ellipsometry measurements with a 3 Å to 4 Å offset. We chose to consider the oxide thickness measured at the center of the wafer by spectroscopic ellipsometry for the analysis of oxidation kinetics in the next sections. The presence of Ge oxide on XPS Ge3d spectra of SiGe layers oxidized by means of dry RTO has been noticed. We therefore sysmetically characterized the presence of Ge oxide during thermal oxidation of SiGe as a function of process parameters.

We performed pAR-XPS characterization of four SiGe samples oxidized by means of dry RTO to measure the depth distribution of each element, and especially Ge oxide. Figure III.6 shows the Ge3d spectrum as a function of the collection angle of a Si 0.7 Ge 0.3 layer oxidized by means of dry RTO at 900°C for 20 seconds. As explained in Section II.3.3., the contribution to XPS spectra of a buried atom decreases when the collection angle increases. This is the case of the Ge3d peak of metallic Ge atoms (bonded to either Si or Ge atoms). On the other hand, the area of the peak associated with Ge oxide is independent from the collection angle, which shows that Ge oxide is exclusively located in the very top part of the oxide.

Reconstruction of depth profiles of atomic concentrations was perfomed from the analysis of Ge3d, Si2p, O1s and C1s spectra. Figure III.7 shows the reconstructed chemical profile of the Si 0.7 Ge 0.3 layer oxidized by means of dry RTO at 900 seconds for 20 seconds. We firstly noticed the usual presence of carbon contamination on top of the sample. We also observe that the proportion of O atoms to Si atoms is roughly 2 to 1 in the oxide, as expected for stoichiometric SiO 2 . The depth resolution is degraded below a few nanometers so that it does not allow very accurate reconstruction of the chemical profile of the SiGe layer below the oxide. The metallic Ge observed in the oxide must not be taken as a piece of evidence of the presence of Ge in the oxide. However, qualitatively speaking, there is no or hardly any metallic Ge atom in the oxide while Ge oxide is exclusively located in the first nanometer of the oxide (referred as "Top of the oxide" on Figure III.7).

We performed the same analysis on three other samples oxidized by means of dry RTO: a Si 0.7 Ge 0.3 layer

III.A.3.2. Ge Oxide Formation During Dry RTO of SiGe

III.A.3.2.1. Depth Distribution of Ge Oxide Measured by pAR-XPS

oxidized at 900°C for 180 seconds, a Si 0.7 Ge 0.3 layer oxidized at 1000°C for 30 seconds and a Si 0.9 Ge 0.1 layer oxidized at 1000°C for 60 seconds. The same conclusion was reached: a small quantity of Ge oxide (a few percent) is located in the very top of the oxide (the top 10 Å). We believe that these four samples are representative enough to extend this conclusion to all SiGe samples of this study oxidized by means of dry RTO. For all investigated samples, we followed the fitting procedure described in Section II.3.3. which uses two peaks for the Ge oxide. We found that the main peak is always centered at a binding energy about 4.25 eV above the one of metallic Ge. Such a chemical shift is associated with the fourth oxidation state of Ge [Prabhakaran00]. Moreover, we also noticed the disappearance of the peak associated with Ge oxide on Ge3d spectra after the sample surface has been rinsed with water (Figure not shown). As GeO 2 is known to be water soluble, it confirms that the Ge oxide observed after dry RTO of SiGe is mainly composed of GeO 2 and that it is located in the very top part of the oxide.

As oxidation proceeds by formation of oxide at the SiGe-oxide interface, the top of the oxide corresponds to the first stages of oxidation. Ge oxide is therefore likely formed at the beginning of the oxidation plateau. In order to precisely determine at which step(s) of the SiGe epitaxy-cleaning process-dry RTO sequence Ge oxide is formed, we performed XPS measurements after each key step of the sequence. XPS measurements were performed after SiGe epitaxy, after the cleaning process prior to oxidation, after the temperature ramp up until 550°C, after the temperature ramp up until 900°C, and after a complete dry RTO process. Figure III.8 (a) shows on a typical temperature-time profile of a dry RTO process at which steps measurements were performed.

Figure III.8 (b) shows Ge3d spectra after each of these steps. No Ge oxide is noticed after SiGe epitaxy and after the cleaning process prior to oxidation while some Ge oxide is observed right after the temperature ramp up until 550°C. The O1s peak is also noticeable after the temperature ramp up until 550°C, which shows that the oxide started to grow even at temperatures below 550°C. The O1s peak area keeps increasing as the temperature goes up and oxidation at the target temperature proceeds. There is still some Ge oxide observed after the temperature ramp up until 900°C and after the complete oxidation process. Comparison of Ge3d 4+ peak areas after each of these process steps show that there are more Ge oxide after the temperature ramp up until 900°C than after the temperature ramp up until 550°C only (Figure III.11 shown and discussed into more details in Section III. A.3.2.4.). Within measurement and fitting errors, no clear difference between peak areas has been evidenced between the temperature ramp up until 900°C and the complete oxidation process.

As a conclusion, Ge oxide is formed during the temperature ramp up until at least 550°C. Some Ge oxide is still formed above 550°C but we cannot precisely determine at which moment they are not formed anymore based on these results only. Figure III.9 allows us to study the influence of the initial Ge concentration on the quantity of Ge oxide formed during dry RTO of SiGe. The metallic Ge peak area depends on various parameters, especially the thickness on top of the SiGe layer because of the depth sensitivity of the XPS technique (see Section II.3.3.) and of the pile-up mechanism which leads to various SiGe composition profiles. Analysis of the area of the metallic Ge peak is thus complex while the one of the Ge3d 4+ and O2s is easier because they both directly depend on the total amount of GeO 2 and SiO 2 in the thermal oxide. We chose to compare samples that have close oxide thicknesses. For all investigated temperatures, the Ge oxide peak area is larger for Si 0.7 Ge 0.3 layers than for Si 0.9 Ge 0.1 layers. All other parameters being equals, there is more Ge oxide after oxidation of SiGe when the initial Ge concentration is higher. As for the effect of the temperature of dry RTO on the quantity of Ge oxide, we observe on Figure III.10 that the area of the Ge3d 4+ peak is slightly higher when the oxidation temperature is lower. Figure III.11 shows the area of the fitted Ge3d 4+ peak after dry RTO of SiGe as a function of the initial Ge concentration and of the oxidation temperature. There are several points for a single temperature and a single initial Ge concentration because measurements were performed for several oxidation durations. No clear evidence of variation of the quantity of Ge3d 4+ peak area with the oxidation duration was found. But due to different Ge3d metallic and O2s peak shapes at different oxide thicknesses, precise analysis of the evolution of the Ge3d 4+ peak area 99/230 Chapter III -SiGe Thermal Oxidation with the oxidation time is less accurate.

III.A.3.2.2. Determination of the Oxidation State of Ge Oxide

III.A.3.2.3. Ge Oxide Formation During the Full Oxidation Sequence

III.A.3.2.4. Influence of Initial Ge Concentration and Oxidation Temperature on Ge Oxide Formation

The measured peak areas after the temperature ramp up until 550°C and 900°C for the case of Si 0.7 Ge 0.3 are also shown as full lines over the total range of temperature for comparison. As already mentionned, Ge oxide is formed below and above 550°C, so that the peak area is higher after a temperature ramp up until 900°C than after a temperature ramp up until 550°C only. Figure III.11 also confirms that the higher the initial Ge concentration is, the higher the quantity of Ge oxide is. More precisely, the quantity of Ge oxide for dry RTO of Si 0.7 Ge 0.3 seems to be roughly three times as high as the one for Si 0.9 Ge 0.1 . The Ge oxide quantity would therefore be proportional to the initial Ge concentration of the SiGe layer. Thirdly, the Ge oxide quantity seems to decrease when the oxidation temperature increases, but the variations are lower than the one caused by the initial Ge concentration. . Such mixed oxide formation occurs when using low temperature and/or high oxidant pressure oxidation conditions. Paine et al. also showed that the mixed oxide formed under these oxidation conditions is mostly composed of SiO 2 -GeO 2 , with a GeO 2 to SiO 2 fraction roughly similar to the Ge to Si fraction in the SiGe layer at the oxidation interface. They argued that when using these oxidation conditions, the interdiffusion rate of Si towards the oxidation interface is much lower than the oxidation rate, so that we can almost consider that diffusion is suppressed. The huge gap between the oxidation rate and the diffusion rate is responsible for the mixed oxide formation.

We showed in Section III.A.3.2.1. that Ge oxide is always located in the top part of the oxide (the top 10 Å) (see Figure III.7). Based on XPS results, no Ge oxide was observed after either SiGe epitaxy and after the standard chemical cleaning process used prior to any oxidation process (see Figure III.8). There is thus no Ge oxide just before thermal oxidation of SiGe. As the top of the oxide layer corresponds to the beginning of oxidation, we argued that Ge oxide is formed during the first stages of oxidation only.

In Section III.A.3.2.2., we showed that in all samples where the presence of Ge oxide is observed by XPS, there is a major contribution of the Ge3d 4+ peak, which makes us think that formed Ge oxide is mainly GeO 2 , in agreement with Paine et al.. We suggest that during the temperature ramp up of dry RTO, the atmosphere being already pure O 2 at 740 Torr, some oxidation occurs even below 550°C and leads to mixed SiGeO x formation. The formation of SiGeO x seems to continue between 550°C and 900°C as the Ge3d 4+ peak area increases between 550°C and 900°C (see

III.A.3.2.5. Discussion about Ge Oxide Formation During Dry RTO of SiGe

Figure III.11). We argue that, during temperature ramp up of dry RTO, the oxidation temperature is low enough and the oxidation rate high enough (the oxide is very thin, which results in a high oxidation rate as will be seen in Section III. A.4.1.) for the oxidation regime to be the one described by Paine et al. and Kilpatrick et al. that corresponds to mixed oxide formation. We also observed that the Ge3d 4+ peak area after temperature ramp up until 900°C is close to the ones of samples oxidized at 900°C for various oxidation durations, while the Ged 4+ peak areas of samples oxidized at higher temperatures are lower (see Figure III.11). It shows that there is no Ge oxide formation after the temperature ramp up.

In summary, we claim that Ge oxide is formed during the temperature ramp up, mostly until 900°C, and that the quantity of formed Ge oxide is be roughly proportional to the initial Ge concentration, in agreement with Paine et al.. Indeed, the GeO 2 quantity for dry RTO of Si 0.7 Ge 0.3 seems to be roughly three times as high as the one for Si 0.9 Ge 0.1 (see Figure III.11). One possible mechanism for the smaller GeO 2 quantity at higher temperatures could be recombination of GeO 2 into GeO during dry RTO followed by desorption of GeO [Prabhakaran00].

III.A.3.3. Ge Oxide Formation during Dry Furnace Oxidation and ISSG Oxidation of SiGe

Figure III.12 (a) and (b) show that there are no Ge oxide after either the temperature ramp up until 900°C or after complete oxidation for the case of dry furnace oxidation. As described in Section II.2.3.2., the atmosphere is pure N 2 until 600°C, so that there is no oxidation during the temperature ramp up until 600°C. Then, the partial pressure of O 2 gradually increases from 600°C to 900°C and oxidation starts. However, because the O 2 partial pressure is increasing into the chamber during the temperature ramp up, oxidation is most likely less important. We argue that there is no Ge oxide formation during the temperature ramp up because the oxidation rate is not high enough compared with the diffusion rate for the mixed oxide formation conditions to be fulfilled. Some Ge oxide contribution on Ge3d spectra is noticed for the temperature ramp up until 900°C of ISSG oxidation and after complete ISSG oxidation at 850°C (see Figure III.12 (a) and (c)). For ISSG oxidation, the chamber is filled with pure O 2 at low pressure until 600°C is reached, and H 2 starts to flow into the chamber when the temperature exceeds 600°C. Some SiGeO x is also formed during this temperature ramp up, but the Ge oxide quantity seems to be smaller than for the temperature ramp up of dry RTO (see Figure III.12 (a)). Further investigation is required to identify the contributions of the temperature ramp up below and above 600°C, which correspond to different oxidation atmospheres (see Section II.2.3.2.), to Ge oxide formation. Similarly to dry RTO, one possible mechanism for the smaller GeO 2 quantity at higher temperatures could be recombination of GeO 2 with other species of the oxide and of the atmopshere (such as O* which has a strong oxidizing power) to form GeO which then desorbs [Prabhakaran00]. We found that in spite of the thermodynamic tendency to form SiO 2 during oxidation of SiGe, mixed SiGeO x may be formed due to kinetic factors. We investigated the formation of Ge oxide by pAR-XPS and XPS. We observed that mixed SiGeO x are formed when the oxidation rate largely exceeds the Ge diffusion rate. These mixed oxide formation conditions are fulfilled during the temperature ramp up of dry RTO under one atmosphere of O 2 . The quantity of formed Ge oxide, found to be mainly GeO 2 , is roughly proportional to the Ge concentration of the SiGe epitaxial layer.

We also claimed that the oxidation rate during the temperature ramp up of dry furnace oxidation is not high enough for Ge oxide to be formed. As for ISSG oxidation, some GeO 2 are formed during the temperature ramp up but disappear afterwards if the oxidation temperature is above 950°C. More experiments are required to clarify GeO 2 formation and disappearance during ISSG oxidation of SiGe.

To conclude, we showed that, except possibly during the very first stages of oxidation, the composition of oxides grown on SiGe is pure SiO 2 for all investigated oxidation processes, in agreement with thermodynamic considerations of Section III.A.1.2.. Having studied the composition of the grown oxide, the next section focuses on the kinetics of oxidation of SiGe.

III.A.3.4. Conclusion of Section III.A.3.

III.A.4. SiGe Oxidation Kinetics and SiGe Composition Evolution III.A.4.1. Oxidation Kinetics Results

III.A.4.1.1. Results

Figure III.13 shows the measured oxide thickness as a function of time for dry furnace oxidation, dry RTO and ISSG oxidation. We observe that, for all investigated oxidation processes, and for identical oxidation temperature and duration, the higher the initial Ge concentration is, the higher the oxide thickness is. However, the differences between the oxide thicknesses for SiGe and for Si are much lower for ISSG oxidation than for dry furnace oxidation and dry RTO.

Figure III.13: (a) Oxide thickness as a function of the oxidation time for dry furnace oxidation of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 at 900°C. (b) (c) and (d) show the oxide thickness as a function of the oxidation time for dry RTO of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 at 900°C, 1000°C and 1100°C respectively. (e) (f) and (g) show the oxide thickness as a function of the oxidation time for ISSG oxidation of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 at 850°C, 950°C and 1050°C respectively. Dots are experimental points and dotted lines are guides for the eye.

The goal of this section is to identify the oxidation mechanisms that govern oxidation kinetics and to correlate oxidation kinetics with the process parameters and with the evolution of the Ge concentration during oxidation. Experimental oxidation rates of dry furnace oxidation, dry RTO and ISSG oxidation of SiGe are firstly compared with the Deal and Grove linear-parabolic regime. A quantitative analysis methodology based on the definition of a term called Growth Rate Enhancement (GRE) is then introduced. After having shown that SiGe oxidation kinetics does not depend on the Ge concentration of the initial SiGe epitaxial film only, we study the evolution of the Ge concentration profile during oxidation as a function of process parameters. The last part of this section shows that oxidation kinetics and the Ge concentration at the oxidation interface are correlated.

We also highlight that the oxidation rate decreases with time during oxidation for all samples. The oxidation regime is thus not linear, in contradiction with the Deal and Grove model for thin oxides (thinner than 20 nm) [Deal65]. However, to highlight in a more quantitative manner the discrepancy with the Deal and Grove model, we tried to fit our experimental data with the Deal and Grove model with two methods: ȃ If we fix the parabolic constant to the value provided by the Deal and Grove model (the parabolic constant can be determined from oxidation kinetics for oxides much thicker than 200 Å according to the Deal and Grove model, which is not the case in our study), the fit quality is highly degraded as illustrated Figure III.14. ȃ If we let free all parameters of a linear-parabolic model to fit our experimental data, the fitting procedure tends to a purely parabolic regime. There is thus hardly any linear contribution, even for short oxidation times.

The Deal and Grove model is thus not adequate for fitting the experimental data. As shown Figure III.15, all oxidation processes follow parabolic regimes with the following equation type:

x t 2 = + α β,
where x is the oxide thickness and α and β are coefficients that depend on the temperature, the initial Ge concentration and the oxidation process type. In this section, we thus showed that firstly, the oxidation regime is parabolic in the whole range of oxide thicknesses. According to the Deal and Grove model, oxidation parabolic regimes refer to an oxidation rate limited by diffusion of the oxidizing species through the thermal oxide [Deal65]. Secondly, the oxidation rate is higher for a higher initial Ge concentration for the three studied oxidation processes.

III.A.4.1.2. Discussion about Temperature Calibration

Prior to any attempt to analyze and discuss differences of oxidation kinetics between Si and SiGe, one question that needs to be answered is that if the observed differences of oxidation kinetics do not arise from different actual oxidation temperatures between Si and SiGe samples.

For furnace oxidation, the equipment used was a batch equipment with a good spatial temperature homogeneity. Since Si and SiGe samples were oxidized in the same oxidation run, the possibility of different temperatures between Si and SiGe samples is thus ruled out.

For dry RTO and ISSG oxidation, we highlight that the measured temperature is the one of the back-side of the wafer, which is identical for Si and SiGe samples. But let us assume that the observed differences in oxidation kinetics are actually a temperature effect, in other words, that the parabolic constant α is only dependent on the oxidation process type and the temperature, and not on the Ge concentration. Knowing the activation energies of the parabolic constants α of dry RTO and ISSG oxidation of Si (they will be discussed in Section III. A.4.1.4.), we can extract the temperature that would correspond to the observed oxidation kinetics of SiGe.

There would be approximately 25°C and 50°C gaps between the set oxidation temperature and the actual one for dry RTO of Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 respectively. Such differences are much higher than possible temperature calibration errors. As for ISSG oxidation, gaps are much lower, between 3°C and 7°C. ISSG oxidations were performed in a similar equipment as dry RTO and in a single run. We are thus confident that the observed differences of oxidation kinetics for ISSG oxidation of SiGe arise from truly different oxidation kinetics between Si and SiGe and not from temperature calibration errors. Therefore, we believe that possible temperature errors have a minor contribution in any oxidation process of this study, and that the differences of oxidation rates arise from truly Ge-based and oxidation process type effects.

III.A.4.1.3. Kinetics of Dry Oxidation of Si

As illustrated Figure III.16, there is hardly any difference between oxidation kinetics of Si oxidized at 900°C by means of dry furnace oxidation and dry RTO. The slightly systematic higher oxide thickness for furnace oxidation compared to dry RTO could arise from the contribution of oxidation during temperature ramps, which are longer for furnace oxidation than for dry RTO. These results suggest that O 2 diffusion properties through the thermal oxide grown on Si are identical for both processes. Deal and Grove found that the activation energy of the parabolic constant is 1.24 eV and assumed that it corresponds to the one of diffusion of O 2 through the thermal oxide, in good agreement with the one reported by Norton (1.17 eV) [START_REF] Oh | Chemical structure of the ultrathin SiO 2 /Si(100) interface: An angle-resolved Si2p photoemission study[END_REF]. The activation energy of the parabolic constant of dry RTO of Si was found to be 2.29 eV in our study, in disagreement with the ones reported by Deal and Grove and Norton for dry furnace oxidation. However, it is in better agreement with the ones reported by Kazor for dry RTO of Si (2.17 eV) and by David et al. for dry RTO of SiGe (2.1 eV) [Kazor95][David16]. As suggested by Kazor, mechanisms specific to one of the two oxidation processes likely take place and affect activation energies of oxidation. The activation energy of dry RTO found in our study is also in fairly good agreement with the activation energy of diffusion of O 2 in strained thermal SiO 2 calculated by Bongiorno et al., Tsetseris et al. and Akiyama et al. (about 2.0 eV) . In keeping with these observations, kinetics of dry RTO of Si would be limited by diffusion of O 2 through the thermal oxide and the activation energy of diffusion would be about 2.3 eV. We also note that the activation energy of dry RTO of SiGe decreases when the initial Ge concentration increases.

As for ISSG oxidation, according to literature, ISSG oxidation is mainly limited by diffusion of O* radicals through the growing oxide, which results in a parabolic regime [Kee00]. The extracted activation energy of the parabolic constant of ISSG oxidation is therefore the one of diffusion of O* in SiO 2 . We observe that ISSG oxidation follows parabolic regimes with activation energies around 1 eV for all initial Ge concentrations.

Si

Si 0.9 Ge To go further into the analysis of oxidation kinetics of SiGe, the next section will quantitatively compare the oxidation rate of SiGe with the one of Si.

III.A.4.2. Growth Rate Enhancement (GRE)

III.A.4.2.1. Analysis Methodology of Oxidation Kinetics

In order to quantitatively examine oxidation kinetics of SiGe with respect to Si, we need to define a meaningful term that compares oxidation rates of SiGe and Si. We could not find such a definition in literature. We define a Growth Rate Enhancement (GRE): the GRE is the ratio of the SiO 2 growth rate for the case of oxidation of SiGe to the one for Si at identical oxide thickness. As a consequence of the observed parabolic regime, the GRE does not depend on the oxide thickness and has a unique value for a unique combination of oxidation process type, initial Ge concentration and oxidation temperature. Finally, the GRE is reduced to:

GRE SiGe initial concentration oxidation process type T = ( ) α , , α α Si oxidation process type T , , , (
)
where SiGe initial concentration is either 10% or 30% Ge concentration in this study, the oxidation process type is either dry furnace oxidation, dry RTO, or ISSG oxidation, T is the oxidation temperature, and α the coefficients extracted from the fitted parabolic laws of We propose this GRE definition for the following reasons. Firstly, in order to quantitatively examine the effect of Ge on oxidation kinetics, what must be considered is the oxidation rate and not directly the oxide thickness. Indeed, whatever the mechanisms that govern oxidation kinetics are, oxidation models such as the ones of Deal and Grove, Watanabe et al. and Cui et al. are based on oxidizing species flux and oxidation rates [Deal65][Watanabe06] [Cui09]. Secondly, because the oxidation regime is parabolic, the oxidation rate tends to decrease with time. With this GRE definition, the contribution of the increasing oxide thickness to the oxidation rate is thus excluded and only the effects of both the oxidation process type and Ge on oxidation kinetics are analyzed.

III.A.4.2.2. GRE Dependence on the Initial Ge Concentration

As seen previously, the initial Ge concentration influences the oxidation rate of SiGe. Figure III.19 shows the GRE as a function of the initial Ge concentration. We observe that, for either dry furnace oxidation, dry RTO or ISSG oxidation, the GRE is always higher for Si 0.7 Ge 0.3 than for Si 0.9 Ge 0.1 (see (1) on Figure III.19). We also observe that the GRE is higher for dry RTO than for dry furnace oxidation (see (2) on Figure III.19). Therefore, for a given oxidizing species, how the thermal energy is provided influences oxidation kinetics of SiGe.

If we now compare the GRE values of dry RTO and ISSG oxidation, two RTO process types, the GRE is much lower for ISSG than for dry RTO. The nature of the oxidizing species thus also influences the GRE. Then, for both dry RTO and ISSG oxidation, the GRE depends not only on the initial Ge concentration, but also on the oxidation temperature (see (3) on Figure III.19). However, the GRE of dry RTO and of ISSG oxidation depend on the oxidation temperature in the opposite way. The higher the oxidation temperature is for a fixed initial Ge concentration, the lower the GRE is for dry RTO. On the contrary, the higher the oxidation temperature is, the higher the GRE of ISSG oxidation is.

Therefore, further investigation is required to understand how the nature of the oxidizing species, the initial Ge concentration and the temperature influence oxidation kinetics of SiGe. As described in Section III. A.1.3.3., the SiGe composition depth profile evolves during oxidation. We hence propose to consider into the next section the Ge redistribution during oxidation of SiGe to further examine the GRE. .20 shows the SIMS profiles of Ge concentration after oxidation of various samples oxidized by means of dry RTO. For clarity considerations, the x axis is shifted so that the SiGe-oxide interfaces are all close to zero. We observe that the Ge concentration just below the oxide, so that at the oxidation interface, is always higher than the initial Ge concentration of the SiGe epitaxial layer. The pile-up mechanism caused by the Si-selective oxidation of SiGe is therefore clearly evidenced. The Ge concentration profile after oxidation, and especially at the oxidation interface, obviously depends on the initial Ge concentration and oxidation temperature.

For dry RTO at 900°C, a few nanometers-thick layer with a high Ge concentration is observed below the oxidation interface. Plateaux at 10% and 30% Ge concentration corresponding to the initial SiGe epitaxial layer and a quite abrupt interface between the SiGe epitaxial layer and the Si substrate are still clearly visible on Figure III.20 (a) and (b). While for dry RTO at 1000°C and 1100°C, no layer with abrupt interfaces and formed by the pile-up mechanism can be identified on Figure III.20 (c),(d) and (e). The plateaux at 10% or 30% Ge concentration have also disappeared and interfaces between the SiGe epitaxial layer and the Si substrate are now much graded. The higher the oxidation temperature is, the less noticeable the pile-up mechanism is. As a consequence, the Ge concentration at the oxidation interface is lower at higher oxidation temperatures.

With the oxidation conditions used in this study, all Ge atoms at the oxidation interface are piled-up below the oxidation front in the underlying SiGe crystal [Paine91] (see Section IV.A.3.). Stated differently, there is no Ge segregation in the oxide. The Ge concentration profile below the oxide, and especially at the oxidation interface, depends on a competition between oxidation that piles-up Ge atoms below the oxidation front and SiGe interdiffusion that tends to smooth the Ge concentration profile below the oxide. Oxidation and diffusion in SiGe are both thermally activated mechanisms that follow Arrhenius-type laws. As shown in Section III. A.4.1.4., the activation energy of dry RTO of SiGe is about 2 eV. While the ones of Ge and Si diffusion in SiGe are in the 4 eV to 5 eV range [Kube10]. Phrased in qualitative terms, the temperature dependence of diffusion is much higher than the one of both dry RTO and ISSG oxidation. Therefore, the diffusion rate increases faster with temperature than the oxidation rate does, which makes the Ge pile-up less pronounced as Ge atoms diffuse away.

The second key point of this section is that the Ge concentration at the oxidation interface seems to be constant after a transient regime. For dry RTO of Si 0.7 Ge 0.3 at 1100°C (see Figure III.20 (c)), the first ten seconds at least likely correspond to a transient regime when the oxidation rate is particularly high because of the very thin oxide. In that case, the oxidation rate largely exceeds the diffusion rate. Therefore, we will not consider this transient regime for the analysis of the GRE, and will assume that the Ge concentration at the oxidation interface is rather constant at about 33% during almost all the investigated oxidation durations range. For dry RTO of Si 0.7 Ge 0.3 at 1000°C (see Figure III.20 (b)), the Ge concentration at the oxidation interface is 41.7% and 42.4% after 60 seconds and 120 seconds of oxidation respectively. As for dry RTO of Si 0.7 Ge 0.3 at 900°C (see Figure III.20 (a)), the Ge concentration at the oxidation interface increases from 45.9% after 120 seconds of oxidation to 49.8% after 1080 seconds of oxidation. Nevertheless, this 3.9% variation of the Ge concentration with time is rather low compared to the gap induced by the oxidation temperature, from 33% Ge concentration at the oxidation interface at 1100°C to more than 46% at 900°C. (c) 1100°C, and of (d) Si 0.9 Ge 0.1 at 900°C and (e) at 1100°C. For clarity considerations, the x axis is shifted so that the SiGe-oxide interfaces are all close to zero.

III.A.4.3. Evolution of the Ge Concentration at the Oxidation Interface

The Ge concentration at the oxidation interface after dry RTO of Si 0.9 Ge 0.1 at 1100°C for 50 seconds and at 900°C for 540 seconds are 16% and 21% respectively. By analogy with the case of Si 0.7 Ge 0.3 , we can reasonably assume that there is also a plateau for the Ge concentration at the interface for dry RTO of Si 0.9 Ge 0.1 , taken as 16% at 1100°C and 21% at 900°C (see Figure III.20 (d) and (e)).

The evolution of the Ge profile below the oxide, and especially at the oxidation interface, depends on the interdependence of the pile-up mechanism caused by the Si-selective oxidation and SiGe interdiffusion. This manuscript does not intend to quantitatively explain the observed quasi-plateau of the Ge concentration at the oxidation interface during dry RTO, but rather keeps to the observed tendencies to discuss oxidation kinetics.

To summarize the main conclusions of this section, we firstly claim that for all investigated dry RTO processes, a pile-up layer is rapidly formed below the oxide, and that after a transient regime, the Ge concentration at the oxidation interface is quite constant with time over the range of studied oxidation durations. We also observe that the higher the oxidation temperature is, the lower the Ge concentration at the oxidation interface is.

III.A.4.4. GRE Dependence on the Ge Concentration at the Oxidation Interface

Following Section III. A.4.3. conclusions,Figure III.22 shows the GRE as a function of the Ge concentration at the oxidation interface measured by SIMS, except for dry RTO of Si 0.7 Ge 0.3 at 1100°C for 10 seconds that likely correspond to a transient regime. We found that the higher the Ge concentration at the oxidation interface is, the higher the GRE for dry RTO is. For dry RTO, the GRE dependence on the initial Ge concentration and on the temperature shown Figure III.19 (a) can be summed up by its dependence on the Ge concentration at the oxidation interface. Due to the GRE definition given Equation III.9, the oxidation rate of SiGe by means of dry RTO is therefore the one of Si oxidized in the same conditions multiplied by a factor that depends only on the Ge concentration at the interface.

As already shown Figure III.19 (a), the GRE of dry furnace oxidation at 900°C is lower than the one of dry RTO at the same temperature. It was also evidenced that the Ge concentrations at the oxidation interface for dry As shown Figure III.21 (a), with a Ge concentration at the oxidation interface of 45.5% for dry furnace oxidation at 900°C of Si 0.7 Ge 0.3 , there is hardly any difference with dry RTO. The differences in terms of total quantities of Ge atoms in both layers arise from a thinner epitaxial SiGe film for the case of dry furnace oxidation. As for ISSG oxidation, the same tendency as for dry RTO is obserbed, with a strong pile-up layer with a maximum Ge concentration of 59.7% observed for ISSG oxidation of Si 0.7 Ge 0.3 at 850°C. Indeed, ISSG oxidation at 850°C is much faster than dry RTO, which increases even more the gap between the oxidation rate that piles up Ge atoms and the diffusion rate. The Ge concentration at the oxidation interface for ISSG oxidation of Si 0.9 Ge 0.1 at 1050°C is 22.3%.

To conclude Section III.A.4., we showed that oxidation of SiGe by means of either dry furnace oxidation, dry RTO and ISSG oxidation is limited by diffusion of the oxidizing species through the growing oxide. The oxidation rate of SiGe was found to be weakly dependent on the presence of Ge for ISSG oxidation, while it is significantly higher for the cases of dry furnace oxidation and dry RTO. A factor called Growth Rate Enhancement (GRE) was then defined to quantitatively compare the oxidation rates of Si and SiGe.

We also evidenced that a Ge pilled-up region below the oxide is rapidly formed during oxidation and that afterwards, the Ge concentration at the oxidation interface is rather constant with time over the range of studied oxidation durations. For dry RTO, our results show that the oxidation rate of SiGe increases with the Ge concentration at the oxidation interface. It was shown that the gap between the GRE of SiGe oxidized by means of dry RTO with the one for dry furnace oxidation cannot be explained based solely on the comparison of the Ge concentrations at the oxidation interface, which was found to be almost equals for both processes. The process type thus influences oxidation kinetics of SiGe.

Various studies showed or suggested that the oxide density can influence diffusion of O 2 through the oxide and that this density is influenced by oxidation process parameters (see Section III. A.1.3.2.). The next section thus discusses further oxidation kinetics of SiGe in light of the measurement of the oxide density by X-Ray Reflectivity (XRR). furnace oxidation and dry RTO are similar. It thus shows that the Ge concentration at the oxidation interface is not the only parameter that affects oxidation kinetics of SiGe for dry oxidation, and that the process type also influences oxidation kinetics of SiGe.

III.A.4.5. Conclusion of Section

Finally, kinetics of ISSG oxidation is only slightly dependent on the Ge concentration at the oxidation interface, as already shown Figure III.19 (b).

III.A.5. Correlation Between SiGe Oxidation Kinetics and Oxide Density

We showed in the previous section that oxidation of SiGe is governed by diffusion of the oxidizing species through the growing thermal oxide. As the density of the oxide is known to influence diffusion of the oxidizing species, this section aims to understand how oxidation kinetics of SiGe depends on the oxide density, and how the process parameters and the Ge concentration at the oxidation interface affect the oxide density. This section starts by describing the methodology of measurement of the oxide density by X-Ray Reflectivity (XRR). Then, we investigate how the process parameters and the Ge concentration at the oxidation interface influence the oxide density. Oxidation kinetics is correlated with the oxide density and discussed in light of the influence of the oxide density on the diffusivity of O 2 through the thermal oxide. Finally, the dependence of O 2

III.A.5.1. Measurement of the Oxide Density by X-Ray Reflectivity (XRR)

III.A.5.1.1. Methodology for Extraction of Scattering Factors from Absorption Spectra

We intended to measure the density of thermal oxides grown on SiGe by Resonant Soft X-Ray Reflectivity (R-SoXR) around the O K-absorption edge. Indeed, R-SoXR can be a powerful technique for measurement of small variations of electron densities (see Section II.4.1.4.). As drawn in Section II.4.1.2., scattering factors of all elements in the sample must be known for analysis of XRR spectra, hence the description of their measurement in this section.

While the tabulated values of the scattering factors do not accurately reproduce the precise variations of the scattering factors close to an absorption edge, they are reliable far from it. Si and Ge scattering factors below and above the O K-absorption edge used in the fitting procedure are thus tabulated values [CXRO].

The following procedure was used to measure O scattering factors around its K-absorption edge: As a result:

ȃ
a f eV f eV I eV I eV = ( )-( ) ( )-( ) 2

( ( )

.

We also highlight that Si and Ge atoms absorb X-Ray photons, and thus contribute to the measured absorption spectra. However, the variations of their scattering factors between 520 eV and 550 eV are much lower than the one of O atoms. Therefore, their contributions in the measured absorption spectra has no effect on calculation of the imaginary part of the O scattering factor using Equations III.10.

Equation III.9

Equations III.11 diffusion on the oxide density, a subject that the study of oxidation kinetics allows to investigate, is discussed. Hardly any difference can be observed between O scattering factors of the three samples for the T.E.Y. mode (see Figure III.23 (a)). The T.E.Y. mode shows that the O scattering factor of the top of the oxide is not or hardly influenced by Ge and by the oxidation temperature. A hollow can be observed at around 528 eV while no noticeable variation is expected below the absorption edge. This hollow is likely due to absorption caused by residual oxygen in the beamline (absorption of the beamline was measured independently from these experiments and shows some absorption at around 528 eV).

Equations III.10

III.A.5.1.2. O Scattering Factor around the K-Absorption Edge

O scattering factors acquired in T.F.Y. mode are shown Figure III.23 (b). The O scattering factor of the Si sample oxidized at 1100°C and measured in T.E.Y. mode is also plotted for comparison. The same hollow is observed at about 528 eV, which supports the suggestion of absorption by residual oxygen in the beamline. In T.F.Y. mode only, a pre-edge "bump" is observed at around 532 eV. This pre-edge "bump" can thus be associated with O atoms close to the interface with Si(Ge). We also observe that the intensities of this pre-edge "bump" and of the maximum of absorption at 535 eV are different between the three samples. It is noticed that the O scattering factor of the Si sample in T.F.Y. mode unexpectedly decreases from about 1.6 at 520 eV to 0.3 at 528 eV. We believe that this is due to measurement issues and that this decrease has no physical meaning. Muller et al. found that the scattering factor of O atoms located at the Si-oxide interface is very different from the one of O atoms in the bulk of the oxide. More precisely, they found that the edge onset is reduced by 3 eV and that there is no sharp peak at 535 eV. They claimed that this sharp peak is associated with O-O scattering and is strongly reduced when the number of O second-nearest neighbours around the excited O atom is reduced. The number of O second-nearest atoms being reduced in the vicinity of the Si-oxide interface, the sharp peak tends to disappear close to the interface.

The energy shift of the pre-edge "bump" with respect to the edge onset of the T.E.Y. mode is lower than but in qualitative agreement with the one observed by Muller et al.. These results support that the pre-edge "bump" observed in T.F.Y. mode arises from O atoms close to the oxide-Si(Ge) interface.

However, it is difficult to perform a quantitative analysis of O absorption at the Si(Ge)-oxide interface such as the one published by Muller et al. using the T.F.Y. mode. Indeed, this mode measures absorption of all O atoms of the oxide, including O atoms at the surface of the oxide which also have a reduced number of second-nearest O neighbours, and O atoms in the vicinity of the Si(Ge)-oxide interface. Moreover, the contribution of O atoms to the spectrum depends on the depth of the O atoms because of absorption of X-Ray photons in the SiO 2 film. A dedicated study would be required to retrieve this dependency.

A simple quantitative method is now attempted in order to draw tendencies. We assume that the T.F.Y. mode simply adds up absorption of O atoms from the bulk of the oxide, whose scattering factor is assumed to be the one of the Si sample in T.E.Y. mode, with the one of O atoms in an interfacial layer where all O atoms have the scattering factor at the interface published by Muller et al.. The equation of this linear combination is then: Absorption being proportional to the total quantity of O atoms, it is also proportional to the thickness of the considered layer. So, the thickness of the interfacial layer is:

f z f z f T F Y T E Y interface
Interfacial thickness ztotal oxide thickness = .

.

Firstly, we chose to apply this method at 532 eV because at this energy, the contribution of O atoms of the bulk of the oxide is minimum and equal to the tabulated values while the one of the interface is maximum. Table III.3 summarizes the total oxide thicknesses of the three samples measured by ellipsometry and the interfacial layer thickness calculated using Equation III.12 and Equation III.13. In summary, the thickness of the layer that has O atoms having scattering factors different from the one of the bulk of the oxide is higher for SiGe than for Si, even if reported values must not be taken as absolute values. In summary, there are more O atoms having scattering factors different from the one of the bulk of the oxide at the SiGe-oxide interface than at the Si-oxide interface.

Equation III.12 Equation III.13

We found no sound results with this procedure when applied to the whole spectra between 520 eV and 550 eV. We believe that the linear combination of Equation III.12 is not sophisticated enough to match reality. For instance, we could imagine a graded O scattering factor from the one at the interface published by Muller et al. to the one of bulk for the O atoms in a sub-oxide region in close proximity with the interface. And the dependence of the contribution of each O atom on depth was neither taken into account.

The fact that we cannot accurately retrieve the O scattering factors as a function of the distance from the Si(Ge)-oxide interface does not allow proper fitting of XRR spectra except below the absorption edge. Indeed, the O scattering factor below the absorption edge is not dependent on the environment of the scattering O atom and is tabulated. As a consequence, XRR spectra acquired at 500 eV only will be analyzed in the next sections using tabulated values of Si, Ge and O scattering factors [CXRO]. All XRR spectra were fitted using Dyna program (see Section II.4.1.6.). The thickness, the density and the roughness of each layer except the Si substrate were free fitting parameters. The roughness of the Si substrate was a free parameter while its density was fixed to its known value. Uncertainties of the measurement of the oxide density were calculated assuming that the statistical error is governed by Poisson statistics, which is the statistical error of a photon source [Tiilikainene07]. We included a synthetic error into the count of photons at each measurement angle of XRR experiments using the following formula:

Si

I I I S θ θ η θ θ ( ) = ( )+ ( ) ( ),
where I, Is, θ and η are the measured count of photons, the synthetic count of photons, the angle between the incoming X-Ray beam and the surface of the sample, and a random number in the [-1,1] interval. We then fitted numerous synthetic XRR spectra to get the statistical error on the oxide density. We report as error bars on all following figures the maximum deviation obtained with this method.

Equation III.14

XRR experimental spectra were fitted with two different models depending on the oxidized material:

ȃ Oxides grown on Si: a single SiO 2 layer on top of the Si substrate. The density, the thickness, the roughness of the oxide and the roughness of the Si substrate were free parameters. The density of the Si substrate was fixed to its known value.

ȃ Oxides grown on SiGe: a single SiO 2 layer on top of two SiGe layers and the Si substrate. The densities, thicknesses and roughnesses of the oxide and of the SiGe layers were free parameters. The density of the Si substrate was fixed to its known value.

Figure III

.25 shows the experimental and fitted XRR spectra of Si and SiGe samples oxidized in various conditions. Some portions of experimental curves are missing because signal jumps were removed. For oxides on Si, very high fit qualities were obtained with a single SiO 2 layer on top of the Si substrate.

As for SiGe, we observe that all spectra have major interference fringes and some minor contributions which warp the main interference fringes. The major fringes are due to reflection at the oxide-SiGe interface and minor fringes are due to reflections in the SiGe and Si layers. The thickness of the major fringes is related to the oxide thickness while their amplitudes are related to the optical contrast between the oxide film and the SiGe layer. For X-Rays, the optical contrast is related to gaps in terms of electron densities. The electron density is also related to the mass density, which is a more common and used parameter. The mass density of thermal oxides is around 2.2 g.cm -3 while the one of SiGe goes from about 2.3 g.cm -3 for Si to 4 g.cm -3 for Si 0.5 Ge 0.5 . Therefore, the higher the Ge concentration at the oxidation interface is, the higher the optical contrast between the oxide III. A.5.1.4

. Density of Thermal Oxides Grown on Si and SiGe

and the SiGe layer is. It explains why fringes are more pronounced for oxides on SiGe than for oxides on Si.

Rather good fit quality are obtained by using a model with a single oxide layer on top of two SiGe layers and the Si substrate. The fit quality is very good around the critical angle, which only depends on the oxide density (see Section II.4.1.3.). This gives us confidence about the reported oxide density values. In general, the fit quality is better for low q values, where spectra are more sensitive to the density contrast between the oxide and the SiGe layer, than for high q values where reflections at deeper interfaces have a higher contribution. We are thus confident about values of the oxide density extracted from these spectra and discussed into the next sections. ȃ 1: Initial Ge concentration: for both dry RTO and ISSG oxidation, the higher the initial Ge concentration is, the lower the oxide density is.

The density of oxides grown on Si and on SiGe is plotted

ȃ 2: Process type: the density of the oxide grown on Si 0.7 Ge 0.3 is higher for dry furnace than for dry RTO.

ȃ 3: Temperature: the oxide density of oxides grown on Si 0.7 Ge 0.3 by means of dry RTO increases with the temperature, as opposed to the case of Si. These results confirm the conclusion of Section III.A.4.2.2.: the oxide density is influenced by a combination of process type, initial Ge concentration and oxidation temperature. Following Section III.A.4. logical procedure, the next section considers the Ge concentration at the oxidation interface to study further SiGe oxidation kinetics.

Figure III.26: Evolution of the density of oxides grown on Si and on SiGe as a function of the oxidation temperature. The effects of the initial Ge concentration, of the process type and of the oxidation temperature on the oxide density are indicated by arrows (see text).

III.A.5.2. Dependence of the Oxide Density on Process Parameters

Figure III.27 shows the measured oxide density as a function of the Ge concentration measured by SIMS at the oxidation interface. We observe that the higher the Ge concentration at the oxidation interface is, the lower the oxide density is for dry RTO and ISSG oxidation. A shift is also observed between dry RTO at 900°C and dry RTO at 1100°C. We suggest that SiO 2 stress relaxation via viscous flow at high temperature (> 960°C) may account for the influence of temperature on the oxide density for the case of SiGe as well.

One mechanism can be proposed to explain the oxide density dependence on the Ge concentration at the oxidation interface. As suggested by Holland et al., the higher lattice parameter of SiGe compared with Si, at least the out-of-plane one for a fully strained SiGe film, would tend to accommodate the mismatch of atomic spacing between Si and SiO 2 [START_REF] Houghton | Strain relaxation kinetics in Si 1-x Ge x /Si heterostructures[END_REF]. As a result, the interfacial compressive stress in the grown oxide would be lower for oxides grown on SiGe. The lower compressive stress at the oxidation interface would correspond to a less dense oxide generated at the oxidation interface. The Ge concentration at the oxidation interface being rather constant during oxidation, except during a transient regime at the beginning of oxidation (see Section III. A.4.3.), the density of the grown oxide would be rather homogeneous in depth.

We also observe that the oxide density for the case of dry furnace oxidation at 900°C is not aligned with dry RTO at 900°C. We suggest that the longer temperature ramp up of furnace oxidation may influence strain in the SiGe layer prior to oxidation. Strain of the SiGe film after dry furnace oxidation could not be measured, so further experiments are required to study this suggestion.

Figure III.27: Evolution of the oxide density as a function of the Ge concentration at the oxidation interface.

To conclude, we discussed in this section the dependence of the oxide density on process parameters, i.e. oxidation process type and temperature, and on the Ge concentration at the oxidation interface. Into the next section, we show a correlation between the measured oxide density and oxidation kinetics by considering the dependence of the GRE on the oxide density.

III.A.5.3. Dependence of Oxidation Kinetics on the Oxide Density

The GRE of ISSG oxidation, dry furnace oxidation and dry RTO is plotted Figure III.28 as a function of the measured oxide density. The GRE of dry RTO at 900°C decreases when the oxide density increases. It is consistent with oxidation kinetics limited by diffusion of O 2 through the oxide and with the O 2 diffusivity that decreases when the oxide density increases. We also observe that the GRE of dry furnace oxidation at 900°C is close to the trend line of the evolution of the GRE of dry RTO at 900°C with the oxide density. It further confirms the dependence of the oxidation rate on diffusion of O 2 through the oxide. A similar decrease of the GRE when the oxide density increases is also observed for dry RTO at 1100°C. However, there seems to be a shift of the GRE of dry RTO at 1100°C to lower values than dry RTO at 900°C. This will be discussed in Section III.A.5.4.. As for ISSG oxidation, the GRE also decreases when the oxide density increases. However, the oxidation rate of SiGe by means of ISSG oxidation is close to the one of Si (the GRE is close to 1) while the oxide density gap is significant between the two measured samples. ISSG oxidation is mainly limited by diffusion of O* through the oxide (see Section III. A.4.1.4.). It thus shows that O* diffusivity is not or hardly dependent on the oxide density. However, we can still discuss the slight GRE dependence on the oxide density. The GRE could be due to a slightly higher O* diffusivity when the oxide density decreases. But O* is not the only oxidizing species for ISSG oxidation [Kee00], so that the GRE of ISSG oxidation could also be due to a higher diffusivity of another oxidizing species (such as O 2 ) when the oxide density decreases.

Figure III.28: Evolution of the GRE as a function of the measured oxide density.

In this section, it was shown that oxidation kinetics depends on the oxide density: the higher the oxide density is, the lower the GRE is. Into the next section, thanks to the dependence of the GRE on O 2 diffusivity and on the oxide density, we will open a window on a mechanism which is not specific to this study: the influence of the oxide density on O 2 diffusivity.

According to thermodynamic considerations about the temperature and pressure dependence of the diffusion coefficient detailed in [Mehrer07], the diffusion coefficient of O 2 in SiO 2 is:

D D H k T D E p V k T B B = -       = - +       0 0 exp e xp , ∆ ∆
where D 0 , ΔH, k B , T, E, p and ΔV are a pre-exponential constant, the activation enthalpy of diffusion, the Boltzmann constant, the temperature in K, the activation energy of diffusion, the pressure and the activation volume. Some authors choose to express pΔV as a function of another parameter because pΔV is sometimes not known and less convenient to use than other related parameters in real diffusion problems. . Whatever the parameter they chose to use is, their models are very similar in their forms and can be generalized to:

D D E k T D B = -       0 0 exp ex through oxide grown on Si p p , - +              E E k T B ∆
through oxide grown on SiGe

where E and ΔE are the activation energy in eV of O 2 diffusion in SiO 2 grown on Si and the diffusion barrier modulation in eV due to the density variation. The density variation due to the temperature only is included into the activation energy E and not in ΔE. In other words, the activation energy E is dependent on T because of the density variation with the temperature. We will use this more generalized model because we do not know stress levels in the oxide and will not make any assumption on the dependence of pΔV on any other parameter such as density.

According to Equation III.8, the GRE is equal to the ratio of O 2 diffusivity in the oxide grown on SiGe to the one in the oxide grown in Si. Combining Equation III.8 with Equation III.16, the GRE is:

III.A.5.5. Conclusion of Section III.A.5.

We previously showed in Section III. A.4. that oxidation kinetics of SiGe is limited by diffusion of the oxidizing species through the oxide. As the oxide density is known to influence diffusion, this section aimed for studying SiGe oxidation kinetics in light of the measurement of the oxide density.

After having presented the experimental method based on XRR measurements, we studied the dependence of the density of oxides grown on SiGe on the initial Ge concentration, on the process type and on the oxidation temperature. We showed that the oxide density decreases when the oxidation temperature increases and suggested that SiO 2 stress relaxation via viscous flow at high temperature (> 960°C) may account for this temperature influence. We also found that the density of the oxide grown on SiGe by means of dry furnace oxidation is higher than for dry RTO, thus showing a strong influence of the process type. And finally, we found that the higher the Ge concentration at the oxidation interface is, the lower the oxide density is. It was suggested that the higher lattice parameter of SiGe with respect to Si would tend to accommodate the lattice mismatch between Si and SiO 2 , and as a consequence lower the density of the growing oxide.

Then, we showed that the lower the density of the oxide is, the higher the oxidation rate for dry oxidation is, in agreement with an oxidation regime that is governed by diffusion of O 2 through the growing oxide.

Finally, we proposed a model of O 2 diffusivity that depends on the oxide density through the modulation of the diffusion barrier. This model is consistent with experimental oxide densities measured on SiGe samples and with the model of O 2 diffusivity proposed by Cui et al. [Cui09].

III.A.6. Conclusion of Section III.A.

In Section III.A., based on a literature review, we defined two problematics:

ȃ What is the composition of the thermal oxide grown on SiGe and does it agree with thermodynamic predictions (Section III.A.3.)? ȃ What are the mechanisms and kinetics of thermal oxidation of SiGe (Section III.A.4. and Section III.A.5.)?

In Section III.A.3., we showed that, except possibly during the very first stages of oxidation (the temperature ramp up), the composition of oxides grown on SiGe is pure SiO 2 for all investigated oxidation processes, in agreement with thermodynamic considerations of Section III.A.1.2.. We found that a very small quantity of Ge oxide (less than 5%) may be formed during the temperature ramp up when the oxidation rate largely exceeds the SiGe diffusion rate.

Then, we showed that oxidation of SiGe by means of dry furnace oxidation, dry RTO and ISSG oxidation is limited by diffusion of the oxidizing species through the growing oxide. A factor called Growth Rate Enhancement (GRE) was then defined to quantitatively compare the oxidation rates of Si and SiGe. We found that the GRE, and consequently oxidation kinetics, is correlated with the oxide density via the modulation of O 2 diffusivity. The influence of the process parameters on the oxide density, and as a result on oxidation kinetics, was finally clarified: the oxide density decreases when the oxidation temperature and the Ge concentration at the oxidation interface increase. Finally, we proposed a model of O 2 diffusivity that depends on the oxide density through the modulation of the diffusion barrier and which is consistent with our experimental data. Some mechanisms that could account for the dependence of the oxide density on the oxidation temperature and on the Ge concentration at the oxidation interface have been suggested. The higher lattice parameter of SiGe with respect to Si would tend to accommodate the lattice mismatch between SiGe and SiO 2 , and as a consequence decrease the density of the oxide. As for the effect of temperature, it was proposed that stress relaxation of SiO 2 via viscous flow could decrease the oxide density.

In the second part of this chapter (Section III.B.), two technological applications based on these results are presented.

III.B. Application to Industrial Processes for Advanced FD-SOI CMOS Manufacturing

The second part of this chapter presents two tools built based on the former oxidation kinetics results. The first one is a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, a key parameter of the evolution of the SiGe composition during oxidation, as a function of process parameters. It is designed for a convenient and quick evaluation of the relative importance of the Ge pile-up mechanism with respect to SiGe interdiffusion. The second one is a numerical simulation tool of oxidation of SiGe which can be used both for a finer understanding of the dynamics of SiGe oxidation and for accurate prediction of output parameters of the condensation process (such as the oxide thickness and the SiGe composition profile in the SGOI film).

III.B.1. Condensation Abacus: Oxide Thickness and Ge Diffusion Length

There are two key variables of the condensation process that can be used for optimizing it:

ȃ The oxide thickness after condensation: indeed, when fixing the epitaxial SiGe layer thickness and Ge concentration, the final SGOI layer thickness and Ge concentration directly depend on the oxide thickness at the end of the condensation process (this will be be shown in Section IV.A.1.). The oxide thickness may therefore be considered as a target specification of the condensation process.

ȃ The self-diffusion length of Ge in SiGe: depending on the needs in terms of SGOI composition after the condensation (for instance, full homogenization or on the contrary hardly any diffusion of Ge into the SOI layer may be required), SiGe interdiffusion must be more or less pronounced. How far Si and Ge atoms diffuse into SiGe describes how homogeneous the SGOI layer will be after condensation. We will use a deliberately simple model for SiGe interdiffusion: the Ge self-diffusion length in SiGe is a fairly good variable of how far atoms diffuse and may therefore be considered as a variable of the condensation process.

Precise prediction of SGOI layer thickness and composition profile requires the use of simulations such as the ones presented in Section III.B.2.. However, these simulations are not appropriate for an easily understandable display of key variables of the condensation process. We thus give Figure III.30 and Figure III.31 two graphic and convenient tools for assessing the key variables of the condensation process as a function of two process parameters: oxidation temperature and duration. These tools are temperature-time charts showing isolines of oxide thicknesses and of Ge self-diffusion lengths for two SiGe compositions relevant to the condensation process (0% and 25%). Isolines of oxide thicknesses and Ge self-diffusion were chosen to be plotted on temperature-time charts with x and y axis ranges relevant to Rapid Thermal Processes (dry RTO and ISSG oxidation), i.e. from 850°C to 1150°C for the oxidation temperature and from 10 seconds to 1000 seconds for the oxidation duration. Isoline values were also chosen to be relevant to the condensation process, from 100 Å to 200 Å for the oxide thicknesses and from 10 Å to 200 Å for the Ge self-diffusion lengths.

Oxidation kinetics and SiGe interdiffusion are inter-dependent. Therefore, the fact that the BOx layer plays the role of a diffusion barrier in the condensation process influences SiGe composition evolution and consequently oxidation kinetics of the condensation process. Oxidation kinetics and SiGe composition evolution are thus different depending on the substrate, either bulk Si or SOI, on which SiGe is epitaxially grown. We compared the oxide thickness after oxidation of SiGe on bulk Si (samples of Section III.A.) and after condensation (samples of Section IV.A.). We found that differences between the two substrate types are negligible for SGOI final Ge concentrations below 50%. And as already stated before, these charts aim at giving a visual assessment of the key variables of the condensation process and not at precisely predicting SGOI composition profiles. Oxide thicknesses are calculated using equations of Section III.A.4.1.1. and Ge self-diffusion lengths are calculated using data published by Kube et al. [Kube10] according to the equation:

L Dt = .
where L is the diffusion length, D the diffusion coefficient and t the oxidation duration. Napolitani et al. have shown

Equation III.18

123/230 Chapter III -SiGe Thermal Oxidation that oxidation of Si and SiGe induces injection of Si interstitials into the Si/SiGe film below the oxidation interface [Napolitani05]. They also observed that the interstitials injection mechanism depends on the Ge content. SiGe interdiffusion being mediated by interstitials and vacancies, the oxidizing environment most likely alters SiGe interdiffusion. We do not have quantitative data on the injection of interstitials for our processes, so that we can only use published data while assuming that tendencies are similar in an oxidizing environment.

The temperature dependence of dry RTO is higher than the one of ISSG oxidation because the activation energy of dry RTO is higher than the one of ISSG oxidation (about 2 eV for dry RTO against 1 eV for ISSG oxidation). At 1100°C, dry RTO and ISSG oxidation have similar oxidation rates. While at lower temperatures, ISSG oxidation yields a much higher oxidation rate, which induces much shorter oxidation durations for ISSG oxidation for the same target oxide thickness. This temperature dependence of ISSG oxidation versus dry RTO can be seen in the slopes of isolines: the steeper the slope is, the lower the temperature dependence is. Dry RTO at temperatures lower than 900°C is not suitable for oxides thicker than 100 Å because oxidation times would have to be higher than 10 minutes, which is out of specification for a Rapid Thermal Process. The activation energy of SiGe interdiffusion is about 4 eV, thus much higher than the one of dry RTO (around 2 eV) and ISSG oxidation (1 eV). Similarly to oxidation, such a temperature dependence can be seen in the slopes of isolines, which are gentler than the ones of dry RTO and ISSG oxidation plotted on Figure III.30. Isolines for diffusion in Si and Si 0.75 Ge 0.25 are almost parallel on this chart because the activation energies of Ge self-diffusion in Si and in Si 0.75 Ge 0.25 are close (4.83 eV for Si against 4.71 eV for Si 0.75 Ge 0.25 ). We emphasize that isolines are highly dependent on the SiGe composition: the higher the Ge concentration is, the higher the Ge self-diffusion length is. For instance, isolines of 50 Å for diffusion in Si and of 200 Å for diffusion in Si 0.75 Ge 0.25 are very close.

Figure III.31: Oxidation temperature-time charts for Ge self-diffusion in SiGe. Isolines of Ge self-diffusion lengths are displayed on chart.

Let us illustrate how these charts can be used with an example relevant to this work. Let us assume that, for some technology-related reasons, we need to grow a 100 Å-thick oxide film on a Si 0.7 Ge 0.3 -SOI stack while keeping a very steep interface between the Si 0.7 Ge 0.3 and the SOI films. The later constraint can be translated into a constraint on the Ge diffusion length in Si, for instance that it cannot exceed 10 Å. This example is illustrated Figure III.32.

We can draw the isoline of 10 Å Ge self-diffusion length in Si on Figure III.31. The isoline defines the process window that respects this criterion and is plotted Figure III.33. Within this process window, we could for instance perform ISSG oxidation at 1035°C for 37 seconds, or at 900°C for 93 seconds, or dry RTO at 1015°C for 82 seconds. This methodology can be extended to other criteria such as the diffusion length of a dopant, or the maximum temperature that does not induce phase change of another material.

III.B.2. Simulation of SiGe Thermal Oxidation

III.B.2.1. Introduction

As already stated in Section I.5.1.2., numerical simulation of thermal oxidation of SiGe has two interests at stake: it can be used for a finer understanding of mechanisms and kinetics of SiGe oxidation; and it could also be used to accurately predict two key outputs of the condensation process: the oxide thickness and the SiGe composition profile. In this section, we present the very first results of simulation of thermal oxidation of SiGe using a FEM simulation tool developed in COMSOL Multiphysics environment. After introducing the main features and assumptions of the model, we compare simulated kinetics and SiGe composition profiles evolution during dry RTO and ISSG oxidation of SiGe with experimental data. As a conclusion, we highlight some key findings about oxidation of SiGe evidenced by simulation as well as current limitations and perspectives about simulation of thermal oxidation of SiGe.

III.B.2.2. Simulation Model

The model was derived from the model built by Pierre Morin [Morin15]. There are three main mechanisms that govern thermal oxidation of SiGe and that must be modeled: ȃ Oxide growth rate: According to the definition of the GRE given Section III.A.4.2.1., the oxide growth rate on SiGe is:

dx dt x dx dt x GRE SiGe Si ( ) = ( )
, where x is the oxide thickness. The oxide growth rate on Si is:

dx dt x x A E k T x Si Si ox B ( ) = = -       α 2 2 exp ,
where α SiGe is the parabolic constant of dry RTO of Si and where A and E ox are the pre-exponential constant and the oxidation activation energy. Both parameters are calculated according to Section III. A.4.1.4. (E ox is equal to 2.29 eV for dry RTO and 0.95 eV pour ISSG oxidation, and A is equal to 1.32×10 10 Å 2 .t -1 for dry RTO and 3.01×10 5 Å 2 .t -1 for ISSG oxidation).

Equation III.19 Equation III.20

For ISSG oxidation, there is no strong correlation between the oxidation rate and the Ge concentration. We thus consider that the GRE is equal to 1.

For dry RTO, we found that the GRE is dependent on the Ge concentration at the oxidation interface (see Section III.A.4.4.). Figure III.34 shows the experimental data along with the best fit we found, which is an exponential function of the Ge concentration at the oxidation interface Ge% ox (in percent): ȃ SiGe interdiffusion: the model of SiGe interdiffusion is based on Darken's equations with a unique self-diffusion coefficient calibrated by Morin et al. [Morin15]:

D x D S x E x kT PV kT E kT S a a C ( ) = ( ) ( ) -( )       -       -  0 exp exp exp e xp       ,
where x is the Ge fraction, k and T are respectively the Boltzmann constant and the temperature, S the SiGe mixture entropy, E a the enthalpy (related to the formation and migration of vacancies in SiGe), P the pressure related to the stress components, V a the activation volume. The effect of strain on diffusion is accounted by the product PV a where P x y z = + + ( )/ σ σ σ 3. Hence the model also calculates the stress components in the SiGe layer according to the formalism presented in Appendice A. The term E c was added to account for the injection of interstitials at the oxidation interface (see Section III.A.1.3.3.). The values of the parameters of the diffusion coefficient have been fitted on published self-diffusion data [Kube10] and experimental data acquired previously to this work at IBM Albany Alliance site.

III.B.2.3. Kinetics of Dry RTO and ISSG Oxidation of SiGe

Equation III.20 shows that the oxide thickness must not be null. We thus enter an initial oxide thickness into the simulation model. Experimental and simulated oxides thicknesses of dry RTO of Si at 1000°C as a function of oxidation time are shown Figure III.35. It appears logical in the first instance to enter the native oxide thickness as the initial oxide thickness (about 1 nm). However, as illustrated Figure III.35, it may not be the initial conditions that correspond to the best fit of experimental data. In the case of dry RTO of Si at 1000°C, the best initial conditions seem to be an initial 2 nm-thick oxide.

Equation III.22

This apparent contradiction is due to the fact that there is a transient regime at the very beginning of oxidation. Indeed, the oxide thicknesses extrapolated from the fitted parabolic laws of Section III.4.1.1. are not equal to zero or one nanometer when the oxidation time is zero. This transient regime may be due to oxidation during the temperature ramp up and/or altered oxidation kinetics for very thin oxides. As we do not know what this regime is, we should enter the initial conditions of the simulation that correspond to the beginning of the actual parabolic regime. A further problem that arises with SiGe is that the Ge concentration at the oxidation interface changes very rapidly during the first stages of oxidation (see Section III. A.4.3.). The SiGe composition profile at the beginning of the parabolic regime is not known either. We chose to start simulations with an initial oxide layer on top of a single SiGe layer in the first instance. We determined for each oxidation conditions the initial oxide thickness that corresponds to the best fit of experimental data. Figure III.36 compares the oxide thicknesses simulated with the best initial oxide thickness with experimental data for all dry RTO and ISSG oxidation conditions. Simulation of dry furnace oxidation of SiGe cannot be performed as we do not know the dependence of the GRE on the Ge concentration at the oxidation interface. Simulation of ISSG oxidation kinetics does not depend on the Ge concentration (the GRE is fixed to one as explained Section III. B.2.2.), hence data of ISSG oxidation of Si only are shown.

We observe a rather good agreement between simulated and experimental data except for the case of dry RTO of Si 0.9 Ge 0.1 at 900°C. It seems that starting simulations with an initial oxide layer on top of a single SiGe layer does not drastically alter the simulated oxide growth kinetics. In summary, the simulation model based on the GRE equation is rather consistent with experimental data in terms of oxide growth kinetics. We examine into the next section the evolution of the Ge concentration profile and compare it with SIMS measurements already shown in Section III.A.4.3..

III.B.2.4. Evolution of the SiGe Composition Profile

Experimental SIMS and simulation profiles of the Ge concentration after various dry RTO and ISSG oxidation processes are shown Figure III.37. Due to simulation constraints, the same initial SiGe epitaxial thickness was used while it was not the case of experiments. However, these differences are not expected to influence the tendencies observed on Figure III.37.

In agreement with the discussion of Section III.A.4.3., we notice that the higher the oxidation temperature is, the lower the simulated Ge concentration at the oxidation interface is. Simulations show the same dependence on temperature of the competition between oxidation that tends to pile-up Ge atoms below the oxidation interface and SiGe interdiffusion that tends to smooth Ge concentration gradients.

Yet, in terms of absolute values, we observe that the Ge concentration gradients are smoother in experiments than in simulations in all cases. The systematically higher Ge concentration at the interface in simulations shows that the pile-up mechanism is more pronounced in simulations. In addition, the Ge concentration gradients between the pile-up layer and the epitaxial SiGe film and between the SiGe film and the Si substrate are also smoother in experiments. Even if SIMS measurements tend to smooth concentration gradients due to the snow-plow effect, we do not expect it to cause such differences.

We have shown in the former section that the experimental and simulated oxidation rates are very similar so that these differences can only arise from SiGe interdiffusion rates slower in simulations than in experiments. The SiGe diffusion coefficient used in simulations was the one calibrated by Morin et al. [Morin15] (see Section III. B.2.2.) with published self-diffusion data [Kube10] and dedicated experiments. These dedicated experiments are different from the experiments performed during this thesis in two ways.

Firstly, oxidation processes performed in this study are different in various ways from the ones performed by Morin et al. (the details cannot be disclosed). As oxidation of Si(Ge) is known to inject a certain amount of Si and/ or Ge interstitials into the Si(Ge) film (see Section III.A.1.3.3.), the concentrations of interstitials and vacancies are most likely influenced by oxidation conditions. SiGe interdiffusion being mediated by diffusion of interstitials and vacancies (see Section I.2.5.2.), the SiGe diffusion coefficient is consequently influenced by oxidation conditions.

Secondly, they calibrated their model with condensation experiments, therefore with a SOI substrate, while these experiments are performed with bulk Si substrates. Si(Ge)-oxide interfaces, as any interface in diffusion experiments, likely play a role in recombination and generation of interstitials and vacancies, and as a result on SiGe interdiffusion.

Whatever the reasons for these differences are, there is a discrepancy between the simulated oxidation rates, which are very close to the experimental ones, with the simulated Ge concentration at the oxidation interface, from which the oxidation rate is calculated in simulations. Due to time constraints, we could not further investigate about this discrepancy. As tendencies are still consistent with experiments, we will discuss into the next section the evolution of the Ge concentration at the oxidation interface during oxidation, and its dependence on process parameters.

III.B.2.5. Evolution of the Ge Concentration at the Interface

Figure III.38 shows the simulated and experimental Ge concentration at the oxidation interface as a function of oxidation time for various dry RTO and ISSG oxidation conditions. In all cases, the simulated Ge concentration at the interface rapidly increases until it reaches a plateau. As previously observed, the higher the oxidation temperature is, the lower the Ge concentration at the oxidation interface is. The same discrepancy in terms of absolute values is noticed: the simulated Ge concentration at the interface is systematically higher than the experimental one.

The only case that is slightly different is dry RTO of Si 0.7 Ge 0.3 at 1100°C: the Ge concentration at the interface starts to decrease after about 30 seconds. It is consistent with experiments, where the Ge concentration at the interface decreases from 38.6% after 10 seconds to 32.9% after 60 seconds.

Investigating why the Ge concentration at the oxidation interface is rapidly increasing until a plateau would require to compare in a quantitative manner the oxidation rate with the diffusion rate as a function of time. Such a problem cannot be solved analytically (hence the simulations) so that further discussion about this phenomenon is difficult at the present time.

In summary, we highlight that despite of the discrepancy between experimental and simulated results in terms of absolute values, simulations are consistent with what was experimentally shown in Section III. A.4.3.: the Ge concentration at the oxidation interface is rather constant after a transient regime.

Figure III.38: Experimental and simulated Ge concentration at the interface as a function of time.

The Ge concentration at the interface is crucial not only for dry RTO kinetics, but also in terms of oxide density (see Section III.A.5.). The oxide density can also influence further processing of CMOS devices (such as chemical etching of the oxide), so that it may be useful to be able to predict the Ge concentration at the interface as a function of process parameters. Even if simulated values must not be considered as absolute values, tendencies are valid and valuable information.

Figure III

.39 shows the simulated Ge concentration at the interface once the plateau is reached, and plotted as a function of the initial Ge concentration and oxidation temperature. We noticed that in all cases, the plateau is reached when the oxide is 13 nm-thick (figure not shown), so that we decided to extract the Ge concentration at the 131/230 Chapter III -SiGe Thermal Oxidation interface at this thickness.

As expected, the higher the initial Ge concentration and the lower the temperature are, the higher the Ge concentration at the interface is. The temperature dependence of ISSG oxidation is higher than of dry RTO. Indeed, as already discussed in Section III. A.4.3. and in Section III.B.1., ISSG oxidation at low temperatures is much faster than dry RTO, which makes the gap between the oxidation rate and the diffusion rate higher, and results in a higher Ge concentration at the oxidation interface. The dependence on the initial Ge concentration seems similar for dry RTO and ISSG oxidation. Then, two trends can be emphasized for both processes:

ȃ The dependence of the Ge concentration at the interface on the initial Ge concentration is higher at higher temperatures. ȃ The dependence of the Ge concentration at the interface on temperature is higher at lower initial Ge concentrations. III.4 summarizes the dependence of the Ge concentration at the oxidation interface on the process type, the initial Ge concentration and temperature.

Initial Ge concentration Oxidation temperature

Dry RTO

+ ++

ISSG oxidation

+ +++

Table III.4: Relative weight of the dependence of the Ge concentration at the interface on the process type (dry RTO and ISSG oxidation), the initial Ge concentration and the oxidation temperature. + refers to a low dependency while +++ to a high one.

III.B.2.6. Conclusion of Section III.B.2.

We implemented the experimentally observed GRE dependence on the Ge concentration at the interface into the simulation model built by Morin et al. [Morin15]. We showed that simulated kinetics of dry RTO and ISSG oxidation of SiGe is in good agreement with experiments, in spite of the difficulty to model the transient regime of the beginning of oxidation. We evidenced a discrepancy between experiments and simulations in terms of SiGe composition profile evolution, especially at the oxidation interface. We believe that this discrepancy is due to a non-adapted calibration of SiGe interdiffusion parameters. Nevertheless, it allowed us to highlight some tendencies, and especially the dependence of the Ge concentration at the oxidation interface on process parameters.

In a summary, the simulation tool already proved its usefulness in studying the dynamics of thermal oxidation of SiGe as it successfully reproduces tendencies in terms of oxidation kinetics and SiGe composition profile evolution. However, as expected given that the simulation model is an initial approach, it requires further improvement to be able to accurately predict a condensation process.

Conclusion of Chapter III

In Section III.A., based on a literature review, we defined two problematics:

ȃ What is the composition of the thermal oxide grown on SiGe and does it agree with thermodynamic predictions (Section III.A.3.)? ȃ What are the mechanisms and kinetics of thermal oxidation of SiGe (Section III. A.4. and Section III.A.5.)?

In Section III.A.3., we showed that, except possibly during the very first stages of oxidation (the temperature ramp up), the composition of oxides grown on SiGe is pure SiO 2 for all investigated oxidation processes, in agreement with thermodynamic considerations of Section III.A.1.2.. We found that a very small quantity of Ge oxide (less than 5%) may be formed during the temperature ramp up when the oxidation rate largely exceeds the SiGe diffusion rate.

Then, we showed that oxidation of SiGe by means of dry furnace oxidation, dry RTO and ISSG oxidation is limited by diffusion of the oxidizing species through the growing oxide. A factor called Growth Rate Enhancement (GRE) was then defined to quantitatively compare the oxidation rates of Si and SiGe. We found that the GRE, and consequently oxidation kinetics, is correlated with the oxide density via the modulation of O 2 diffusivity. The influence of the process parameters on the oxide density, and as a result on oxidation kinetics, was finally clarified: the oxide density decreases when the oxidation temperature and the Ge concentration at the oxidation interface increase. Finally, we proposed a model of O 2 diffusivity that depends on the oxide density through the modulation of the diffusion barrier and which is consistent with our experimental data. Some mechanisms that could account for the dependence of the oxide density on the oxidation temperature and on the Ge concentration at the oxidation interface have been suggested. The higher lattice parameter of SiGe with respect to Si would tend to accommodate the lattice mismatch between SiGe and SiO 2 , and as a consequence decrease the density of the oxide. As for the effect of temperature, it was proposed that stress relaxation of SiO 2 via viscous flow could decrease the oxide density. As a perspective to this work, we suggest that studying oxidation kinetics of fully relaxed SiGe films and varying the conditions of the temperature ramp up of each oxidation process (which are largely different from dry RTO to dry furnace oxidation) could shed light on the mechanisms that induce this dependence.

In the second part of this chapter (Section III. B.), two technological applications based on the former results were presented. The first one is a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, a key parameter of the evolution of the SiGe composition during oxidation, as a function of process parameters. The tool compares the relative importance of the Ge pile-up mechanism caused by the Si-selective oxidation with respect to SiGe interdiffusion. And secondly, a numerical simulation tool of thermal oxidation was adapted from the one of Morin et al. [Morin15]. The simulation model proved its worth in bringing more understanding of the dynamics of SiGe oxidation in spite of some remaining discrepancies with experimental data. We concluded with some suggestions that could improve the model, i.e. dedicated experiments to calibrate the diffusion model on the specific processes performed in this work.

IV.B. Strain Relaxation at the Edge of Active Areas

IV.B.1. Problem Definition

IV.B.2. Chemical Bonding at the SGOI-BOx interface Investigated by XPS

IV.B.3. Measurement of Hydrogen Concentration by Neutron Reflectivity

Coupled with XRR

Conclusion of Chapter IV

Introduction to Chapter IV

After Chapter III has brought about understanding of the mechanisms and kinetics of thermal oxidation of SiGe, this chapter strives to study strain and crystal quality of SGOI films fabricated by condensation. Indeed, we highlighted in Section I.5.2. that high strain levels and high crystal quality are needed for increasing the hole mobility of SGOI channels.

Strain of full-sheet SGOI films is evaluated as a function of process parameters. The conditions for strain relaxation and its correlation with the formation of extended crystal defects are discussed for both dry Rapid Thermal Oxidation (RTO) and In-Situ Steam Generation (ISSG) oxidation processes. The process conditions and mechanisms that lead to strain relaxation and formation of dislocations are discussed. Then, the amount of atoms displaced from lattice sites in SGOI films is measured as a function of depth by the Medium Energy Ion Scattering (MEIS) technique. This amount of atoms displaced from lattice sites is a crystal quality indicator. The effects of the condensation process and of the Ge concentration on SGOI crystal quality are discussed.

Finally, the second part of this chapter starts by describing a particular SGOI strain relaxation phenomenon occurring during CMOS fabrication after the condensation process. The mechanical properties of the BOx-SGOI and SGOI-thermal oxide interfaces are thought to play a key role in this phenomenon. Two possible causes of these weaker mechanical properties are suggested and studied in Section IV. B.2. and Section IV.B.3. respectively. Section IV.B.2. studies the chemical bonding at the BOx-SGOI interface by X-Ray Spectroscopy (XPS) and Section IV. B.3. examines by neutron reflectivity the presence or not of hydrogen in BOx-SGOI-thermal oxide stacks. Where t Si , t SiGe , x, t Ox , t SGOI and y are the SOI thickness, the epitaxial SiGe thickness, the Ge content in the epitaxial SiGe film, the final oxide thickness, the final SGOI thickness and the Ge content in the SGOI film.

IV.A.1. Principles of the Condensation Technique

Si

When assuming negligible atomic density changes with Ge content (the lattice volume of Si and Ge are very close), the conservation of Ge can be written as:

t x t y

SiGe SGOI = .

And within the same assumptions and because 1 volume of Si is converted into 2.25 volumes of SiO 2 , the conservation of Si can be written as:

t t x t y t Si SiGe SGOI Ox + -= -+ ( ) ( ) . . 1 1 2 25
As the oxide thickness t Ox may be a reference parameter when calibrating an oxidation or a condensation process, it is useful to predict the final SGOI thickness t SGOI and Ge content y for a given structure as a function of the thermal oxide thickness t Ox . Combining Equation IV.1 and Equation IV.2, we provide the following equations for the SGOI thickness and Ge content in the SGOI layer as a function of the initial structure parameters and of the oxide thickness: Let us now consider for example that a 60 Å-thick SGOI layer (t SGOI = 60 Å) with 40% Ge concentration (y = 0.4) is targeted. We could consider the Ge total amount instead of the Ge concentration for this example. But for clarity considerations, we assume that the SGOI layer is fully homogeneous at the end of the condensation process, and that we can therefore use the Ge concentration, a more striking and used parameter. Let us also assume that

t

Equation IV.1 Equation IV.2 Equations IV.3

The condensation technique is based on concurrent Si-selective thermal oxidation of a SiGe layer epitaxially grown on a SOI substrate and homogenization of the SGOI layer by interdiffusion. As shown by other groups and in Section IV.A.3., there is no significant Ge loss during condensation, except possibly at the very beginning of the process (see Section III.A.3.). It was also shown that there is no Ge diffusion into the BOx layer within measurement error. The Ge quantity provided by the SiGe epitaxy process can thus be considered to be fully conserved in a well-defined SGOI layer sandwiched by the thermal oxide and the BOx layer. Moreover, thermal annealing processes that follow the condensation process for fabrication of functional devices are important enough to fully homogenize the SGOI layer if some composition gradient remains after the condensation process.

From the point of view of technologists, there are a few key variables related to both the SiGe epitaxy process and the condensation process that need to be well controlled. Especially, at the end of the condensation process, the Ge concentration and the thickness of the SGOI film are crucial because the final transistor performances directly depend on them. In this work, various Ge concentrations and SGOI thicknesses were targeted, so that a clear understanding of how these variables are interdependent was required for calibrating the condensation processes. For this purpose we consider the following model: the Ge concentration of the epitaxial SiGe layer x is fixed to either 25% or 35% (x = 0.25 or 0.35) and that the SOI thickness t Si is 70 Å. According to Equation IV.1, the thickness of the epitaxial SiGe layer t SiGe is 96 Å and 69 Å for 25% and 35% respectively. We will refer to C1 and to C2 for the cases of 25% and 35% initial Ge concentrations respectively.

IV.A. Strain and Crystal Quality of Full Sheet SGOI Films Fabricated by the Condensation Technique

Figure IV.2 shows graphic representations of the model of condensation for C1 and C2 cases.

Based on Equation IV.1, the Ge concentration of the SGOI layer as a function of its thickness is plotted Figure IV.2 (a). The more the SiGe layer is oxidized, or "condensed", the lower the SGOI layer thickness is because of Si atoms incorporation in the oxide. At the same time, the Ge concentration of the SGOI layer increases during condensation. Figure IV.2 (a) also shows the target point of the condensation process. We highlight that both C1 and C2 have the same evolution of the Ge concentration as a function of the SGOI thickness. Indeed, they have the same total Ge quantity in the starting epitaxial SiGe layer, with C1 having a SiGe layer that has a higher thickness but a lower Ge concentration. Now, in order to reach the target SGOI specifications, we need to calculate the target oxide thickness for C1 and C2 cases according to Equation IV.2. Figure IV.2(b) shows the evolution of the oxide thickness as a function of the Ge concentration of the SGOI layer, along with the target oxide thicknesses, which are 239 Å and 177 Å for C1 and C2 respectively. As long as oxidation proceeds, the thickness of the oxide and the Ge concentration of the SGOI film increase. We notice that the targeted oxide thickness of C1 is higher than C2. Indeed, the total Si quantity of C1 is higher than the one of C2 while the Ge quantity of C1 and C2 are identical. So, in order to reach the same SGOI target specifications, more Si atoms must be consumed by the oxidation process for C1 than for C2. In other words, the grown oxide must be thicker for C1 than for C2.

All condensation processes presented Section IV.A.2. and performed in this work were calibrated based on this model of condensation. Having presented the condensation technique into more details, the next section describes the experiments conducted to address the problematic defined in Section I.5.2.1..

IV.A.2. Experimental

IV.A.2.1. SiGe Epitaxy and Condensation Process

SOI substrates with a 100 Å-thick SOI layer and a 200 Å-thick BOx layer were used. Dry furnace oxidation was performed in order to decrease the SOI layer thickness to about 75 Å. After removal of the grown thermal SiO 2 layer, three types of SiGe layers were epitaxially grown full sheet, in other words on the whole surface of the wafer. The compositions and thicknesses of the grown layers were measured by ellipsometry.

Condensation was performed by either dry RTO or ISSG oxidation. At each moment of the condensation process, the temperature must be kept below the melting temperature of the SGOI layer in order to prevent melting of (a) (b)

SiGe. As shown Section IV.A.1., if we assume that the SGOI composition is homogeneous at every moment of the condensation process, the SGOI composition as a function of the oxide thickness is known. Then, as SiGe oxidation kinetics can be predicted (see Chapter III), the evolution of the SGOI composition during condensation can be predicted too.

For Ge concentrations below 50%, the melting temperature is higher than the maximum oxidation temperature used in this study (1100°C). However, above 50% Ge concentration, the condensation temperature must be decreased. Most of the oxidation recipes were designed so that the temperature decreases stepwise from 1100°C to 1000°C and finally to 900°C, as illustrated Figure IV.3. Temperature ramp down between temperature steps were 25°C.s -1 . A margin of about 50°C was kept at every moment between the condensation temperature and the SGOI predicted melting temperature because of three sources of bias on the prediction of the SGOI melting temperature: ȃ During condensation, the Ge concentration is never fully homogeneous because of the pile-up mechanism, even when using high temperatures such as 1100°C. The maximum Ge concentration of the SGOI layer is the one at the oxidation interface and is higher than the one calculated when assuming a homogenous composition of the SGOI layer. The predicted melting temperature of the SGOI layer when assuming homogeneous SGOI layers is thus lower than the actual melting temperature of the top part of the layer.

ȃ The type of substrate, bulk or SOI, influences SiGe interdiffusion (the BOx layer actually acts as a diffusion barrier, and interfaces are known to play a role in interstitials and vacancies generation and recombination, so on diffusion processes as well), which may in turn affect the Ge concentration at the oxidation interface.

As we know from the results of Chapter III that kinetics of dry RTO depend on the Ge concentration at the oxidation interface, dry RTO kinetics thus depend on the type of substrate.

ȃ Oxidation kinetics can be predicted according to the results of Chapter III for Ge concentrations at the oxidation interface below 50% only. However, some SGOI layers of this study have Ge concentrations above 50%, especially at the oxidation interface. As we cannot predict oxidation kinetics for Ge concentrations above 50%, we cannot either predict the evolution of the SGOI composition during condensation above 50% Ge concentrations. Oxidation process durations are limited by the specifications of the Rapid Thermal Processes equipment used for performing dry RTO and ISSG oxidation. The required process durations to form high-Ge content SGOI layers by a single condensation process exceed the equipment specifications. In those cases, two condensation processes which both comply with the equipment specifications were performed consecutively. A HF treatment was used to thin down the oxide between the two condensation processes in order to decrease the duration of the second condensation process. Indeed, the oxidation rate was shown to decrease when the oxide thickness increases (see Chapter III). Therefore, the thinner the oxide thickness before the second condensation process is, the shorter the second condensation process is for a fixed target SGOI composition. This sample preparation sequence is schematically illustrated Figure IV.4. For low Ge content films (lower than 50% Ge concentration), the temperature of oxidation (1100°C) is high enough for the SGOI film to be homogeneous in depth. And when, the temperature is decreased to 1000°C for high Ge content films, the diffusion rate increases fast enough with the Ge concentration (see Section III. B.1.) to still

IV.A.2.2. Characterization of SGOI Films

A dedicated method was used to measure SGOI layer thicknesses and compositions. Spectroscopic ellipsometry could not be used because SiGe optical constants are not known for Ge concentrations above 54%. On the other hand, while compositions cannot be accurately determined by XRR measurements, XRR is a reliable technique for measuring the thickness of SGOI and oxide layers for the whole range of SGOI compositions. The roughness of the SGOI-BOx and SGOI-thermal oxide interfaces and of the oxide surface increases with the Ge concentration, especially when strain relaxation occurs for high Ge-content SGOI layers (this will be shown later). The roughness of these interfaces were therefore treated as unknown during the fitting process of XRR spectra. For the sake of comparability, the SGOI layer thicknesses mentioned in this manuscript are the ones measured by XRR even for Ge concentrations below 54%.

Then, the SGOI composition was calculated assuming no Ge loss during condensation, as described Section IV.1.. We mention that methods based on other techniques, such as X-Ray Fluorescence (XRF) coupled with XRR and a technique based on hybridization of XRR and XPS [Fauquier17], allow measurement of SGOI thickness and/ or composition as well. However, they could not be used during the whole duration of this study and were thus ruled out for the sake of comparison.

Then, various techniques were used to study strain and crystal quality of SGOI films fabricated by condensation.

Top surface Atomic Force Microscopy (AFM) imaging was performed on 5 x 5 μm 2 SGOI surfaces after oxide removal by HF treatment in order to reveal dislocations. As will be shown Section IV. A.4.1.3., dislocations may consider that the SGOI film is homogeneous at the end of the condensation process.

induce atomic steps on the surface of the SiGe film along specific crystallographic directions, and AFM can be used to reveal these specific surface steps.

Strain levels were measured by Raman Spectroscopy assuming that the SGOI composition is the one determined by the method described above. Si-Si and Ge-Ge raman peak shifts were used for Ge concentrations below and above 60% respectively (see Section II.3.4.). Uncertainties in strain levels were calculated assuming the uncertainty of wavelength measurement is the monochromator resolution, i.e. 1 cm -1 . X-Ray Diffraction (XRD) could not be used because SGOI films are too thin for accurate strain determination with this technique. HRTEM images of a few representative samples were acquired. GPA of these images, taking as a reference an unperturbed area of the Si substrate below the BOx layer was performed to measure the in-plane and out-of-plane deformation of SiGe for comparison with Raman spectroscopy.

The MEIS technique was used to study the composition and the depth distribution of crystal defects of three reprentative samples. Angle configurations and details of the samples will be given Section IV.A.5.1.. The model of condensation described in Section IV.A.1. relies on Si and Ge conservation during thermal oxidation of SiGe. Other groups have already shown that there is no significant Ge loss during condensation, and we also checked that this is also the case for the condensation processes used in this study. Spectra of the Ge K-α and L-α spectral lines were measured by Wavelength Dispersive X-Ray Fluorescence (WDXRF) to compare Ge doses before and after condensation (see Appendice C for the description of the WDXRF technique). Ge K-α and L-α intensities of all three samples are very close to each other, which shows that there is no significant variation of the Ge amount before and after condensation. We showed in Section III.A.3. that there is significant Ge in the thermal oxide. Therefore, all Ge atoms are repelled below the growing oxide during oxidation and conserved in the sample. We could not perform the same measurement for condensation by means of ISSG oxidation but we will assume that there is no significant Ge loss for this process as well. In summary, there is no significant Ge loss during condensation of SiGe and that the model of condensation presented in Section IV.A.1. is valid. RTO at 900°C and 1100°C. Figures IV.6, IV.7, IV.8 and IV.9 show HRTEM images of various samples along with the in-plane and outof-plane deformation maps measured by GPA. Some caution must be observed when analysing GPA deformation maps because of all perturbations that alter the results. This is illustrated on Figure IV.7 (d) by the local fluctuations of the compressive in-plane deformation averaged over the thickness of the layer and plotted along an in-plane profile.

IV.A.3. Ge Conservation during Condensation

The contrast difference of the HRTEM image of the "Post Epitaxy" sample shows the presence of two distinct layers. The quite homogeneous green color for the in-plane deformation of the SOI-SiGe bilayer indicates that the in-plane deformation is close to zero, as expected for a fully strained SiGe layer. The out-of-plane deformation is close to zero (in green) in the bottom part of the layer, which corresponds to the SOI film, and positive in the top part of it (in red) because of the Poisson effect in the SiGe layer. The out-of-plane deformation is averaged along the in-plane direction over the whole HRTEM image and is plotted as a function of depth on Figure IV.6 (d). The spatial and strain resolutions of the GPA technique for this particular study do not allow a precise reconstruction of the out-of-plane deformation depth profile, but gives rather good qualitative agreement with what can be expected for a fully strained SiGe layer. Indeed, the out-of-plane deformation goes from about zero at the SOI-BOx interface to about 2% in the SiGe layer, in agreement with theoretical calculations detailed in Section I.2.2.3.. HRTEM images and in-plane and out-of-plane deformation maps of SGOI layers with Ge concentrations of 26%, 57% and 77%, referred to as SGOI26, SGOI57 and SGOI77, are also shown on Figures IV.7, IV.8 and IV.9. For both SGOI26 and SGOI57, the dominant color on the in-plane deformation map is green, hence the inplane deformation is close to zero. SGOI26 and SGOI57 would therefore have no or limited strain relaxation. As for the out-of-plane deformation, it is positive for SGOI26 and SGOI57 samples due to the Poisson effect.

The red shift of the in-plane deformation map for SGOI77 indicates that it is positive for the SGOI77 film, which thus shows strain relaxation. If it is fully relaxed, the in-plane and out-of-plane deformations should be identical (the in-plane and out-of-plane lattice parameters would be equal) while if the film is partially relaxed only, the out-of-plane deformation is expected to by higher because of the Poisson effect. However, we cannot determine if it is partially or fully relaxed because of the high uncertainty on absolute deformation values determined by GPA of HRTEM images.

The section starts by examining the evolution of elastic strain of SGOI films fabricated by condensation based on high temperature oxidation by means of either dry RTO or ISSG oxidation. Results are compared with literature state-of-the-art and discussed in terms of strain energy balance. Condensation based on low temperature ISSG oxidation is then performed to investigate the influence of the temperature and of the nature of the interfaces with the SiGe film on strain. Finally, the formation of dislocations in partially strain relaxed SGOI films is evidenced by AFM imaging of SGOI surfaces. Averaged values of compressive in-plane strains measured by GPA and Raman spectroscopy are reported on Figure IV.10. As described in Section II.3.4., the Si-Si and the Ge-Ge modes were used for extraction by Raman spectroscopy of the in-plane strain of SGOI layers with Ge concentrations below and above 60% respectively. Due to the dependence of the uncertainty calculation on the mode used for extraction of strain, the uncertainty is higher with the Ge-Ge mode than with the Si-Si one. The line called "Fully strained SiGe" is the theoretical strain of a fully strained SiGe film calculated based on Poisson theory (see Section I.2.2.3.).

As for GPA, for each SGOI layer, the in-plane strain was extracted from the measurement of the in-plane deformation averaged over two HRTEM images acquired at two different locations of the TEM lamella, minus the areas where the measurement is obviously altered by either defects such as dislocations or by a contrast too low for robust GPA. We chose to use the in-plane deformation instead of the out-of-plane deformation because there are more atomic planes analyzed in the in-plane direction than in the out-of-plane one (about 40 nm is analysed with the GPA method in the in-plane direction instead of less than 10 nm in the out-of-plane direction), and thus a better statistical significance. Moreover, there may be some Ge concentration gradients in the out-of-plane direction because of the condensation process, while the Ge concentration is quite uniform in the in-plane direction thanks to the good within wafer uniformity of the SiGe epitaxy and condensation processes. This concentration gradient may cause a gradient in the out-of-plane deformation because of the Poisson effect, which in turns increases the uncertainty of averaging the out-of-plane deformation over a full HRTEM image. We finally highlight that quantitative uncertainties cannot be defined for the GPA method (see Section II.3.5.) so that values are to be considered with some caution.

It is firstly confirmed that the SiGe layer of the post epitaxy sample is fully strained. As condensation proceeds and as the Ge concentration increases, the compressive in-plane strain increases according to the theoretical strain of a fully strained SiGe film up to about 55-60% Ge concentration and even starts to decrease above 63% Ge concentration. Up to at least 45% Ge concentration, SGOI layers are fully strained, while the SGOI57 layer has a compressive in-plane strain slightly lower than the one of a fully strained SiGe layer. However, the difference between the two values corresponds to the measurement uncertainty so that we cannot claim for certain that this SGOI layer is partially relaxed. In short and in qualitative terms, the SGOI strain follows a "bell shape" as a function of the Ge concentration with a maximum strain value found to be 2.1% at 63% Ge concentration.

The GPA measurement of the same SGOI57 film gives the same strain value, which tends to confirm that this sample is likely to be slightly relaxed. GPA measurement gives a compressive in-plane value lower than the one extracted from Raman spectroscopy for the SGOI77 film. GPA measurement for this sample is likely less robust than Raman spectroscopy because of two reasons. Firstly, the TEM lamella of this sample was too thick for optimum contrast of the SGOI layer, and as a consequence for the GPA method. Secondly, as will be shown in Section IV. A.4.3., the numerous extended defects of the SGOI77 layer alter the GPA measurement. Whatever the robustness of the GPA measurement for this SGOI layer is, both techniques qualitatively agree and show a strong strain relaxation of the SGOI77 film. A study similar to the one conducted for dry RTO was conducted for condensation based on ISSG oxidation. Figure IV.11 shows the measured compressive in-plane strain of SGOI films as a function of the Ge concentration for both dry RTO and ISSG oxidation. The SGOI31 layer fabricated by the latter process type is fully strained. However, the SGOI48 film seems to be already partially relaxed. As for the SGOI73 layer, it is also partially relaxed but still features a high compressive in-plane strain value of 2.2 %. The evolution of strain as a function of the Ge concentration follows similar trends for dry RTO and ISSG oxidation, but with a slight difference shown by the dotted lines on Figure IV.11. ISSG oxidation-based condensation seems to lead to strain relaxation for Ge concentration as low as 48% and to a maximum strain level attained at a higher Ge concentration than for dry RTO.

Compressive

Two possible reasons for this difference are suggested. Firstly, and similarly to dry RTO, the temperature of condensation had to be lowered to 1000°C for fabrication of SGOI48 and SGOI73 by ISSG oxidation (see Section IV. A.2.1.). ISSG oxidation is faster than dry RTO at 1000°C (see Section III. B.1.), so that the SGOI film at the end of the condensation process is likely not as homogeneous as SGOI films fabricated by condensation based on dry RTO because of a stronger pile-up effect. As shown by Takagi et al., Ge concentration gradients tend to favor SGOI strain relaxation. Therefore, the fact that the SGOI48 sample is already partially relaxed could be due to a higher Ge concentration gradient for the case of ISSG oxidation than for dry RTO.

Secondly, oxidation of SiGe by means of dry oxidation is known to possibly inject Si and/or Ge interstitials into the SiGe film (see Section III.A.1.3.3.). We found no data on the injection of interstitials during ISSG oxidation of SiGe. If we assume that the mechanisms of strain relaxation of SGOI films (they will be discussed in Section IV. A.4.1.5. and Section IV.A.4.1.6.) are influenced by the injection of Si and/or Ge interstitials and that the two processes have different injection rates of interstitials, strain relaxation could be influenced by the process type. Strain relaxation might be favored in the case of ISSG oxidation compared with dry RTO for Ge concentration around 45%-50%, which would explain strain relaxation of the SGOI48 sample, while it might maintain a higher strain for Ge concentrations above 60%, which would correspond to the case of the SGOI73 sample. Especially, injection of interstitials is known to possibly play a role in the formation of extended defects in crystals, which are associated with plastic strain relaxation (the formation of extended defects will be discussed in Section IV. A.4.3.).

Studying more SGOI films with Ge concentrations above 50% is needed for a more accurate determination of the strain evolution of SGOI films formed by condensation based on high temperature ISSG oxidation. . They also found that the compressive strain of the SGOI layer increases with the Ge content until a peak strain value. Above the Ge concentration that corresponds to this peak value, the compressive in-plane strain decreases. Figure IV.12 shows the compressive in-plane strain values published by Takagi et al. and obtained with their best condensation process, and values from this work. We observe that the "bell shapes" of the evolution of the compressive in-plane strain as a function of the Ge concentration are very similar, with a maximum compressive strain of 2.2% at 60% Ge concentration for Takagi et al. and 2.1% at 63% Ge concentration for this work. Takagi et al. [Takagi10] and obtained with their best condensation process, and values from this work measured by Raman spectroscopy.

IV.A.4.1.2. Comparison between Dry RTO and ISSG Oxidation

Figure IV.12: Compressive in-plane strain values published by

In this first approach, it seems that there is no difference between dry furnace oxidation and dry RTO in terms of strain evolution of SGOI films as a function of the Ge concentration. However, Takagi et al. pointed out that the analysis of the evolution of strain of SGOI films cannot be based on the consideration of the Ge concentration only. of the film especially must be considered. A more relevant analysis framework of these results is the one of elastic strain energy, as will be discussed into the next section.

IV. A.4.1.4

. Comparison with SiGe Equilibrium Models

It appears necessary to identify the mechanisms that govern the evolution of strain of SGOI films. As already highlighted by Takagi et al., a common framework of the analysis of elastic strain is the one of the balance of elastic strain energy with plastic strain energy, the so-called equilibrium models of the critical thickness (see Section I.2.4.2.). These models were developed for SiGe epitaxial films and not for SGOI films but we will still compare Takagi et al. results and ours with these models. We suggest that the nature of the interfaces of the SiGe films, being with a Si crystal for the case of epitaxial SiGe systems while being with amorphous oxides for the case of SGOI films, influences strain relaxation of the SiGe crystal. We propose that SiGe-oxide interfaces might alter either the balance between elastic strain energy and plastic strain energy, which would explain the discrepancy with equilibrium models of epitaxial systems, or the generation, recombination and gliding of dislocations, which are known to nucleate at interfaces.

IV.A.4.1.5. Discussion about the Influence of the SGOI Thickness on Strain

Even if strain of SGOI films seem to disagree with the most accepted equilibrium models of epitaxial SiGe-Si systems, Takagi et al. still showed that the evolution of strain is not dependent on the Ge concentration only, as could imply Figure IV.12.. Takagi et al. discussed the effect of the thickness of SGOI films on strain using two key methods:

Chapter IV -Strain and Crystal Quality of SGOI Films Formed by the Condensation Technique 151/230 ȃ They considered that the thickness of initial epitaxial SiGe films and SGOI films obtained by condensation of these initial films can be represented by the thickness of their equivalent GOI film. The equivalent GOI film thickness of an initial epitaxial SiGe film is the thickness of the GOI film obtained by condensation of this film until 100% Ge concentration is reached. For instance, a 40 nm-thick epitaxial Si 0.72 Ge 0.28 film corresponds to a 11.2 nm-thick GOI film after condensation (see Section IV.A.1. for calculations) while a 40 nm-thick epitaxial Si 0.78 Ge 0.22 film corresponds to a 8.8 nm-thick GOI film after condensation. We will also use this analysis methodology to discuss the effect of the thickness of SGOI films on strain.

ȃ They studied SGOI films fabricated on tensily-strained SOI substrates (sSGOI). Tensily-strained SOI substrates (sSOI) are substrates whose SOI film is tensily strained with respect to relaxed Si.

Firstly, the comparison of strain of SGOI and sSGOI films made by Takagi et al. obviously shows that strain relaxation of SGOI layers is not dependent on the Ge concentration only. Indeed, the authors showed that the "bell shape" of the evolution of strain of sSGOI films is very similar to the one of condensation experiments performed with relaxed SOI substrates, except that it is shifted to higher Ge concentrations. We highlight that the authors found that the maximum strain values of SGOI and sSGOI films are similar (between 2.2% and 2.5%).

Secondly, as illustrated on Figure IV.14, Takagi et al. found that thicker sSGOI films yielded lower compressive strain at identical Ge concentrations. We also notice on Figure IV.14 that the effect of the sSGOI thickness on elastic strain seems to be less and less important as sGOI thicknesses are reduced. If we assume that this tendency is also observed for SGOI films, we can suggest that as the GOI equivalent thickness is reduced to lower and lower values, the evolution of strain as a function of the Ge concentration gets closer to some sort of "asymptotic bell shape" curve with a maximum strain value between 2.2% and 2.5%. It also corresponds to a decreasing influence of the SGOI thickness on strain as this thickness is reduced. If we now look back to Figure IV.11, the fact that our results are aligned with Takagi et al. in terms of evolution of strain as a function of the Ge concentration is consistent with the idea of an "asymptotic bell shape". Indeed, data of Takagi et al. showed Figure IV.11 corresponds to their thinner SGOI films and SGOI films of this study are even thinner. Takagi et al. [Takagi10]. sGOI equivalent thicknesses are indicated in the legend (see text for the definition of the sGOI equivalent thickness).

Figure IV.14: Compressive in-plane strain values of sSGOI films published by

IV.A.4.1.6. Conclusion about Strain of SGOI Films Fabricated by High Temperature Condensation

In summary, the comparison of our results with Takagi et al. ones and with equilibrium models of epitaxial SiGe-Si models showed that SGOI strain relaxation is not governed by equilibrium models of epitaxial SiGe-Si systems. We argued that the Ge concentration alone is not a relevant parameter and that the SGOI thickness influences strain of SGOI films. However, it seems that the effect of the thickness of the SGOI film on strain relaxation is likely weak for SGOI thicknesses in the 10 nm range, a range which is relevant in light of today's technology needs. We finally highlight that the strain level is likely the key parameter of strain relaxation of thin SGOI films, with a threshold in the 2.2%-2.5% range for an optimized condensation process based high temperature oxidation. In terms of strain relaxation mechanisms, we highlighted that it is likely highly dependent on the nature of the SiGeoxide interfaces.

As elastic strain relaxation is usually accompanied by generation of crystal defects such as dislocations, we study into the next section the formation of dislocations in SGOI films fabricated by condensation based on dry RTO.

IV.A.4.2. Strain of SGOI Films Fabricated by Low Temperature Condensation

As emphasized into the last section, SGOI strain relaxation is not governed by equilibrium models of epitaxial SiGe-Si systems and the nature of the interfaces with the SiGe film likely influences the mechanisms of strain relaxation. We performed condensation experiments at low temperature by means of ISSG oxidation at 850°C in order to study the influence of the temperature and of the nature of the interface with the SiGe film on strain relaxation. The process is schematically illustrated on Figure IV.15. As shown in Section III. A.4.3., ISSG oxidation at 850°C of a Si 0.7 Ge 0.3 film leads to formation of a pile-up layer with a Ge concentration of about 60%. Then, as condensation proceeds, the thickness of the pile-up layer increases while its Ge concentration is rather constant. The oxidation temperature being low, there is hardly any interdiffusion between the epitaxial layer and the SOI film. Finally, after the pile-up layer has "consumed" all the epitaxial SiGe layer, it starts to extend into the SOI layer until the bottom interface of the pile-up layer reaches the interface with the BOx layer. Oxidation durations were chosen based on kinetics results of Section III.A.4.1. so that it stops when the bottom interface of the pile-up layer reaches the interface with the BOx layer. As recipes are based on results of ISSG oxidation kinetics of SiGe layers on bulk Si substrates, there is some uncertainty on condensation kinetics (see section IV. A.2.1.). Two oxidation durations were thus chosen and final Ge concentrations were 53% and 58% based on measurement of the SGOI layer thicknesses by XRR.

As shown Figure IV.16, Both samples have strain levels well below the ones of fully strained SiGe layers with the same SiGe compositions and obtained with high temperature oxidation. It is suggested that a SiGe layer on SOI has a crystalline interface that might favor strain relaxation via dislocations formation (a case that could be comparable to an epitaxial SiGe film on a Si bulk substrate). While for high temperature condensation, interdiffusion is fast enough for a quite homogeneous SGOI film to be formed quickly after the beginning of condensation. The interface between the strained SiGe layer and the underlying film is not a crystalline one anymore, but an interface with an amorphous oxide that might hinder strain relaxation up to higher strain levels.

IV.A.4.3. Extended Defects Revelation

AFM images of the surface of SGOI layers after oxide removal are shown on Figure IV.17. No specific pattern is observed for Ge concentrations up to 57% and Root Mean Squared (RMS, see Appendice E) are between 1.5 Å and 2 Å, which are typical values of SOI substrates. However, characteristic cross-hatched patterns are observed for SGOI63 and SGOI77 films. These patterns are formed by lines oriented in the <110> direction, and the spacing between lines is estimated to about 50 nm. These lines are due to atomic steps on the surface of the SiGe layer induced by dislocations gliding on {111} planes [Gunji11]. Strain relaxation of SGOI63 and SGOI77 samples measured by either Raman spectroscopy or TEM GPA (see Section IV. A.4.1.1.) is thus accompanied by generation of dislocations on {111} planes.

<110> (001)

We also acquired TEM images of some samples of this study to have a look at the films fabricated by condensation (and also to measure strain of SGOI films by GPA as shown in Section IV. A.4.1.1.). Indeed, while AFM is a convenient technique to identify the presence of dislocations at the surface of SGOI films, TEM is required to study the bulk of the SGOI films. These TEM images can be also used to obtain complementary information on the dislocations. Note that comparison between AFM and TEM results are not easy due of the different scales and views (edge view for the TEM section and plane view for AFM). From TEM images, dislocations can be evidenced using the Moiré technique. Numerical Moiré patterns of a chosen plane, such as (111), can be obtained by Fast Fourier Transform (FFT) of the HRTEM image, followed by filtering around the chosen diffraction spot, and finally inverse FFT to reconstruct the chosen planes. Figure IV.18,IV.19,IV.20 and IV.21 show HRTEM images and the (111) numerical Moiré patterns of Reference, SGOI25, SGOI57 and SGOI77 samples. No dislocations are observed in numerical Moiré patterns of Reference and SGOI25 samples while some are for SGOI57 and SGOI77 ones. TEM lamella preparation may be the source of some dislocations so that we cannot claim for sure that the dislocations observed in HRTEM images and corresponding Moiré patterns are caused by the condensation process. However, TEM lamella preparation cannot make dislocations disappear so that we can claim that there is no dislocations in the bulk of the films of Reference and SGOI25 samples in the limited probed area of a TEM lamella. Interestingly, dislocations were also observed in the Si substrate of the SGOI77 sample. If these dislocation are due to the condensation process and not to the TEM lamella preparation, we suggest that the substrate might be slightly strained by the SGOI layer even with the 20 nm-thick buried oxide film between the two layers. The Si substrate would be tensily strained and not compressively strained in this case. Properties of tensile Si might be a lot different from the ones of compressive Si, and might explain why there are dislocations into the substrate for low tensile strain levels. However, an in-depth study, based on dedicated experiments and mechanical simulations, seems required to further the discussion on these dislocations. In summary, we showed that elastic strain relaxation of SGOI films measured by Raman spectroscopy is accompanied by formation of dislocations on {111} planes. Identifying the precise nature of these dislocations is however more difficult as TEM images are 2D projections of 3D samples, and as a consequence 2D projections of 3D dislocations. The absence of a sound consensus in the literature about the type of defects in partially relaxed SGOI layers also confirms the difficulty to study the nature of dislocations

[Bedell04][Nakaharai06][Vincent07] [Gunji11].

IV.A.4.4. Conclusion of Section IV.A.4.

We showed that strain of SGOI films fabricated by condensation based on high temperature dry RTO increases with the Ge concentration until a strain peak value of about 2.1% at 60% Ge concentration. It was also evidenced that the strain level is likely the key parameter of strain relaxation of thin SGOI films and not the Ge concentration, with a threshold in the 2.2%-2.5% range for an optimized condensation process based on dry oxidation at high temperature. In terms of strain relaxation mechanisms, we emphasized that the oxide-SiGe interfaces likely play a role in strain relaxation by hindering dislocations formation up to higher elastic strain values than at SiGe-Si interfaces. Finally, AFM imaging of the surface of SGOI films revealed that strain relaxation is accompanied by formation of dislocations in {111} planes.

IV.A.5. Characterization of SGOI Composition and Crystal Quality by MEIS

The former section focused on the evolution of elastic strain of SGOI films and its correlation with the formation of dislocations. The process conditions and mechanisms that lead to strain relaxation and formation of dislocations have been discussed. However, a finer characterization of the crystal quality as a function of depth into the SGOI film would be of great interest as the SiGe crystal quality influences its electrical properties. This is the object of this section. The composition and crystal quality of SGOI films fabricated by condensation is studied by MEIS. The three studied samples and experimental conditions of MEIS experiments are firstly introduced. In particular, the analysis methodology of MEIS random and channeling spectra is presented. The composition depth profiles are firstly studied because they must be fixed for the analysis of channeling spectra that allows to study the crystal quality of SGOI films. This section ends with the evolution during condensation of the crystal quality of SGOI films as a function of depth.

Sample name

"Reference" "Low Ge" "High Ge"

Random 45-90

Channeling 45-90

Random 66-90 MEIS spectra were acquired on the three different samples and in the three angular configurations shown Table IV.1. These three samples are not the same as the ones studied in Section IV.A.4. (except for one) and will thus be called differently. The three samples are: ȃ "Reference": it corresponds to a sample after SiGe epitaxy, so before condensation. ȃ "Low Ge": this SGOI sample has been fabricated by a condensation process close to the one used for fabrication of standard SGOI layers at STMicroelectronics Crolles fab (dry RTO at 1100°C). The 25% Ge concentration is also in the range of standard Ge concentrations of SGOI layers. ȃ "High Ge": it is the SGOI layer formed by condensation based on dry RTO that features the highest Ge concentration without any visible defects on AFM images (see Figure IV.14) and with a strain level very close to the one of a fully strained SiGe layer (see Figure IV.10). This sample is the SGOI57 film already studied in Section IV.A.4..

For both "Low Ge" and "High Ge" samples, the thermal oxide grown during condensation was thinned down by HF treatment in order to decrease the "de-channeling" effect of amorphous layers located on top of multilayer systems (see Section II.4.3.4. for more details about de-channeling).

The MEIS angular configurations used in this study are (details about "random" and "channeling" modes are given Section II.4.3.4.):

ȃ "Random 45-90": 45 corresponds to 45° between the incident ion beam and the normal to the surface of the sample, and 90° is the scattering angle. This configuration corresponds to the random mode. ȃ "Channeling 45-90": The angle configuration is similar to the one of "Random 45-90" but in channeling conditions. It corresponds to channeling in the [111] direction of the Si/SiGe crystal. ȃ "Random 66-76": the incident angle with respect to the normal of the surface of the sample is 66°.

Channeling conditions were chosen to be the ones that feature the smallest Ge peak.

Table IV.1: MEIS acquisition conditions used for each sample.

IV. A.5.1.1. Studied Samples and Experimental Configurations

IV.A.5.1. Experimental

Section II.4.3. introduced the principles of the MEIS technique and discussed the dependence of the energy and of the probability of ion backscattering on the nature and depth of the scattering atom. The analysis of experimental MEIS spectra that is performed into the next sections requires the identification of the respective contributions of each element of each layer to the MEIS spectra for the specific samples of this study. This identification is conducted by simulation of each one of these contributions.

Figure IV

.20 shows a simulated MEIS spectra of a thermal oxide-SGOI stack where the contribution of each element of each layer is shown. As explained in Section II.4.3.3., the higher the atomic number of the scattering IV. A.5.1.2

. Identification of Peaks on MEIS Spectra

element of each layer is shown. As explained in Section II.4.3.3., the higher the atomic number of the scattering element is, the higher the energy of backscattered ions is. We thus observe that the Ge peak of the SiGe layer is at a higher energy than the Si one of the same SiGe layer. We also notice that the peak area of Ge is higher than the one of Si while the simulation was performed for a Si 0.75 Ge 0.25 layer which has more Si atoms than Ge atoms. Indeed, the probability of backscattering also increases with the atomic number, which makes the number of ions backscattered by Ge atoms higher than the one backscattered by Si atoms. These two effects are also evidenced when comparing the Si peak of the thermal oxide with the O peak of the same layer.

The dependence of the energy of backscattered ions on the depth at which they are backscattered is also observed. Ions backscattered by Ge atoms located at the top of the SiGe layer have the highest energy (around 92 keV for the example of Figure IV.20) while ions backscattered by Ge atoms at the bottom of the layer escape the top surface with a lower energy (around 85 keV for the example of Figure IV.20). The order of magnitude of the thickness dependence of the energy of backscattered ions is 1 keV for 1 nm. There are Si atoms in all layers of the simulated multilayer system so that Si peaks of each layer are consecutive from about 81 keV. We can also notice that the O peak of the thermal oxide layer, located at around 66 keV, is at the same energy as the Si peak of the bottom of the BOx layer. As a conclusion, the dependence of the energy of backscattered ions and of the scattering probability on both the atomic number of the scattering element and the thickness at which ions are backscattered is clearly evidenced for samples and MEIS conditions used in this study. This identification allows us to analyse the experimental 2D MEIS maps presented into the next section.

Experimental MEIS spectra at a specific scattering angle were extracted from 2D maps such as the one shown

IV.A.5.2. Experimental and Simulated Spectra

Experimental and simulated spectra are shown Figure IV.22. Due to the experimental constraints described in Section II.4.3., a trade-off between the scanned energy range and the counting time had to be chosen. Hence the scanned energy range is different from one spectrum to another depending on this trade-off. For "Reference 45-90" and "Low Ge 45-90" random and channeling spectra and "High Ge 45-90" channeling spectra, the scanned energy range was fixed to 65-92 keV. These spectra show the Ge peak and the Si peaks of the thermal oxide, of the SiGe layer and of a part of the BOx layer. While we can observe the O peak at about 67 keV for the "Low Ge 66-76" random spectrum. As for the "High Ge" random spectrum, the scanned energy range was 55-92 keV and thus allows the O peak at about 57 keV to be visible. All simulated spectra were simulated according to the procedure described in Section II.4.3.6..

Figure IV.22:

Experimental MEIS spectra and simulated ones for all configurations described in Section IV. A.5.1.1.: (a) "Reference" in random and channeling 45-90 configurations, (b) "Low Ge" in random and channeling 45-90 configurations, (c) "Low Ge" in random 66-76 configuration, and (d) "High Ge" in random and channeling 45-90 configurations.

Random spectra were firstly analyzed in order to retrieve the composition as a function of depth, which is the object of the next section. Then, the comparison with channeling spectra allows us to focus on the crystal quality in Section IV. A.5.4.. This section discusses the chemical composition profiles of the three studied samples. Fixing the composition of the samples is indeed required for analysing the channeling spectra in Section IV. A.5.4. (see Section II.4.3.6.).

The composition as a function of depth was extracted from MEIS spectra acquired in random conditions. A limited number of layers was used to fit experimental spectra so that composition gradients cannot be precisely retrieved. The composition of the top oxide and of the BOx layer were fixed to SiO 2 while the one of the SOI layer was fixed to pure Si. The Ge concentrations as a function of depth, taking the BOx-SGOI interface as zero depth and the x axis towards the surface of the sample, are plotted Figure IV.23. The BOx-SGOI interface was chosen to be the zero reference for the sake of consistency with the crystal quality results presented in Section IV.A.5.4.. While a single homogeneous SiGe layer was used for fitting ellipsometry and XRR spectra of the "Reference" sample, two SiGe layers were used for fitting the MEIS random spectrum. Due to a limited number of layers used for fitting MEIS spectra (see Section II.4.3.6.), the inhomogeneity of the layer is represented by two SiGe layers with different concentrations. This SiGe bilayer is most certainly not real and a smooth gradient is more likely. Similarly, the values of the Ge concentrations of the two layers are not to be taken as absolute values but as average concentrations over the thickness of each of the two SiGe layers used for fitting.

As well as for the "Reference" sample, we used two SiGe layers to fit MEIS spectra in random 45-90 and 66-76 configurations for the "Low Ge" sample. However, even if this two SiGe layers system gives better fits than a single SiGe layer, the bottom SiGe layer is very thin and its composition is very close to the one of the top SiGe layer.

IV.A.5.3. Chemical Composition Analysis

As a consequence, it seems that the SiGe layer is fairly homogeneous. Such a homogeneous SGOI layer was expected because of the high temperature used for the condensation process (which tends to homogenize the SGOI layer as seen in Section III. A.4.3.).

Three SiGe layers were introduced to fit the experimental random spectrum of the "High Ge" sample. The top of the layer is found to feature a higher Ge concentration than the bottom (from 62% Ge concentration on top to 40% at the bottom). Again, a smoother composition gradient is more likely than three well-defined and homogeneous SiGe layers. We observe that the top of the SGOI layer seems quite homogeneous with a Ge concentration close to 62% while the Ge concentration over the deepest 3 nm of the SGOI layer is significantly lower. Contrary to the "Low Ge" sample, the composition inhomogeneity is significant. We propose that the SGOI layer was most likely homogeneous as long as the condensation temperature was 1100°C (until about 40% Ge concentration). Then, as the temperature was decreased to 1000°C, the pile-up effect was more pronounced and led to higher Ge concentration below the oxide. At the same time, SiGe interdiffusion drastically increases above 50% Ge concentration [Kube10][David15], which may lead to stabilization of the Ge concentration of the pile-up layer while interdiffusion is slowed down at the bottom of the layer. Hence the quite homogeneous SiGe layer on top (here with 62% Ge concentration) and the lower Ge concentration at the bottom (probably as low as 40%).

The chemical profiles of all three samples retrieved from random spectra were fixed for the analysis of channeling spectra performed into the next section.

IV.A.5.4. Crystal Quality Analysis

IV.A.5.4.1. Introduction and χ min Definition

As described in Section II.4.2.4., channeling is a mechanism that takes place when the ion beam is aligned with a crystallographic direction of the probed crystal. We evidenced in the previous section that there are Ge concentration gradients from the top to the bottom of SGOI films. These gradients induce out-of-plane lattice parameter changes from the top to the bottom of SGOI films. It results in deviations of crystallographic directions that must be evaluated for a rigorous analysis of MEIS channeling spectra, which is the subject of Section IV. A.5.4.2.. Then,in Section IV.A.5.4.3., we evaluate the surface peak caused by backscattering by the first atomic layers of the SGOI film, which are not included in the shadow cone (see Section II.4.3.4.).

Having assessed the two former effects that affect the channeling spectra, Section IV.5.4.4. discusses the depth profiles of the χ min factor. As already introduced in Section II. 4.3.4., the χ min factor is the percentage of atoms displaced from theoretical lattice sites. In the case of SiGe films, the theoretical lattice sites are the ones taking into account the Poisson effect.

IV.A.5.4.2. Channeling Angle and Angular Tolerance

The channeling angle along the [111] direction of an unstrained Si crystal is 45°. However, for the case of a fully strained SiGe film, the [111] channeling direction is disoriented with respect to 45° due to the Poisson effect. Channeling may still be observed if the incident beam is only slightly disoriented with respect to the considered crystal direction, here the [111] one. The angular tolerance is defined as the angular difference between the angle that gives the lowest χ min factor (in other words the best channeling angular conditions) and the one that corresponds to an increase of the χ min of 50% of the gap between the lowest χ min factor and the one of random conditions. The angular tolerance is [START_REF] Kahng | Electric field controlled semiconductor device[END_REF]:

ψ πε 1 2 2 0 2 = Z Z e Ed I A
, where Z I is the atomic number of the incident ion, Z A is the average atomic number of the considered material, e is the elementary charge, ε 0 is the dielectric constant of vacuum, E is the kinetic energy of incident ions and d is the average distance between atoms along atomic rows. Let us consider the case of a fully strained SiGe layer, which features the larger angular deviation with respect to unstrained Si. Figure IV.24 shows the angular deviation with respect to 45° of the channeling angle along the [111] direction and the angular tolerance as a function of the Ge concentration. The angular tolerance and the angular deviation increase with the Ge concentration because of the increase of the average atomic number and because of the Poisson effect respectively. For Ge concentrations between 0% and 60%, the angular tolerance is always at least three times (which is the case of 60% Ge concentration) as large as the angular deviation. Chapter

Equation

. The Shadowing Effect and the Surface Peak

As briefly described in Section II.4.3.4., in channeling conditions, deflection of incoming ions by the first nuclei of atomic rows is at the origin of a shadow cone (the so-called shadowing effect) that prevents ions from being backscattered by deeper atoms. The scattering yield of the first atomic layers is however the one of an amorphous layer as scattering of incoming ions is not prevented by this shadowing effect. It results in a peak, called surface peak, in spectra acquired in channeling conditions whose contribution must be quantified for a rigorous analysis.

The number of atoms per row to participate to the surface peak has been measured and simulated for various crystals, channeling directions and temperatures and was found to follow a trendline that only depends on the ratio ρ/R m , where ρ is the amplitude of the thermal vibrations perpendicular to the channeling direction and R m is the radius of the shadow cone [START_REF] Kahng | Electric field controlled semiconductor device[END_REF]. We calculated the radius of the shadow cone for a fully strained SiGe crystal as shown in [START_REF] Kahng | Electric field controlled semiconductor device[END_REF] and found that it is about 0.20 Å and rather constant over the full range of Ge concentration. Using the amplitude of thermal vibrations at room temperature of a Si crystal perpendicular to the [110] direction (we could not find values of amplitudes of thermal vibrations of a SiGe crystal in the perpendicular direction to [111], so that we assume that it is close enough to the case of a Si crystal in [110] channeling conditions to use the latest value), we can show that the ratio ρ/R m is about 0.39. It corresponds to only one atom per atomic row to participate to the surface peak whatever the Ge concentration is.

In summary, with the conditions used for MEIS measurements, the surface peak of SGOI films is limited to backscattering of the first atomic layer of the film.

IV. A.5.4.4. Crystal Quality of SGOI Films Figure IV.25 shows the Ge concentration and χ min factor profiles extracted from random and channeling MEIS spectra. Some caution must be observed when discussing the χ min factor profiles. Indeed, as we could not explore all possible solutions due to the manual fitting method, we cannot guarantee that the solutions reported here are the best ones and the closest to reality. Indeed, the contributions of each atomic layer overlap with each other because of the straggling effect (see Figure IV.20 which shows that peaks overlap with each other). As a consequence, determination of the best χ min factor profile is more difficult by a manual fitting method. If we look for instance at the "Reference" sample, we found that the top 2 nm of the SiGe films has a χ min factor of 75 %. But the best solution could also be that the top 1.6 nm of the SiGe layer has a χ min factor of 85 % as far as we cannot compare all χ min factor profiles.

However, what we are more confident about is the total amount of atoms displaced from their lattice positions because this amount is more directly linked to the total area of the peak. For the sake of a more quantitative comparison between crystal qualities of samples, we define a term called Amorphous Atomic Layer Equivalent Number (AmALEN) which is proportional to the total amount of displaced atoms of the studied layer:

AmALEN th d = × χ min ,
where th is the thickness of the considered layer and d the distance between two atomic layers (in this case between two {111} planes because the chosen channeling direction is [111]). When the considered layer is the top part of the SiGe film, the AmALEN of the surface peak (see Section IV. A.5.4.3.), which is equal to one, must be subtracted to the total AmALEN of the film.

Let us comment each sample:

Equation IV.5

As a result, when the angle is fixed so that the χ min factor is minimum in the SiGe layer of the "Reference" sample, the channeling conditions are still very close to the best ones for the SOI layer. And as for the "Low Ge" and "High Ge" samples, non-homogeneities of the Ge concentration are neither high enough to induce important variations of the χ min factor. In summary, the χ min factor is weakly sensitive to the elastic strain level variations of this study. And as a result, only atoms displaced from lattice sites in a significant manner, such as atoms of a dislocation core and in the immediate vicinity of it, or atoms in interstitial positions, significantly influence the χ min factor.

ȃ Figure IV.25 (a): "Reference": SiGe epitaxial layer on SOI: the χ min factor is found to be 5% in the whole SOI layer. It thus shows that the SOI crystal is not altered by the SiGe epitaxy process and still has a very low amount of atoms displaced from lattice sites. The bottom part of the epitaxial SiGe layer has a higher amount of atoms displaced from lattices than the SOI layer, even if it can be considered a high-quality crystal with a χ min factor of 10%. Finally, the top part has an important amount of atoms displaced from lattice sites, with an AmALEN of about 3. The fact that the χ min factor of the SOI layer, which is below the SiGe film, is as low as 5% rules out any disorientation with respect to channeling conditions as well as de-channeling of the ion beam through the top amorphous oxide as the possible explanations of such a high AmALEN. Neither roughness can account for such a high AmALEN because roughness only is not associated with SiGe crystal distortion. Roughness would result in the surface peak being "spread" over a wider energy range. Buckling, however, can be seen as roughness associated with crystal distortion and can contribute to such a high AmALEN.

ȃ Figure IV.25 (b): "Low Ge": Low Ge fraction SGOI films: the SiGe crystal sandwiched between the perturbed top part and the 1 nm-thick bottom part of the SGOI layer is as perfect as the SOI crystal. It shows that SiGe interdiffusion does not damage the crystal quality of the SOI layer during condensation.

The AmALEN of the top part of the SiGe layer decreases from 3 for the "Reference" sample to 1. Two mechanisms are proposed to account for this decrease. Firstly, thermal oxidation may contribute to reorganize atoms of the top part of the SiGe crystal so that there are less atoms displaced from lattice sites after condensation. And secondly, the top part of the SGOI layer of the "Low Ge" sample corresponds to the former SOI film, whose χ min factor is likely not altered by SiGe interdiffusion as suggested above.

As for the bottom part of the SGOI layer, it seems that the presence of Ge at the interface with the BOx layer is associated with a degradation of the crystal in the immediate vicinity of the interface. Some authors showed that the higher the Ge concentration at an interface with a thermal SiO 2 film is, the higher the number of Ge dangling bonds is. They argued that these dangling bonds are formed in order to accommodate the stress at the semiconductor-oxide interface [Houssa11]. We suggest that while the crystal is not stressed at the interface with the BOx layer for the case of the SOI-BOx interface, it is stressed for the one of the SGOI-BOx interface, and that this stress may cause crystal distortion and/or damage. This mechanism could also account for at least a part of atoms displaced from lattices sites in the top part of the SGOI film.

ȃ Figure IV.25 (c): "High Ge": High Ge fraction SGOI films: The AmALEN of the top part of the SGOI layer is about 2. The fact that it is lower than the one of the "Reference" sample suggests that thermal oxidation tends to decrease the χ min factor by reorganizing atoms below the oxidation front. Indeed, here the SGOI is thicker than the initial SOI so that the second mechanism proposed above for the case of the "Low Ge" sample can be ruled out in this case. The χ min factor is also found to be higher close to the interface with the BOx layer. Finally, apart from the areas close to interfaces, the χ min factor increases from 5% for low Ge fraction to about 10%, which shows that the SiGe crystal starts to be slightly degraded, possibly because of dislocations such as those evidenced by TEM GPA in Section IV. A.4.1.3.. Having discussed in this section the crystal quality of the three studied samples individually, the key evolutions of the crystal quality of SGOI films during condensation are summarized in a concise and schematic manner into the next section. We firstly discussed the composition profiles retrieved by the analysis of MEIS spectra. The composition of the SiGe films was found to be rather homogeneous. Then, the crystal quality of SGOI films, represented by χ min factor which is the percentage of atoms displaced from lattice sites at a certain depth, was studied. Based on a schematic illustration,

Reference

Crystal quality

Distance from BOx towards the surface (a.u.)

Perturbation of the top part of SiGe crystal

• The SOI crystal quality is not altered by the SiGe epitaxy process.

• The SiGe crystal is degraded in the top 1-2 nm of the epitaxial SiGe layer.

No perturbation of the crystal during SiGe interdiffusion in the SOI film

After condensation -Ge% < 50%

Crystal quality

Distance from BOx towards the surface (a.u.) • Thermal oxidation may contribute to atomic rearrangement below the oxidation front and may tend to decrease the χ min factor below it.

Perturbation of

• The SOI crystal quality is not altered by SiGe diffusion for low Ge fraction SGOI layers.

• The presence of Ge, even at low fraction, seems to induce some crystal damage in the immediate vicinity of the SGOI-BOx interface.

• When the Ge concentration increases, the SiGe crystal quality starts to degrade, even far from interfaces with oxides.

SiGe Si

IV.A.6. Conclusion of Section IV.A.

We showed that strain of SGOI films fabricated by condensation based on high temperature dry RTO increases with the Ge concentration until a peak value of about 2.1% at 60% Ge concentration. We showed that the parameter that has the strongest influence on strain relaxation is the strain level and not the total elastic strain energy nor the Ge concentration and that strain relaxation is accompanied by formation of dislocations in {111} planes. Using Takagi et al. data for comparison [Takagi10], mechanisms of strain relaxation were discussed, and we emphasized that that the oxide-SiGe interfaces likely play a role in strain relaxation by hindering dislocations formation up to higher elastic strain values than at SiGe-Si interfaces. Low temperature condensation by means of ISSG oxidation was also performed but found to be not adequate for fabricating high Ge content fully strained SGOI films at the moment.

A finer characterization of the crystal quality as a function of depth into the SGOI film was performed by MEIS using the channeling effect. The evolution of the crystal quality of SGOI films during condensation is summarized in a few key points: ȃ The epitaxial SiGe crystal is already degraded in the top 1-2 nm. ȃ Thermal oxidation may contribute to atomic rearrangement below the oxidation front and may tend to decrease the χ min factor below it. ȃ The SOI crystal quality is not altered by SiGe interdiffusion for low Ge fraction SGOI layers. ȃ After condensation, the presence of Ge, even at low fraction, seems to induce some crystal damage in the immediate vicinity of the SGOI-BOx interface. ȃ At high Ge concentrations (> 55%), the SiGe crystal quality starts to degrade, even far from interfaces with oxides.

IV.B.1. Problem Definition

Strain of SGOI channels should be uniformly maintained in all pMOS active areas during fabrication of functional CMOS devices in order to fully benefit from the gain of strain on hole mobility. However, partial relaxation of the SGOI film at the frontier between pMOS active areas and Sallow Trench Isolation (STI) areas (see Figure IV.27) has been evidenced after the STI etch process step and shown to degrade electrical performances of transistors fabricated close to these frontiers [Andrieu16][Berthelon16]. It is crucial to understand and prevent as much as possible these effects as they induce a transistor performance dependency on chip design. Boureau showed that introducing a "sliding" mechanism at the thermal oxide-SGOI and BOx-SGOI interfaces in FEM simulation allows to reproduce experimental results (see Figure IV.28). As pointed out by Boureau, the agreement between the experimental results and simulations with the "sliding" mechanism suggests a special influence of the mechanical properties of the thermal oxide-SGOI and BOx-SiGe interfaces, even if it does not bring a physical meaning to the "sliding" mechanism. Usuda et al. also suggested that the BOx-SiGe interface may have mechanical properties that differ from the ones of the elastic theory.

IV.B. Strain Relaxation at the Edge of Active Areas

Figure IV.28: In-plane deformation measured by DFEH along with FEM simulations based on the elastic theory only, reffered as "classic FEM", and on the elastic theory with a sliding mechanism introduced at the SGOI-thermal oxide and SGOI-BOx interfaces, referred to as "FEM with the sliding mechanism". The standard condensation process was performed at 1100°C. Adapted from [START_REF] Cooper | Strain mapping of semiconductor specimens with nm-scale resolution in a transmission electron microscope[END_REF].

DFEH Classic FEM FEM with the sliding mechanism

The condensation process allows fabrication of the SGOI film before the STI etch process step, and may alter the mechanical properties of the interfaces of the SGOI film with the thermal oxide and the BOx film. In Section IV.B., we thus try to correlate the strain relaxation experimental results with the condensation process conditions and with the properties of the SGOI film and of its interfaces with the thermal oxide and the BOx layer. We identified two possible causes for this strain relaxation phenomenon and performed dedicated experiments to study them. They are presented in Section IV. B.2. and Section IV.B.3 . The presence of Ge-O bonds could induce weaker mechanical properties at the SiGe-oxide interface and account for the "sliding" mechanism suggested by Boureau [START_REF] Cooper | Strain mapping of semiconductor specimens with nm-scale resolution in a transmission electron microscope[END_REF]. No Ge-O bonds were evidenced at the SiGe-thermal oxide interface (see Section III.A.3.). However, the SiGe-BOx interface may be different from the top SiGe-thermal oxide interface because of different formation processes. The SGOI-thermal oxide one is the oxidation interface while the SGOI-BOx one is modified by diffusion of Ge atoms in its vicinity. We thus decided to investigate the chemical bonding properties of the SiGe-BOx interface by XPS. This study required a specific specimen preparation which is detailed into the next section.

IV.B.2.2. Experimental

Three specimens were studied: ȃ "Reference": a SOI substrate was oxidized by means of dry furnace oxidation at 900°C. This oxidation process is systematically performed prior to the oxide removal-SiGe epitaxy-condensation sequence. We chose to perform this process to discriminate the effect of the condensation process on the evolution of the chemical bonding at the SGOI-BOx interface from the effects of the inital oxidation process of the SOI substrate.

ȃ "SGOI 1100°C": this sample is close to the standard condensation process and is known to suffer from the problem of strain relaxation at the edge of active areas. After epitaxial growth of SiGe on SOI, condensation was performed by means of dry RTO at 1100°C. An anneal in inert environment was performed in the same equipment at 1100°C for one minute to fully homogenize the SGOI film between the thermal oxide and the BOx layer. The final SGOI Ge concentration is about 25%.

ȃ "SGOI 900°C annealed": a SiGe film was epitaxially grown on a SOI substrate and condensation was performed by means of dry RTO at 900°C. A furnace anneal in inert atmosphere for 8 hours at 900°C was then performed to fully homogenize the SGOI layer. This sample was fabricated in order to discriminate the effect of a lower oxidation temperature on chemical bonding at the SGOI-BOx interface. SiO 2 has been shown to be subjet to viscous flow at temperatures as low as 960°C, so that we decided to perform condensation below 960°C to discriminate the effect of viscous flow on strain relaxation [EerNisse77]. The final SGOI Ge concentration is also about 25%.

Then, the specimens were prepared for XPS measurements by the sequence described in The chemical etching of Si was performed using Tetramethylammonium hydroxide (TMAH) and thinning of the BOx layer using HF solution. The HF etching process was calibrated on a dedicated sample and the residual BOx thickness was measured by ellipsometry.

IV.B.2.3. Ge-O Bonding

Figure IV.30 shows the experimental Ge3d and Ge2p XPS spectra of the two SGOI samples along with fitted peaks. During the photoelectron emission process, a part of the energy of the incoming X-Ray photons is transferred to the emitted electron as kinetic energy. The kinetic energy of electrons ejected from the Ge2p orbital is lower than the one of electrons ejected from the Ge3d spectrum (see also Section II.3.3. for a more detailed description of the XPS technique). As a consequence, the probing depth of the Ge2p spectrum is lower than the one of the Ge3d one (2.5 nm and 8.3 nm for the Ge2p and Ge3d orbitals respectively at a collection angle of 45°).

No Ge oxide peak is evidenced on any XPS spectrum except on the Ge2p spectrum of the "SGOI 1100°C" sample. The peak is very weak and no Ge oxide is evidenced on the Ge3d spectrum of the same sample. The Ge2p spectrum has a lower probing depth than the Ge3d one, which results in a lower contribution of metallic Ge atoms of the SGOI layer to the spectrum. The Ge oxide peak is therefore likely too weak to emerge from the foot of the metallic Ge peak on the Ge3d spectrum. As for the origin of this small amount of Ge oxides, we noticed that a first specimen preparation led to complete oxidation of Ge, most likely due to the chemical oxidation during BOx thinning. The small amount of Ge oxides observed on the Ge2p spectrum of the "SGOI 1100°C" sample might be due to the sample preparation.

We conclude on the absence of a noticeable quantity of Ge-O bonds at the SGOI-BOx interface for both investigated oxidation temperatures. We also seeked to compare the amount of Si sub-oxides at the SOI/SGOI-BOx interfaces, as a high amount of Si sub-oxides may be associated with a graded and/or a perturbed interface.

Experimental and fitted Si2p spectra are shown Figure IV.31. Neither precise evaluation of the quantity of Si sub-oxides nor reconstruction of depth profiles of each oxidation state is possible with a single XPS spectrum (see Section II.3.3. about the reconstruction of concentration depth profiles by Parallel Angle-Resolved XPS). Direct comparisons of peak areas is also difficult because of the varying oxide thicknesses from one sample to another one. We can reasonably assume that Si sub-oxides are sandwiched between a residual SiO 2 film and the SGOI/ SOI film. In this case, the SiO 2 layer is on top of the multilayer system and its residual thickness can be evaluated by comparing Si 4+ peak areas. The "SOI" and "SGOI 900°C annealed" have more similar residual SiO 2 thicknesses while the "SGOI 1100°C" has a lower SiO 2 residual thickness. We can reasonably give a direct comparison of the Si sub-oxides peak areas of the "SOI" and "SGOI 900°C annealed" samples. There is a higher amount of Si sub-oxides, especially Si 3+ , for the case of the "SGOI 900°C annealed" sample than for the "SOI" one.

A semi-quantitative comparison of sub-oxide quantities with the "SGOI 1100°C" is now described. The "SGOI 900°C annealed" sample has a thicker SiO 2 residual film than the "SGOI 1100°C". It results in Si metallic atoms to be deeper and to have a lower contribution to the total spectrum for the "SGOI 900°C annealed" sample. Therefore, the "SGOI 900°C annealed" peak has a higher Si 4+ peak area and a lower metallic Si peak area than the "SGOI 1100°C" sample. As a consequence, we suggest that the sum of the metallic Si peak area with the one of the Si 4+ is rather similar in all three samples. The fact that the residual SiO 2 film is thicker also induces the sub-oxide peak area to be lower. However, we believe that the ratio of the total Si sub-oxides peak area to the one of the sum of the Si 4+ and metallic Si can be considered a trend indicator of the total Si sub-oxide quantity. These ratio are 7%, 19% and 16% for the "SOI", "SGOI 900°C annealed" and "SGOI 1100°C" samples respectively. There is more Si sub- oxides in the "SGOI 900°C annealed" sample than in the "SOI" one in agreement with the first qualitative comparison made in the former paragraph. There seems to be roughly the same quantity of Si sub-oxides in the "SGOI 1100°C" sample as in the "SGOI 900°C annealed" one.

IV.B.2.4. Si-O Bonding

To summarize, the SGOI-BOx interface fabricated by condensation has more Si sub-oxides than the SOI-BOx. one.

IV.B.2.5. Conclusion of Section IV.B.2.

A dedicated specimen preparation has been developed in order to study chemical bonding at the SGOI/ SOI-BOx interface by XPS. It was firstly shown that there is no noticeable quantity of Ge-O bonds at the SGOI-BOx interface. It rules out weakened mechanical properties due to soft Ge-O bonds as a possible cause of the anomalous strain relaxation phenomenon described in Section IV.B.1.. Then, a higher quantity of Si sub-oxides has been evidenced after condensation. LeGoues et al. and Kilpatrick et al. found that the SiGe-thermal SiO 2 interface has more Si atoms in intermediate oxidation states than the Si-thermal SiO 2 interface and that the higher the Ge concentration is, the higher this amount is [LeGoues89'] [Kilpatrick08]. It shows that even if the SOI-BOx interface is not the oxidation interface, diffusion of Ge atoms in the vicinity of the BOx film also raises the amount of Si sub-oxides.

We suggest that if there are Ge atoms in the last complete crystal monolayer at the interface with the BOx film, then there are less Si atoms in the last monolayer, and as a result less bonds between Si atoms of the SiGe crystal and Si and O atoms of the BOx film. This limited amount of bonds between the Si atoms of the SiGe crystal and the BOx film seems consistent with a higher amount of Si sub-oxides, even if atomic scale models of the SiGe-oxide interface are required to further confirm these results.

We showed by MEIS that the SiGe crystal is degraded at the interface with the BOx layer (see Section IV. A.5.4.) and Afanas' ev et al. showed that there is a high amount of Ge dangling bonds at the SiGe-oxide interface [Afanas' ev09]. If they are Ge atoms in the last complete crystal monolayer at the interface with the BOx film, and if these Ge atoms are not bonded to any O atom of the BOx film as shown by XPS, Ge dangling bonds are indeed expected. As soon as they are dangling bonds in the crystal, some local bonds may be strained and atoms displaced from their expected lattice sites to compensate these dangling bonds in terms of total energy. In other words, come crystal disorder is expected in the vicinity of the interface with the BOx layer. ȃ "Reference epi": a SOI substrate was firstly oxidized by means of dry furnace oxidation at 900°C, and then a standard SiGe epitaxy process was performed (see Section II.1.).

IV

ȃ "SGOI 1100°C": this sample is the same as the one presented in Section IV. B.2.2.. This sample is close to the standard condensation process and is known to suffer from the problem of strain relaxation at the edge of active areas. After epitaxial growth of SiGe on SOI, condensation was performed by means of dry RTO at 1100°C. An anneal in inert environment was performed in the same equipment at 1100°C for one minute to fully homogenize the composition of the SGOI film between the thermal oxide and the BOx layer.

ȃ "SGOI 900°C": a SiGe film was epitaxially grown on a SOI substrate and condensation was performed by means of dry RTO at 900°C. There is no or hardly any Ge at the interface with the BOx film while there is a high Ge concentration at the oxidation interface (about 45%) due to the low oxidation temperature (see Section III.A.4.3.). This sample was chosen because it was suggested that H might accumulate at oxide-SiGe interfaces only and not at oxide-Si ones.

Each specimen was characterized with XRR to measure the mass density of each layer and with neutron reflectivity to measure the neutron scattering length density of each layer.

IV.B.3.2. Experimental

Segura-Ruiz et al. coupled XRR with neutron reflectivity to evidence the presence of hydrogen accumulation at the interface between a SiO 2 film and an amorphous carbon film [Segura15]. They showed that this H accumulation is correlated with higher local strain and hence a greater susceptibility to a mechanical weakness of the interface. As H is present in SiGe epitaxy processes (see Section II.1.), it was suggested that H might be incorporated into the SiGe-SOI-BOx stack during SiGe epitaxy and that it is accumulated at thermal oxide-SGOI and SGOI-BOx interfaces after condensation. This H accumulation could then weaken the mechanical properties of these interfaces and cause strain relaxation at the edge of active areas.

H is difficult to detect. Especially, H is too light to significantly influence the mass density of SiGe and SiO 2 films, so that it cannot be detected by XRR experiments. H background levels in SIMS measurements are too high to measure H concentrations below 10 19 atoms.cm -3 , which is much higher than what can be expected in thin semiconductor films. However, as explained in Section II.4.2.2., as opposed to O, Si and Ge elements, H has a strong negative neutron scattering length, which makes neutron reflectivity experiments very sensitive to its presence. As will be detailed in Section IV. B.3.3., H concentration cannot be determined based solely on neutron reflectivity experiments, but must be coupled with at least another technique. We coupled neutron reflectivity experiments with XRR to measure H concentrations in three representative samples.

Experimental details on samples, neutron reflectivity and XRR are presented into the next section and results are shown and discussed in Section IV. B.3.3.. Figure IV.33 shows experimental and fitted XRR and neutron reflectivity spectra of all samples. While the composition of oxides was assumed to be SiO 2 for fitting XRR spectra, fitting neutron reflectivity spectra with neutron scattering length densities does not require any assumption on the composition of the films. For neutron reflectivity, we only assumed a limited number of homogeneous layers. High fit qualities are obtained using the following models:

ȃ "Reference epi": native oxide (assumed to be SiO 2 ) on top of a homogeneous SiGe film on a SOI substrate. A four layers model was used for neutrons (corresponding to a SiGe-SOI-SiO 2 -Si substrate stack ; we indeed noticed that neutron reflectivity spectra are not sensitive to layers as thin as native oxides).

ȃ "SGOI 1100°C": SiO 2 oxide on top of a SGOI film on the BOx layer. A four layers model was used for neutron reflectivity corresponding to a SiO 2 -SGOI-SiO 2 -Si substrate stack.

ȃ "SGOI 900°C": a SiO 2 -SiGe-Si-BOx-Si substrate model was used. A five layers model was used for neutron reflectivity corresponding to the same SiO 2 -SiGe-Si-BOx-Si substrate stack. Mass densities determined by fitting XRR data and neutron scattering length densities determined by fitting neutron reflectivity data are shown Table IV On the first approach, we will assume that there is no H in any sample and that all layers are either Si 1-x Ge x or SiO 2 films, and see if measured neutron scattering length densities agree with predictions. The neutron scattering length densities of Si 1-x Ge x or SiO 2 are [Daillant08]:

IV.B.3.3. Results

Si

SLD N x b xb x M xM Si Ge A Si Ge Si Ge x x 1 1 1 - = - ( ) + - ( ) + ρ SLD N b b M M SiO A Si O
Si Ge Table IV.4 compares the measured SiGe SLDs with the theoretical ones calculated according to Equation IV.6 and taking as input the measured SiGe mass density. The Ge content x of Equation IV.6 is calculated according to Poisson theory (see Section I.2.2.3.) and oxides are assumed to be stoichiometric SiO 2 . We observe that the measured SLD is systematically higher than the calculated one from XRR results. H has a negative scattering length so the presence of H in any of these layers could only decrease the measured SLD below the SLD calculated from XRR, as opposed to what is observed. Therefore, there is neither H after SiGe epitaxy nor after condensation.

Neutron reflectivity

Reference epi SGOI 1100°C SGOI 900°C

Scattering length density (×10 Differences between measured and calculated SLDs are indicated in Table IV.5. Interestingly, we observe that the differences are similar between the three different samples for the same layer. It is quite low for the BOx film (between 0.01×10 -6 Å -2 and 0.06×10 -6 Å -2 ) while it is about 0.22×10 -6 Å -2 and 0.25×10 -6 Å -2 for the thermal oxide and the Si layer respectively. To explain this discrepancy, it was suggested that there might be another element at low concentrations in the SiGe film which has a scattering length high enough to significantly raise the scattering length density. Among the elements that have high positive scattering lengths (see Section II.4.2.2.) and that are present in processes performed to fabricate the samples, O and Cl are the only candidates. Cl is used in SiGe epitaxy processes (see Section II.1.), and there is always a more or less important O residual partial pressure in all semiconductor processes. Figure IV.34 shows ToF-SIMS measurements of O and Cl concentrations in the "Reference epi" sample. The concentrations of both elements in the SiGe films are below the detection limit (< 0.1% concentration) hence well below levels that could raise the scattering length density of more than 0.1×10 -6 Å -2 .

Table IV.5: Difference betwenn measured SLD and SLD calculated from XRR results.

IV.B.3.4. Conclusion of Section IV.B.3.

We performed neutron reflectivity coupled with XRR to measure H concentrations in three representative samples. We showed that there is neither H after SiGe epitaxy on SOI substrates nor after condensation. To conclude, H is not the cause of strain relaxation of SGOI films after STI etch.

We observed that the measured scattering length density of SiGe films is systematically higher than the one calculated from XRR measurements. O and Cl presence in SiGe to explain this higher scattering length density has been dismissed by ToF-SIMS measurements. Further investigation is required to understand the anomalously high scattering length density of SiGe and thermal oxide films.

Conclusion of Chapter IV

In this chapter, we firstly studied strain levels and crystal qualities of SGOI films fabricated by various condensation processes based on dry RTO and ISSG oxidation.

We firstly discussed strain levels of SGOI films fabricated by condensation based on high temperature dry RTO and ISSG oxidation. We showed that in the case of SGOI films in the 10 nm thickness range, the parameter that has the strongest influence on strain relaxation is the strain level and not the total elastic strain energy nor the Ge concentration. In terms of strain relaxation mechanisms, we suggested that if strain relaxation is governed by strain energy balance, this balance is influenced by the nature of the underlying layer: a thermal oxide (BOx) for the case of SGOI films instead of crystalline Si for the case of epitaxial SiGe-Si systems. We also suggested that dislocation nucleation at SiGe-oxide interfaces may be hindered up to higher elastic strain values than at SiGe-Si interfaces. Low temperature condensation by means of ISSG oxidation was also performed but found to be not adequate for fabricating high Ge content fully strained SGOI films at the moment.

For technology applications, we highlight that the best condensation process that we studied is dry RTO at high temperature, with a threshold value of more than 2% for SGOI films in the 10 nm thickness range. Such a strain value corresponds to a Ge concentration of about 60%.

A finer characterization of the crystal quality as a function of depth into the SGOI film was performed by MEIS. Unfortunately, we cannot discuss optimization of the condensation process in terms of SGOI crystal quality with the characterization of three samples only. However, two conclusions can be drawn: ȃ The key crystal areas that are degraded and that must be particularly looked at for optimization of the condensation process are the interfaces of SiGe with oxides.

ȃ The SOI crystal quality is not altered by SiGe interdiffusion for low Ge fraction SGOI layers, except when the Ge concentration increases above about 55%.

In the second part of this chapter, it was shown that the SGOI film is partially relaxed after the STI etch process step at the frontier between pMOS active areas and STI areas. Two possible causes to this strain relaxation phenomenon, and related to the condensation process, were suggested and investigated:

ȃ The presence of Ge-O bonds at oxide-SGOI interfaces was thought to weaken the mechanical properties of oxide-SGOI interfaces and to cause a "sliding" mechanism of the SGOI film. We found that there is no noticeable quantity of Ge-O bonds at the SGOI-BOx interface and thus ruled out that proposition. Nevertheless, we observed a higher quantity of Si sub-oxides at the BOx-SGOI interface after condensation and discussed it in view of the crystal disorder evidenced by MEIS.

ȃ H accumulation at SGOI-oxide interfaces was though to possibly induce mechanical weakness of the interfaces. By coupling neutron reflectivity with XRR, we found that there is neither H after SiGe epitaxy on SOI substrates not after condensation. To conclude, H is not the cause of strain relaxation of SGOI films after STI etch.

In summary, we ruled out two possible causes of the strain relaxation phenomenon and further investigation is therefore required to identify the cause of the strain relaxation mechanism.

Conclusion 181/230

Conclusion

The demand for higher speed and lower power consumption Integrated Circuits (ICs) has roused much interest into techniques that offer higher electron and hole mobility into channels of MOSFETs. Thanks to its high hole mobility, compressively strained SiGe is one of the most promising candidates for p-type MOSFET channel material. Moreover, advantages of strained SiGe in terms of hole mobility can be combined with the benefits of On Insulator (OI) structures, i.e., lower leakage current and better electrostatic control of the channel. Therefore, fabrication of SGOI structures may be a key milestone in enhancing performances of pMOSFETs. SGOI layers can be fabricated by the condensation technique. However, several challenges still need to be taken up for an efficient and optimized use of the condensation technique in an industrial environment. In order to complete these challenges, this thesis strived to bring understanding about thermal oxidation of SiGe, upon which the condensation technique is based, and solutions for optimizing the condensation process itself.

Chapter I has described the context, the theoretical and technological underpinnings of this thesis work. More precisely, the MOSFET technology has been introduced and its limitations to further downsizing have been described. New transistor architectures and high mobility channels, two main paths for enhancement of transistor performances, were reviewed. Among these performance "boosters", the benefits of using compressively strained SGOI films as channel material for p-type MOSFETs were presented. Specificities of the condensation technique, which allows the fabrication of SGOI films with possibly high strain levels, were described. Finally, the two main challenges of this thesis, which are addressed in Chapter III and Chapter IV, were defined:

• Understanding and controlling the mechanisms and kinetics of thermal oxidation of SiGe.

• Fabricating highly strained and high crystal quality SGOI films by the condensation technique.

Chapter II has reviewed the various processes, characterization techniques and the simulation tool used during this work. Techniques I personally used, either for the measurement and/or the analysis, were described in Chapter II while the other ones are presented in Appendices. Processes to fabricate samples, with a focus on the thermal oxidation processes used in this work, i.e. dry furnace oxidation, dry Rapid Thermal Oxidation (RTO) and In-Situ Steam Generation (ISSG) oxidation, were firstly presented. Then, characterization techniques using equipment considered as "common" equipment were reviewed: ellipsometry and XRR for thickness measurement, XPS for studying the chemical composition of thermal oxides grown on SiGe, Raman spectroscopy for measurement of strain of SGOI films, and GPA of HRTEM images for measurement of strain levels as well as for revealing crystal defects. Techniques that required the use of large-scale facilities were also reviewed: XRR experiments performed on SEXTANTS beamline at Soleil synchrotron in order to measure the density of thermal oxides, neutron reflectivity to identify whether or not there is hydrogen in SGOI films, and finally MEIS experiments to study the depth distribution of crystal perturbations and defects in SGOI films. This chapter ended with a description of the FEM and the environment of the developed simulation tool.

In the first part of Chapter III (Section III.A), an in-depth study of the mechanisms and kinetics of SiGe oxidations by means of three technology-relevant processes was carried out.

It was shown that, except possibly during the very first stages of oxidation (the temperature ramp up), the composition of oxides grown on SiGe is pure SiO 2 for all investigated oxidation processes, in agreement with thermodynamic calculations.

Then, we showed that oxidation of SiGe by means of dry furnace oxidation, dry RTO and ISSG oxidation is limited by diffusion of the oxidizing species through the growing oxide. A factor called Growth Rate Enhancement (GRE) was then defined to quantitatively compare the oxidation rates of Si and SiGe. We found that the GRE, and consequently oxidation kinetics, is correlated with the oxide density via the modulation of O 2 diffusivity. The influence of the process parameters on the oxide density, and as a result on oxidation kinetics, was finally clarified: the oxide density decreases when the oxidation temperature and the Ge concentration at the oxidation interface increase. Finally, we proposed a model of O 2 diffusivity that depends on the oxide density through the modulation of the diffusion barrier and which is consistent with our experimental data.

In the second part of this chapter (Section III.B), two technological applications based on the former results were presented. The first one is a graphic assessment of the oxide thickness and of the Ge diffusion length in SiGe, a key parameter of the evolution of the SiGe composition during oxidation, as a function of process parameters. The tool compares the relative importance of the Ge pile-up mechanism caused by the Si-selective oxidation with respect to SiGe interdiffusion. And secondly, a numerical simulation tool of thermal oxidation was adapted from the one of Morin et al. [Morin15]. The simulation model proved its worth in bringing more understanding of the dynamics of SiGe oxidation in spite of some remaining discrepancies with experimental data.

High strain levels and high crystal quality of SGOI films are needed for increasing the hole mobility. We thus compared in the first part of Chapter IV (Section IV.A) strain levels and crystal qualities of SGOI films fabricated by two industry-compatible condensation processes.

We firstly discussed strain levels of SGOI films fabricated by condensation based on high temperature dry RTO and ISSG oxidaion. We showed that in the case of SGOI films in the 10 nm thickness range, the parameter that has the strongest influence on strain relaxation is the strain level and not the total elastic strain energy nor the Ge concentration. Low temperature condensation by means of ISSG oxidation was also performed but found to be not adequate for fabricating high Ge content fully strained SGOI films. For technology applications, we highlight that the best condensation process that we studied is dry RTO at high temperature, with a threshold value of more than 2% for SGOI films in the 10 nm thickness range. Such a strain value corresponds to a Ge concentration of about 60%.

A finer characterization of the crystal quality as a function of depth into the SGOI film was performed by MEIS. Unfortunately, we cannot discuss optimization of the condensation process in terms of SGOI crystal quality with the characterization of three samples only. However, two conclusions can be drawn: the key crystal areas that are degraded and that must be particularly looked at for optimization of the condensation process are the interfaces of SiGe with oxides, and the SOI crystal quality is not altered by SiGe interdiffusion for low Ge fraction SGOI films.

In the second part of Chapter IV (Section IV.B), it was shown that the SGOI film is partially relaxed after the STI etch process step at the frontier between pMOS active areas and STI areas. Two possible causes to this strain relaxation phenomenon, and related to the condensation process, were suggested and investigated:

• The presence of Ge-O bonds at oxide-SGOI interfaces was proposed to weaken the mechanical properties of oxide-SGOI interfaces and to cause a "sliding" mechanism of the SGOI film. We found that there is no noticeable quantity of Ge-O bonds at the SGOI-BOx interface and thus ruled that proposition. Nevertheless, we observed a higher quantity of Si sub-oxides at the BOx-SGOI interface after condensation and discussed it in view of the crystal disorder evidenced by MEIS.

• H accumulation at SGOI-oxide interfaces was though to possibly induce mechanical weakness of the interfaces. By coupling neutron reflectivity with XRR, we found that there is H neither after SiGe epitaxy on SOI substrates not after condensation. To conclude, H is not the cause of strain relaxation of SGOI films after STI etch.

This work has brought more understanding about the mechanisms and kinetics of thermal oxidation of SiGe and about the evolution of the strain level and crystal quality of SGOI films fabricated by condensation. Perspectives to this work and questions left hanging in the air are numerous:

• What are the mechanisms that could account for the dependence of the oxide density on the oxidation temperature and on the Ge concentration at the oxidation interface? We suggested that the higher lattice parameter of SiGe with respect to Si could tend to accommodate the lattice mismatch between SiGe and SiO 2 , and as a consequence decrease the density of the oxide. As for the effect of temperature, it was proposed that stress relaxation of SiO 2 via viscous flow could decrease the oxide density. We suggest that studying oxidation kinetics of fully relaxed SiGe films and varying the conditions of the temperature ramp up of each oxidation process (which are largely different from dry RTO to dry furnace oxidation) could shed light on the mechanisms that induce this dependence.

• What are the causes of the discrepancies of the simulation model and how could it be improved? It was suggested that dedicated experiments to calibrate the diffusion model on the specific processes performed in this work would help improve the model.

• Optimizing the condensation process requires a sound understanding of the mechanisms of strain relaxation of SGOI films fabricated by condensation experiments, which are not fully identified yet. In terms of strain relaxation mechanisms, we suggested that if strain relaxation is governed by strain energy balance, this balance is influenced by the nature of the underlying layer: a thermal oxide (BOx) for the case of SGOI films instead of crystalline Si for the case of epitaxial SiGe-Si systems. We also suggested that dislocation nucleation at SiGe-oxide interfaces may be hindered up to higher elastic strain values than at SiGe-Si interfaces. Condensation based on other processes, such as dry furnace oxidation and ISSG oxidation with a different O 2 /H 2 ratio for instance could contribute to a better understanding of these mechanisms.

• The cause to the SGOI strain relaxation phenomenon at the edge of active areas has not been determined yet.

Appendices

A. Stress and Strain

This appendice complements Section I.2.2. by giving a more general and complete description of stress and strain laws. Such a description is needed as soon as studied systems are more complex than full-sheet layer. Investigating the mechanical behaviour of 2D or 3D nanoscale transistors, such as in [Morin15'], indeed requires the full knowledge of elastic moduli or compliance tensors.

Strain is defined as the relative lattice displacement created by deformation of the crystal structure of the material. Strain is fully characterized by the following second-order tensor: The strain tensor is symmetric:

ε ε ε ε ε ε ε ε ε ε =          
ε ε αβ βα α β β α = = ∂ ∂ + ∂ ∂         1 2 u x u u .
Stress is defined as the force responding to strain in a unit area and is fully defined by the second-order stress tensor: Within the linear elastic theory, stress and strain are related by Hooke's law:

σ ε = C. or ε σ = S. ,
where C and S are the elastic moduli and compliances fourth-order tensors. C and S are inverse to each other.

Symmetry relations induce:

C C C C ijkl jikl ijlk klij = = = .
The elastic moduli tensor has therefore 21 independent components. A matrix notation, called Voigt notation, is usually used to take into account symmetry considerations. It uses the indices I, J according to the rule:

(i,j)/(k,l) (1,1) (2,2) (3,3) (2,3) (3,1) (1,2) I/J 1 2 3 4 5 6
Table I.1: Voigt notation.

Hooke's law is then expressed in Voigt notation as:

σ σ σ σ σ σ xx yy zz yz zx xy C C C C C C C                     = 11
C C C C C C C C C C xx y                     ε ε y y zz yz zx xy ε ε ε ε 2 2 2                     .
Using Voigt notation, symmetry relations of cubic systems such as Si and SiGe reduce the elastic moduli tensor to:

B. Secondary Ion Mass Spectroscopy (SIMS)

Secondary-Ion-Mass Spectroscopy (SIMS) is a chemical composition characterization technique based on sputtering by ion bombardment. The incident ions, called primary ions, are focused on the sample with a kinetic energy from a few keV to a few hundreds keV, which induce sputtering of various particles, among which ions which are called secondary ions, from the surface of the sample. The chemical composition of the surface of the sample is extracted from the analysis of secondary ions by a mass spectrometer. A schematic view of the SIMS apparatus is shown Figure B.1. As the surface of the sample is eroded by the incident ion beam, chemical composition profiles as a function of depth can be extracted assuming sound understanding of the ionization probabilities of the studied materials and proper calibration of the sputter rate.

Time-of-Flight SIMS (ToF-SIMS) uses a Time-of-Flight mass spectrometer instead of a magnetic spectrometer. Separation of elements is performed by measuring the time of flight between the sample and the detector, which depends on the mass of each element. ToF-SIMS thus uses a pulsed primary ion beam, usually with a frequency in the 10 kHz range.

SIMS and ToF-SIMS using a 300 V O 2 + beam (CAMECA IMS Wf) were performed to measure the Ge concentration as a function of depth after oxidation. Ge concentration quantification was performed on Si 0.8 Ge 0.2 well known samples prior to any SIMS characterization. The variation of the sputter rate with the Ge concentration is low for Ge concentrations in the 20%-30% range. We can therefore reasonably extract Ge concentrations as a function of depth for most of the samples of this study. No conclusion is based on the depth dependence of the Ge concentration when the Ge concentration is higher, even if Ge concentration profiles are plotted as a function of depth as well for direct comparison. When an incoming X-Ray photon is absorbed by the material, its energy can be transfered to the core electron of an atom. The core electron then goes onto an unoccupied higher energy level or becomes a free electron. The formerly occupied level is almost immediately filled with an electron of a higher energy level. The energy gap between these two levels is evacuated through the emission of either a X-Ray photon or an electron called Auger electron. X-Rays emitted through this process have energies which are specific to the probed atom energy levels and thus allow identification of the probed element Fluorescence spectral lines, which are specific to each element, are usually named according to Siegbahn notation. X-Rays are hardly absorbed in films a few nanometers-thick so that we can assume that the intensity of the spectral lines is proportional to the total quantity of the probed element. As a result, in identical acquisition conditions, comparison of the areas of Kα and Lα spectral lines of Ge gives a relative comparison between the total where λ is the X-Ray wavelength, θ the angle between the incident wave vector and the family of diffracting planes {hkl}, and n the diffraction order. While there are many possible XRD experimental conditions, only one experimental set-up was used in this work to extract the composition and thickness of strained SiGe layers. As illustrated Figure D.2, the analysis was performed in a θ-2θ scan along the (004) direction perpendicular to the surface. A typical XRD spectra of a fully strained and homogeneous SiGe layers on a Si substrate is shown Figure D.3. In this case, the out-of-plane SiGe lattice parameter is higher than the one of Si according to Poisson's effect (see Section I.2.2.3.). It results in the {004} family of SiGe planes to have Bragg's conditions different from the ones of the Si substrate, and as a consequence a (004) diffraction peak shifted with respect to the one of Si.

C. Wavelength Dispersive X-Ray Fluorescence (WDXRF)

D. X-Ray Diffraction (XRD)

The (004) SiGe peak position is only dependent on the out-of-plane SiGe lattice parameter at a fixed X-Ray wavelength. Assuming that the SiGe layer is fully strained, the Ge concentration can be extracted from Dismuke's Equation D.1 quantities of Ge atoms in various SiGe films. XRF Kα and Lα Ge lines were measured in Wavelength Dispersive mode (based on the selective reflection of a specific wavelength on a crystal according to Bragg's law) on a Rigaku AZX400 equipped with a Rhodium source, ten different crystals for the wavelength selection and two different detectors (a scintillation counter and a gaseous proportional counter). law and Poisson's effect [Dismukes64]. The SiGe layer thickness can be extracted from the width of the main peak and of the interference fringes. The width of the main peak is inversely proportional to the thickness of the SiGe layer. One limitation of XRD is that the signal-to-noise ratio becomes too low for proper fitting for very thin films because the number of diffracting (004) planes is reduced to a few ones. Moreover, the width of the main peak is inversely proportional to the thickness of the SiGe layer so that interference fringes disappear when the SiGe thickness decreases. This effect also reduces fitting robustness.

XRD measurements were performed on a commercial Jordan Valley JVX7300 equipment in STMicroelectronics Crolles clean room. Atomic Force Microscopy (AFM) is a surface characterization technique based on the measurement of interaction forces, among which electrostatic forces, van der Walls forces and short-range chemical forces, between the surface of the sample and a sharp tip attached to the free end of a cantilever. While many properties of the surface of the layer can be measured by AFM, such as viscoelasticity or adhesion strength, topography measurements only were performed in this study. As illustrated Figure E.1, the interaction forces induce a deflection of the cantilever, which is then measured thanks to the deflection of the reflected laser beam. Measurements were performed on 5x5 μm 2 squared surfaces on a Bruker Icon Head tool in tapping mode. There are two main formula of interest used in this manuscript about AFM characterization: ȃ Root Mean Squared (RMS): standard deviation of the Z (height) values with a given area:

RMS Z Z N i a ve i N = - ( ) = ∑ 1 ,
where Z i is the current Z value, Z ave the averaged Z with a given area and N the number of points with a given area.

ȃ R max : difference in height between the lowest and highest point on the surface relative to the mean plane.

E. Atomic Force Microscopy (AFM)

F. Transmission Electron Microscopy (TEM)

Transmission Electron Microscopy (TEM) is a technique that allows a broad range of characterization techniques, from imaging to diffraction and chemical analysis, at the nanometric scale. TEM is described in depth in many textbooks such as [START_REF] Zienkiewicz | The finite element method[END_REF] and the apparatus of a transmission electron microscope is shown Figure F.1. High Resolution TEM (HRTEM) with atomic resolution, which is a specific mode of TEM imaging, was performed in this thesis. In HRTEM mode, a parallel electron beam is formed through specific electromagnetic lenses and directed towards a thin lamella (less than 100 nm thick). The contrast of the image is a phase contrast due to interferences between the transmitted and diffracted beams. HRTEM is particularly dependent on the quality of the TEM lamella, and interpretation of HRTEM images must be cautiously done.

TEM lamella were prepared using standard Focused Ion Beam (FIB) and imaged in a JEOL 2010 equipped with a LaB 6 electron gun, and with an accelerating voltage of 200 kV.

Résumé en français

Sommaire Introduction

Entre une minute à consulter les réseaux sociaux sur son smartphone et une autre à lire le dernier best-seller sur sa liseuse, levons les yeux vers la façade d'un immeuble d'une mégalopole recouvert de panneaux publicitaires géants, approchons notre carte de transport d'une borne magnétique dans un tramway d'une ville moyenne, ou encore écoutons les signalements sonores indiquant le passage d'un train à une intersection avec une route de campagne, et nous voilà face à une débauche de systèmes électroniques aussi variés dans leur fonctionnement que dans leur application. Un point commun toutefois à ces systèmes ? Oui, le transistor. La complexification galopante des systèmes électroniques repose en effet principalement sur la miniaturisation de la brique élémentaire de ces systèmes électroniques qu'est le transistor.

Jusque dans les années 2000, la réduction des dimensions des transistors a obéit à une loi, ou plutôt une prophétie auto-réalisatrice, édictée par Gordon Moore en 1965 et nommée d'après lui. La loi de Moore prédit que la taille des transistors est divisée par deux tous les deux ans. Cette loi, dont la force est sa capacité à constituer un objectif clair et simple dans sa formulation pour l'industrie de la microélectronique, se heurte toutefois à des obstacles croissants à mesure que les dimensions des transistors se rapprochent de celles de quelques couches atomiques. La loi de Moore est désormais poussée dans ses retranchements grâce à l'introduction dans le transistor d'éléments nouveaux permettant d'augmenter les performances et/ou de réduire la consommation de ce dernier.

Parmi ces éléments figurent l'intégration d'un nouveau matériau, le SiGe contraint en compression, dans le canal d'un transistor dont l'architecture non-conventionnelle est de type FD-SOI (Fully-Depleted Silicon-On-Insulator). L'intérêt de films SiGe ultra-minces, contraints en compression, et sur isolant (SGOI: SiGe-On-Insulator) est double : ils bénéficient de la forte mobilité des trous dans le SiGe contraint en compression ainsi que du meilleur contrôle électrostatique des structures dites « sur isolant ». Les films SGOI présentant une forte concentration en Ge et une importante contrainte peuvent être fabriqués par une technique industrielle appelée condensation. Cette technique repose sur deux processus simultanés : l'oxydation thermique et sélective du SiGe (seul le Si est oxydé) et l'inter-diffusion du SiGe entre l'oxyde thermique et l'oxyde enterré qui se comporte comme une barrière à la diffusion.

L'utilisation de cette technique dans un environnement industriel nécessite de relever deux défis : maîtriser les mécanismes et la cinétique d'oxydation, et atteindre les plus fortes contraintes et qualités cristallines pour les films SGOI. L'organisation de ce résumé en français est la suivante :

• Le Chapitre I présente le contexte dans lequel s'inscrit ce travail de thèse et les fondements de ce dernier.

• Le Chapitre II traite des mécanismes et de la cinétique d'oxydation du SiGe.

• Le Chapitre III porte sur la caractérisation des contraintes et de la qualité cristalline des films SGOI fabriqués par condensation. Quant aux performances intrinsèques du transistor, elles sont fortement dépendantes au contrôle électrostatique du canal par la grille et à la mobilité des trous dans le canal. Tandis que l'architecture de type FD-SOI permet un meilleur contrôle électrostatique du canal par rapport aux architectures de transistor traditionnelles, les films SiGe contraints en compression présentent des mobilités de trous supérieures [START_REF] Prabhakaran | In situ oxidation of a thin layer of Ge on Si (001): Observation of GeO to SiO 2 transition[END_REF]. Afin de bénéficier de ce double avantage, il est nécessaire de pouvoir fabriquer des films SiGe ultra-minces, contraints en compression, et sur isolant (SGOI : SiGe-On-Insulator) à l'aide d'un procédé industriel: la condensation [START_REF] Bongiorno | Multiscale modeling of oxygen diffusion through the oxide during silicon oxidation[END_REF]. Cette technique, illustrée Figure I.2, repose sur deux processus simultanés : l'oxydation thermique et sélective du SiGe (seul le Si est oxydé) et l'inter-diffusion du SiGe entre l'oxyde thermique et l'oxyde enterré qui se comporte comme une barrière à la diffusion. Trois caractéristiques font de ce procédé l'un des plus pertinents pour la fabrication de ces films dans un environnement industriel : ȃ La condensation permet la fabrication de films SGOI dans les zones de transistors de type p uniquement (les films SGOI ne présentent pas d'intérêt pour les transistors de type n)

Chapitre I -La technique de condensation appliquée aux technologies CMOS FD-SOI

ȃ Des films SGOI non-relaxés (ou autrement dit totalement contraints à la maille du Si), de quelques nanomètres seulement, et présentant des concentrations en Ge de l'ordre de 50% peuvent être fabriqués.

ȃ La condensation ne nécessite l'usage que de substrats conventionnels et de procédés compatibles avec le reste de la chaîne de production des systèmes électroniques. Le traitement mathématique de ces étapes conduit à l'établissement d'une loi cinétique dite linéaire-parabolique : La cinétique de l'oxydation sèche en four et par RTP et de l'oxydation ISSG est tout d'abord décrite en fonction de l'évolution de la composition du film SiGe sous l'oxyde. La deuxième partie de ce chapitre se penche sur la corrélation entre la cinétique d'oxydation et la densité de l'oxyde thermique en croissance. Enfin, la dépendance de la diffusion de l'O 2 dans SiO 2 , dont dépend la cinétique de l'oxydation sèche, est discutée. Deux hypothèses de travail, sur lesquelles repose le principe de la technique de condensation, ont été vérifiées au préalable: ȃ Une étude par spectroscopie de photoélectrons à rayons X (XPS : X-Ray Photoelectron Spectroscopy) nous a permis de montrer que l'oxyde thermique formé sur SiGe est le SiO 2 (dioxyde de silicium) en phase amorphe quel que soit le procédé d'oxydation utilisé.

Si

ȃ La conservation de la totalité du Ge, repoussé dans le film SiGe à l'interface avec l'oxyde lors de l'oxydation, a été démontrée par fluorescence de rayons X (XRF : X-Ray Fluorescence).

II.1.3. Régime d'oxydation

II.1.6. Evolution de la composition du film SiGe pendant l'oxydation

Deux mécanismes thermiquement activés sont à l'origine de l'évolution du profil en concentration en Ge du film SiGe : ȃ L'oxydation sélective du Si repousse le Ge sous l'oxyde en croissance et tend donc à augmenter la concentration en Ge à l'interface oxyde-SiGe : c'est l'effet « pile-up ». ȃ L'inter-diffusion du SiGe, à l'inverse, tend à lisser le profil en concentration en Ge.

L'évolution du profil en concentration en Ge, et en particulier la concentration à l'interface oxyde-SiGe, dépend de la compétition entre ces deux mécanismes. A basse température, l'inter-diffusion dans le SiGe est très faible et l'effet « pile-up » dominant, ce qui conduit à la formation d'une couche riche en Ge sous l'oxyde (entre 45% et 50% de concentration en Ge pour le cas de l'oxydation sèche par RTP de Si 0.7 Ge 0.3 à 900°C), comme illustré en Figure II.6 (a). Lorsque la température augmente, les deux mécanismes s'accélèrent. Du fait d'une importante énergie d'activation (autour de 4 eV [Kube10]), l'inter-diffusion accélèrent de façon plus marquée que l'oxydation (qui présente des énergies d'activation de l'ordre de 2 eV et 1 eV pour l'oxydation sèche par RTP et l'oxydation ISSG respectivement).

La Figure II.6 nous permet de souligner un deuxième point : la variation de la concentration en Ge à l'interface oxyde-SiGe est faible au cours du temps et reste, pour le cas de l'oxydation sèche par RTP, autour de 33% à 1100°C, 42% à 1000°C, et 48% à 900°C. La cinétique d'oxydation est faiblement dépendante de la concentration en Ge pour le cas de l'oxydation ISSG tandis qu'elle se trouve accélérée par la présence de Ge pour le cas de l'oxydation sèche. Ayant montré que la cinétique d'oxydation est limitée par la diffusion de l'espèce oxydante à travers l'oxyde, il semble que la présence de Ge influence la diffusion de l'espèce oxydante à travers l'oxyde. Il a été démontré dans la littérature que la diffusion d'O 2 dans SiO 2 dépend de la densité de l'oxyde. C'est pourquoi nous nous penchons dans la partie suivante sur la dépendance de la densité de l'oxyde thermique crût sur SiGe aux conditions du procédé d'oxydation.

II.2. Corrélation entre cinétique d'oxydation et densité de l'oxyde

II.2.1. Protocole expérimental

Les densités des oxydes thermiques crûs sur Si et SiGe par différents procédés d'oxydation ont été mesurées par réflectivité de rayons X à 500 eV sur la ligne SEXTANTS du synchrotron Soleil. Les spectres de réflectivité ont ensuite été ajustés aux spectres théoriques à l'aide du programme dyna.

II.2.2. Densité des oxydes thermiques

L'évolution de la densité en fonction des paramètres du procédé d'oxydation et de la concentration en Ge est mise en évidence en Figure II.8 : ȃ 1: Concentration initiale en Ge : pour l'oxydation sèche par RTP et l'oxydation ISSG, la densité de l'oxyde diminue lorsque la concentration initiale en Ge augmente. ȃ 2: Type de procédé : la densité de l'oxyde crû sur Si 0.7 Ge 0.3 par oxydation sèche en four est plus élevée que celle de l'oxyde crût par oxydation sèche par RTP. ȃ 3: Température : la densité des oxydes crûs sur Si 0.7 Ge 0.3 par oxydation sèche par RTP augmente avec la température, à l'inverse des oxydes crûs sur Si par le même procédé. Dans la continuité de la conclusion de la partie II.1.5., nous ne pouvons pas identifier les mécanismes à l'origine directe des variations de densité observées en Figure II.8. Nous considérons dans la partie suivante la concentration en Ge à l'interface oxyde-SiGe.

II.2.3. Dépendance de la densité de l'oxyde aux paramètres du procédé d'oxydation

La Figure II.9 montre que plus la concentration en Ge à l'interface est élevée, plus la densité de l'oxyde est faible. Tel que proposé par Holland et al. [Holland87], nous suggérons que le plus grand paramètre de maille du SiGe pourrait réduire le stress compressif dans l'oxyde à l'interface oxyde-SiGe engendré par la différence de densité atomique entre les deux matériaux. La réduction de ce stress serait alors à l'origine de la croissance d'un oxyde moins dense. La densité systématiquement plus faible à 1100°C qu'à 900°C pour l'oxydation sèche par RTP pourrait quant à elle provenir du fluage de l'oxyde qui tendrait à diminuer le stress et par conséquent la densité de l'oxyde [EerNisse77]. Enfin, nous soulignons que, tous paramètres égaux par ailleurs, l'oxyde crû par oxydation sèche en four est plus dense que celui crû par oxydation sèche par RTP.

Figure II.9: Evolution de la densité de l'oxyde en fonction de la concentration en Ge à l'interface.

Ayant montré et commenté les influences des paramètres matériaux et du procédé d'oxydation sur le densité de l'oxyde, nous nous penchons désormais sur la dépendance de la cinétique d'oxydation à cette densité.

II.2.4. Dépendance de la cinétique d'oxydation à la densité de l'oxyde

Conformément à une oxydation limitée par la diffusion d'O 2 à travers l'oxyde, la cinétique de l'oxydation sèche augmente lorsque la densité de l'oxyde diminue (voir Figure II.10). En effet, les molécules d'O 2 diffusent à travers le SiO 2 de site interstitiel en site interstitiel d'autant plus rapidement que la densité de l'oxyde est faible et donc les sites interstitiels grands [START_REF] Bongiorno | Reaction of the Oxygen Molecule at the Si(100)-SiO 2 Interface During Silicon Oxidation[END_REF]. Le fait que le cas de l'oxydation sèche en four à 900°C soit aligné avec l'oxydation sèche par RTP à la même température confirme la dépendance directe de la cinétique d'oxydation et de la diffusion d'O 2 à travers l'oxyde à la densité. Quant au décalage précédemment mentionné dû à la seule température, il sera discuté en partie suivante.

La densité de l'oxyde n'influence presque pas la cinétique d'oxydation ISSG. Il est logique, compte tenu de la petite taille des radicaux d'oxygène par rapport aux molécules d'O 2 , que la diffusion de ces radicaux ne soit pas affectée, ou ne le soit que faiblement, par la densité de l'oxyde.

Grâce à la dépendance du GRE au coefficient de diffusion de l'espèce oxydante dans l'oxyde, nous pouvons discuter spécifiquement de l'influence de la densité de l'oxyde sur la diffusion de l'O 2 dans SiO 2 .

II.2.5. Dépendance de la diffusion de l'O 2 à la densité de l'oxyde

Des considérations thermodynamiques quant à la dépendance du coefficient de diffusion à la température et à la pression nous permettent de proposer le modèle suivant pour le coefficient de diffusion d'O 2 dans SiO 2 L'influence de chacun des paramètres d'oxydation (chimie d'oxydation : oxydation sèche versus ISSG, type de procédé : four versus RTP, concentration initiale en Ge, température et concentration en Ge à l'interface) a été discutée. Nous retiendrons en particulier que tous paramètres égaux par ailleurs, plus la concentration en Ge à l'interface oxyde-SiGe est importante, moins l'oxyde formé est dense. La densité de l'oxyde est également d'autant plus faible que la température d'oxydation est élévée.

La dépendance directe de la cinétique d'oxydation sèche à la densité a ensuite été mise en évidence : moins l'oxyde est dense, plus la diffusion d'O 2 dans SiO 2 , et par conséquent la cinétique d'oxydation, sont rapides. Quant à l'oxydation ISSG, elle ne dépend pratiquement pas de la densité de l'oxyde, indiquant ainsi la faible dépendance de la diffusion des radicaux d'oxygène O* à la densité de l'oxyde.

Enfin, un modèle du coefficient de diffusion de l'O 2 dans SiO 2 a été proposé. Ce dernier dépend de la densité de l'oxyde à travers la modulation de la barrière à la diffusion, modulation qui se trouve observer une variation linéaire par rapport à la densité. Nous avons dans un premier temps montré que pour les trois procédés d'oxydation étudiés, l'oxydation sèche en four, l'oxydation sèche par RTP et l'oxydation ISSG, la cinétique d'oxydation est déterminée par la diffusion de l'espèce oxydante à travers l'oxyde. Tandis que l'oxydation ISSG est peu dépendante à la présence de Ge, l'oxydation sèche du SiGe est plus rapide que celle du Si. La définition d'une méthode d'analyse pertinente au regard du régime d'oxydation observé nous a permis de mettre en évidence une corrélation entre la cinétique d'oxydation, la diffusion d'O 2 dans SiO 2 et la densité de l'oxyde. La diffusion des radicaux d'oxygène O*, qui sont les principales espèces oxydantes de l'oxydation ISSG, est largement indépendante de la densité de l'oxyde. Nous avons ensuite clarifié l'influence du Ge et de la température d'oxydation sur la densité de l'oxyde: la densité diminue lorsque la concentration en Ge à l'interface avec l'oxyde et la température augmentent. Le plus grand paramètre de maille du SiGe pourrait être à l'origine de la variation de densité de l'oxyde due à la présence de Ge tandis que le fluage de l'oxyde à haute température a été suggéré en tant que mécanisme à l'origine de l'influence de la température sur la densité. Enfin, cette étude s'est conclue avec la proposition d'un modèle de coefficient de diffusion de l'O 2 dans SiO 2 , qui se trouve dépendre de la densité à travers la modulation de la barrière à la diffusion. Nous commençons par décrire l'évolution de la contrainte des films SGOI fabriqués par condensation à haute température en fonction de la concentration en Ge. Les mécanismes de relâchement des contraintes sont discutés à la lumière de la comparaison avec la littérature, de la nature des interfaces du film SGOI et des contraintes des films SGOI fabriqués par condensation à basse température. Ensuite, le nombre d'atomes déplacés des sites cristallins des films SGOI, qui est un indicateur de la qualité cristalline, est mesuré en fonction de la profondeur dans le film. Les effets du procédé de condensation et de la présence intrinsèque de Ge sur la qualité cristalline sont soulignés. Afin d'évaluer l'influence de la nature des interfaces des films SiGe (cristal-cristal dans le cas de films SiGe épitaxiés sur Si contre cristal-oxyde amorphe dans le cas de films SGOI), nous avons comparé les contraintes des échantillons fabriqués par condensation à haute température avec celles des films SGOI fabriqués par le procédé de condensation à basse température décrit en partie III.1.1.. Tandis que la condensation à haute température conduit rapidement à la formation d'un film SGOI relativement homogène en composition, la condensation à basse température permet la formation par oxydation d'une couche SiGe à 60% de concentration en Ge environ sur le cristal de Si 0.7 Ge 0.3 puis sur celui de Si du SOI. La nature des interfaces lors de la condensation est donc radicalement différente entre les deux procédés.

Conclusion du Chapitre II

Chapitre III -Déformation et qualité cristalline de films SGOI fabriqués par condensation

Les contraintes des films fabriqués par condensation à basse température sont fortement relâchées et inférieures aux contraintes des films SGOI de composition identique mais fabriqués par condensation à haute température (voir Figure III.5). Les interfaces SiGe-Si semblent donc favoriser le relâchement plastique des contraintes par rapport aux interfaces SiGe-oxyde amorphe. Nous suggérons que ces interfaces SiGe-oxyde amorphe modifient l'équilibre entre énergie élastique et énergie plastique et/ou retarde la formation de dislocations, qui sont les défauts cristallins accompagnant le relâchement des contraintes comme nous le verrons en partie suivante, dans le film SGOI. En résumé, le relâchement des contraintes s'accompagne de la formation de dislocations dans les plans {111} du SiGe. Le cas du film SGOI à 57% de concentration en Ge est difficile à trancher du fait de la contradiction entre l'image AFM qui ne révèle pas de dislocations sur une surface d'échantillon de 5×5 μm 2 , la GPA qui en révèle sur des lames minces pour TEM.

III.1.6. Conclusion de la partie III.1.

Nous avons montré que la contrainte des films SGOI fabriqués par condensation à haute température augmente avec la concentration en Ge jusqu'à une valeur maximale d'environ 2.1% à 60% de concentration en Ge, puis diminue au-delà de cette concentration. Le relâchement des contraintes est associé à la génération de dislocations Résumé en français 221/230 dans les plans {111} du cristal de SiGe. Toutefois, les mécanismes du relâchement des contraintes demeurent indéterminés, même si nous avons montré que l'influence de la nature des interfaces du cristal SiGe, avec un oxyde amorphe dans le cas de films SGOI, semble être majeure, et que le paramètre déterminant est la contrainte en ellemême, avec un maximum entre 2.2% et 2.5%.

Après avoir identifié les paramètres permettant l'obtention de contraintes maximales dans les films SGOI continus et fabriqués par condensation, nous nous attelons dans la partie suivante à caractériser finement la qualité cristalline de ces films.

III.2. Mesure de la qualité cristalline des films SGOI

III.2.1. Introduction de la technique MEIS

La technique Medium Energy Ion Scattering (MEIS), ou rétrodiffusion d'ions de moyenne énergie, permet de déterminer le profil de composition en profondeur et certaines propriétés physiques, telle que la qualité cristalline, de couches minces [START_REF] Kahng | Electric field controlled semiconductor device[END_REF]. La technique repose sur la mesure du nombre d'ions rétrodiffusés par les atomes de l'échantillon, ainsi que de leur angle de rétrodiffusion et leur énergie. A angle fixé, la mesure MEIS se présente sous la forme de spectres en énergie tel que celui représenté en Figure III.7, où la contribution de chaque élément de chaque couche est identifiée. ȃ « Faible Ge » : le film SGOI, dont la composition est de l'ordre de 25% en concentration de Ge, a été réalisé par un procédé de condensation reposant sur une oxydation sèche par RTP à 1100°C.

ȃ « Fort Ge » : d'une concentration moyenne en Ge de 57%, il s'agit du même film SGOI que celui discuté en partie III.1.. Ce film SGOI est celui présentant le maximum de contrainte sans révélation de dislocations par mesure AFM dela surface du film.

III.2.2. Protocole expérimental III.2.3. Spectres MEIS

Les spectres MEIS expérimentaux et ajustés pour chaque échantillon et chaque condition (conditions de canalisation ou non) sont représentés en Figure III.9. Nous observons pour chacun des échantillons une réduction plus ou moins prononcée du nombre d'ions rétrodiffusés aux énergies correspondant aux couches SOI, SiGe et SGOI (entre 85 keV et 90 keV et autour de 72.5 keV). Chacun des spectres expérimentaux a été ajusté à l'aide d'un programme dont les modèles des empilements sont définis par un nombre fini de couches homogènes en composition et en χ min . Même si un nombre fini de couches ne peut pas reproduire exactement la réalité des profils en composition et en χ min , l'ajustement aux spectres expérimentaux est jugé suffisamment bon avec un nombre fini de couches pour en extraire des tendances de variation de χ min en profondeur. Ces tendances sont présentées et discutées en partie suivante. • La qualité cristalline du SOI n'est pas dégradée par la diffusion de Ge pour de faibles valeurs de concentration en Ge.

• La présence de Ge, même à faible concentration, semble provoquer une dégradation de la qualité cristalline du film SGOI à l'interface avec l'oxyde enterré (BOx) sur environ 1 nm.

• Lorsque la concentration en Ge augmente vers de fortes valeurs (50% et plus), la qualité cristalline du film SGOI commence à diminuer, et ce loin des interfaces avec les oxydes thermique et enterré. Nous avons discuté de l'évolution des contraintes en fonction des paramètres du procédé de condensation, et sommes parvenus à identifier le procédé parmi ceux étudiés conduisant à une contrainte maximale d'environ 2.1% à 60% de concentration en Ge. Il s'agit de l'oxydation sèche par RTP dont la température est aussi haute que possible sans faire entrer le SiGe en fusion.

SiGe Si

Deux aspects clés du procédé de condensation et des films SGOI ont été soulignés grâce à la caractérisation de la qualité cristalline de ces films :

ȃ Les zones dégradées du cristal, et qui méritent donc de recevoir une attention particulière lors de l'optimisation du procédé de condensation, sont les interfaces avec les oxydes thermique et enterré.

ȃ L'excellente qualité cristalline du SOI est maintenue lors de la condensation jusqu'à des concentrations en Ge de l'ordre de 50%, et ce malgré la diffusion de Ge dans le film.

Conclusion du Chapitre III

Conclusion

L'intégration de films SiGe contraints en compression sur isolant (SGOI) dans le canal du transistor de type p apparait comme l'une des alternatives aux transistors traditionnels les plus pertinentes. L'utilisation dans un milieu industriel de la technique de condensation, qui permet de fabriquer de tels films, nous astreint à maîtriser à la fois la cinétique de l'oxydation thermique du SiGe et la contrainte et la qualité cristalline du film SGOI ainsi formé.

Dans le Chapitre II, nous nous sommes penchés sur le premier de ces défis. Nous avons identifié les mécanismes d'oxydation du SiGe pour trois procédés d'oxydation appropriés à la condensation. Nous avons ainsi relié la cinétique d'oxydation avec la diffusion de l'espèce oxydante à travers l'oxyde en croissance et avons corrélé cette diffusion à la densité de l'oxyde. L'application de ces résultats fondamentaux à la condensation est double : ȃ Nous pouvons estimer graphiquement la cinétique d'oxydation du SiGe à l'aide d'abaques, dont la construction n'est pas détaillée dans ce résumé, et ce pour trois procédés d'oxydation et de larges gammes de concentrations en Ge et de températures.

ȃ Un programme de simulation par éléments finis, basé sur celui développé par Morin [Morin15], permettrait de poursuivre l'étude de la dynamique de l'oxydation du SiGe et de prédire avec précision les épaisseurs et compositions des films SGOI fabriqués par condensation.

Le Chapitre III a porté quant à lui sur l'évolution de la contrainte et de la qualité cristalline des films SGOI. Les paramètres de condensation permettant d'atteindre les plus fortes contraintes, supérieures à 2% relativement au SiGe relâché, ont été discutés. La qualité cristalline des films SGOI a pu être finement caractérisée, et les zones du cristal dégradées et les causes de dégradation, partiellement identifiées. Cette étude soulève une autre question : comment la contrainte et la qualité cristalline des films SGOI changent pendant la suite de la fabrication des transistors, et quel est le procédé de condensation adéquat au regard de toute la chaîne de fabrication ? Cette problématique a fait l'objet de deux études, l'une portant sur la nature des liaisons chimiques entre le film SGOI et l'oxyde enterré, et l'autre sur l'incorporation involontaire d'hydrogène au cours de la fabrication des films SGOI.

Les réponses aux problématiques précédemment définies ont aussi soulevé de nouvelles questions :

ȃ Nous avons montré dans le Chapitre II que plus la concentration en Ge à l'interface et la température d'oxydation sont élevées, moins l'oxyde est dense. Les mécanismes qui en sont à l'origine n'ont toutefois pas été identifiés, malgré deux suggestions, à savoir que le plus grand paramètre de maille du SiGe et le fluage de l'oxyde à haute température tendrait à réduire la densité de l'oxyde.

ȃ S'il a été mis en exergue, le rôle que joue la nature des interfaces du cristal de SiGe avec les autres couches dans les mécanismes de relâchement des films SGOI n'a pas été déterminé et mérite donc une attention particulière.

Figure

  Figure I.1: The First transistor fabricated at the Bell Labs.

Figure I. 2 :

 2 Figure I.2: Schematics of a pMOSFET (left) and a nMOSFET (right).

Figure I. 3 :

 3 Figure I.3: Schematic of a pMOSFET in off state.

Figure I. 4 :

 4 Figure I.4: Schematic of a pMOSFET in weak inversion regime.

  Figure I.5: Schematic of a pMOSFET in strong inversion regime.

Figure

  FigureI.6: Schematic (left) and drain current versus drain voltage (right) of a pMOSFET in the linear regime.

Figure I. 7 :

 7 Figure I.7: Schematic (left) and drain current versus drain voltage (right) of a pMOSFET at the pinch-off point.

Figure I. 8 :

 8 Figure I.8: Schematic (left) and drain current versus drain voltage (right) of a pMOSFET in the saturation regime.

  Figure I.10: Schematic of a MOSFET showing the pitch and half-pitch.

Figure I. 11 :

 11 Figure I.11: Evolution of the transistor density and minimum feature size (transistor gate length) in microprocessors. From [Ferain11].

Figure

  Figure I.12: Schematic view of a conventional bulk MOSFET. The depletion regions are in dark-gray. The width of the drainjunction depletion region extends into the channel as the drain voltage increases, causing DIBL. From [Ferain11].

FinFET

  Figure I.13: Schematics of FinFET and FD-SOI architectures. FD-SOI and FinFET architectures ultimately converge, differing only from a 90° rotation, as the Buried Oxide (BOx) thickness is thinned down. From [Ouellette15].

  Figure I.14: Transition from bulk to FD-SOI MOSFET brings better electrostatic control over the channel and lower leakage currents. From [Ouellette15].

  Figure I.15: FD-SOI and FinFET target markets. The FD-SOI technology addresses power sensitive markets whileFinFET is more suitable for high-performance markets. From[START_REF] O P Paine | Oxidation of Si 1-x Ge x alloys at atmospheric and elevated pressure[END_REF].

Figure

  Figure I.17: TEM image of a FD-SOI pMOS transistor fabricated at STMicroelectronics Crolles site. Internal document.

Figure I. 18 :

 18 Figure I.18: Schematic illustration of the dependence of the mobility on the effective gate field of a transistor. It shows the three scattering mechanisms and at which field they limit the mobility. From [Robertson15].

  [Yang03]. Advantageous substrate-channel orientations are illustrated Figure I.20 and have been used in commercial devices.

Figure I. 19 :

 19 Figure I.19: Calculated bulk electron and hole mobility of silicon, germanium and some III-V materials as a function of the bandgap. Open symbols indicate hole mobility and full ones indicate electron mobility.From[START_REF] Prabhakaran | In situ oxidation of a thin layer of Ge on Si (001): Observation of GeO to SiO 2 transition[END_REF].

Figure

  Figure I.21 shows calculated electron mobility gain in silicon for various strain types (uniaxial and biaxial), levels and crystallographic orientations.

  Figure I.22).

Figure I. 21 :

 21 Figure I.21: Calculated electron mobility enhancement as a function of tensile strain. From [Uchida05].

Figure

  Figure I.22: pMOSFET from Intel with compressively strained Si channel using SiGe source and drain (left). And nMOSFET from Intel with tensely strained Si channel using a SiN capping layer (right). From [Thompson04].

Figure I. 23 :

 23 Figure I.23: Schematic of a FD-SOI transistor showing the main application of this thesis work for advanced FD-SOI technologies.

.

  Figure I.24: Schematic of the diamond-like lattice structure of SiGe.

Figure

  Figure I.26: Band diagrams of (left) silicon and (right) germanium. From[IOFFE].

  β = , , and where u α is the displacement of the lattice point along x α .

.

  As schematically representedFigure I.29, σ xy (or τ xy ) is the force applied in the x direction to a unit area in the plane which is normal to the y direction.

  Figure I.29: Schematic representation of stress components on surfaces of an infinitesimal cube. From [Sun09].

Figure

  Figure I.30: Schematic illustration of a relaxed SiGe layer and a relaxed Si layer (left) and a pseudomorphic (fully strained) SiGe layer on relaxed Si (right).

Figure

  Figure I.31: Out-of-plane deformation of a pseudomorphic SiGe layer as a function of the Ge concentration.

  Figure I.33: Model of hole mobility variation in SiGe versus stress based on experimental piezoresistance coefficients. From [Andrieu14].

Figure I. 34 :

 34 Figure I.34: Schematic illustration of the epitaxy process.

  Figure I.35: Schematic representation of a pseudomorphic SiGe layer on top of a Si layer.

  Figure I.36 shows a schematic representation of the surface ondulations observed on a elastically relaxed SiGe layer on top of a Si layer.

Figure I. 36 :

 36 Figure I.36: Schematic representation of the surface ondulations observed on an elastically relaxed SiGe layer on top of a Si layer.

  Figure I.37: Schematic representation of a plastically relaxed SiGe layer on top of a Si layer.

Figure I. 39 :ȃ

 39 Figure I.39: Critical thickness of an epitaxial SiGe layer versus the Ge concentration calculated from the Matthews and Blakeslee model [Matthews74].

Figure

  Figure I.41: Definition of the burgers circuit and the burgers vector by comparison of (a) a crystal without any dislocation with (b) the same crystal with a dislocation.

  Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 50/230 I.2.5.

Figure

  Figure I.44: Schematic representations of vacancy and interstitial diffusion mechanisms.

Figure

  Figure I.45 and Figure I.46 show relaxed SiGe interdiffusivity and strained SiGe interdiffusivity versus Ge content at different temperatures and published by Dong et al. group.

Figure I. 45 :

 45 Figure I.45: Calculated SiGe interdiffusivity versus Ge content and comparison with various experimental data from literature. From [Dong12].

Figure I. 46 :

 46 Figure I.46: Comparison of SiGe interdiffusivity versusGe content in a SiGe layer on a fully relaxed Si 0.60 Ge 0.4 substrate under fully relaxed and fully strained conditions and at two different temperatures. From[START_REF] Daillant | X-ray and neutron reflectivity: principles and applications[END_REF].

Figure I. 47 :

 47 Figure I.47: Schematic view of the tetrahedron structure of thermal SiO 2 .

ȃ

  Oxidation states of Si as a function of depth into the interfacial layer have been investigated by X-ray Photoemission Spectroscopy (XPS) [Himpsel88][Rochet97][Oh01][Hirose07]. Oh et al. showed that Si 1+ and Si 2+ are located in the first layer of the interfacial layer while Si 3+ species location extends into the second and third layers. ȃ The density of thermal SiO 2 have been measured by X-Ray Reflectivity (XRR) [Ajawi96][Kosowsky97]. They

Figure I. 49 :

 49 Figure I.49: Schematic representation of the condensation technique.

  Figure I.51: STEM-EDX picture of the transition between the SGOI and SOI regions as well as the future location of the STI. From [Gourhant14].

Figure I. 50 :

 50 Figure I.50: FD-SOI process flow showing co-integration of SGOIbased pMOS and Si-based nMOS. The figure does not show intermediary process steps, such as SiN deposition, for clarity considerations.

Figure

  Figure II.1: Picture of a RTO tool in open situation. The honeycomb lamp head is on top while the processed wafer is below.

  Figure II.2.

FigureFigureψ

  Figure II.5: Illustration of ellipsometry principles. θ is the angle of incidence. The sample is optically described by its Jones matrix. Δ and Ψ are the ellipsometry angles.

  Photoelectron (1s) total energy

  Figure II.8 shows an experimental spectrum with the associated background and the peaks extracted from the fitting procedure.

Figure

  Figure II.8: Spectrum around the energy of the Ge3d XPS peak of a Si 0.7 Ge 0.3 sample oxidized at 1100°C for 10 seconds by means of dry RTO. The measured oxide thickness is 75 Å. The envelope refers to the fitted spectrum.

  FigureII.11: (a) 1D lattice fringes of a HRTEM image,(b) contrast of the HRTEM image and of the reference area and (c) phase substraction that allows to retrieve the displacement u(x).

Figure

  Figure II.12: (a) HRTEM image of a fully strained SiGe epitaxial layer on a SOI substrate. (b) Diffractogram of the HRTEM image showing the selected g vectors and the masked areas around these frequencies. (c) and (d) are phase maps associated with the two g vectors and showing the chosen reference area in the unstrained Si substrate.And (e)

Figure

  Figure II.13: Schematic illustration of (a) the scattering geometry of θ-2θ reflectivity and (b) bilayer interference.

Figure

  Figure II.15: Nuclear scattering lengths as a function of the atomic number Z. From [pd.chem].

Figure

  Figure II.18: Kinematic factor as a function of the scattering angle for O, Si and Ge target atoms and a He + ion beam.

Figure

  Figure II.19: Schematic illustration of the channeling mechanism showing the formation of a shadow cone behind the first atomic layer.

  Simulation of MEIS spectra was performed using a homemade program developed by Denis Jalabert from INAC [Jalabert17]. This program is based on the mechanisms and equations introduced in the former sections. The program inputs are: ȃ General inputs which are fixed in this work: • Particle type (He+) • Ion energy (100 keV) • Detector resolution (0.6%) ȃ Angle configuration: • Incident angle • Detection angle, which is also the scattering angle because it takes the incident angle as the reference Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies 80/230 ȃ Inputs for each layer of the model:

  stricte ne vous guérit pas. Non content d'être malade, on veut encore contaminer les autres, leur passer nos miasmes. » Oxidation Kinetics and SiGe Composition Evolution III.A.5. Correlation Between SiGe Oxidation Kinetics and Oxide Density III.A.6. Conclusion of Section III.A.

Figure

  Figure III.1: The Deal and Grove model for the oxidation of silicon. Adapted from [Deal65].

Figure

  Figure III.2: Experimental oxidation rate in dry environment and at various temperatures. Solid lines are the oxidation rate predicted by the Deal and Grove model and dotted lines are experimental data. From [Massoud85].

Figure

  Figure III.3: Schematic illustration of interstitials injection during oxidation of Si and SiGe in dry and wet oxidation conditions.

Figure

  FigureIII.4: Experimental and fitted (004) rocking curves of a Si 0.7 Ge 0.3 layer after epitaxy on bulk Si and after dry RTO at 1100°C for 40 seconds. Curves after RTO are vertically shifted for clarity considerations. The Si peak located at 0° was cut on experimental curves.

  apparatus are presented in Section II.4.1. and results are shown and discussed in Section III.A.5..

Figure

  Figure III.5: Oxide thickness measured by ellipsometry along the radius of the wafer for two oxidized samples.

  Figure III.6: Ge3d spectra of a Si 0.7 Ge 0.3 film oxidized by means of dry RTO at 900°C for 20 seconds for different collection angles (indicated on the Figure).

Figure

  Figure III.7: Chemical profile of the same sample reconstructed from Ge3d, Si2p, O1s and C1s spectra.

Figure

  FigureIII.8: (a) Typical temperature-time profile of a dry RTO process showing at which moment oxidation was stopped to perform XPS measurements. And (b) Ge3d spectra acquired after each key step of the SiGe epitaxy-cleaning process-dry RTO sequence.

  Figure III.9: Ge3d XPS spectra acquired after dry RTO of SiGe layers at (a) 1100°C, (b) 1000°C and (c) 900°C.

Figure

  Figure III.10: Ge3d XPS spectra acquired after dry RTO of SiGe layers with initial (a) 10% Ge concentration and (b) 30% Ge concentration.

Figure

  Figure III.11: Areas of the fitted Ge3d 4+ peak for all SiGe samples oxidized by dry RTO. Dotted lines are linear trend lines for each initial Ge concentration. Full lines are for the temperature ramp up and are plotted over the full range of oxidation temperature for comparison only.

Figure

  FigureIII.12: (a) Ge3d spectra acquired on a Si 0.7 Ge 0.3 layer which underwent the temperature ramp up until 900°C of dry furnace oxidation, and (b) which underwent a full dry furnace oxidation at 900°C for 540 seconds. And (c) Ge3d spectra acquired on Si 0.7 Ge 0.3 layers oxidized by means of ISSG oxidation performed at three different temperatures.

Figure

  Figure III.14: Experimental oxide thickness as a function of time of Si 0.7 Ge 0.3 oxidized by means of dry RTO at 900°C along with curves fitted with the two methods described in text.

Figure

  Figure III.16: Oxide thickness as a function of oxidation time for dry furnace oxidation and dry RTO of Si at 900°C.III.A.4.1.4. Activation Energies of Oxidation of SiGeAs shown Figure III.17, the parabolic constants α of dry RTO and ISSG oxidation both follow Arrhenius laws (only one oxidation temperature was investigated for dry furnace oxidation, which does not allow extraction of an activation energy). Extracted activation energies are shown in TableIII.2.

  Figure III.18: Illustration of the GRE for the case of Si 0.7 Ge 0.3 oxidized at 1000°C after 90 seconds. The oxide thicknesses mentioned in text for illustration are 67 Å and 93 Å. Arrows are tangents to the curves and represent the growth rates of Si 0.7 Ge 0.3 and Si at these oxide thicknesses.

Figure

  Figure III.19: GRE as a function of the initial Ge concentration for (a) dry RTO and furnace oxidation, and (b) ISSG oxidation of SiGe. The influences indicated by numbers on the figure are explained in text.

Figure

  FigureIII.20: SIMS profiles of the Ge concentration below the oxide after dry RTO of (a) Si 0.7 Ge 0.3 at 900°C, (b) 1000°C, (c) 1100°C, and of (d) Si 0.9 Ge 0.1 at 900°C and (e) at 1100°C. For clarity considerations, the x axis is shifted so that the SiGe-oxide interfaces are all close to zero.

Figure

  FigureIII.21: SIMS profiles of the Ge concentration below the oxide after (a) dry furnace oxidation of Si 0.7 Ge 0.3 at 900°C, (b) ISSG oxidation of Si 0.7 Ge 0.3 at 850°C, (c) and ISSG oxidation of Si 0.9 Ge 0.1 at 1050°C. For clarity considerations, the x axis is shifted so that the SiGe-oxide interfaces are all close to zero.

  III.A.4.

Figure

  Figure III.22: GRE of dry furnace oxidation, dry RTO and ISSG oxidation of SiGe as a function of the Ge concentration at the oxidation interface.

  Measurement of absorption spectra: absorption of several samples was measured in Total Electron Yield (T.E.Y.) and Total Fluorescence Yield (T.F.Y.) modes between 520 eV and 550 eV. Differences between the two modes are introduced in Section II.4.1.5. in terms of experimental set-up and will be discussed in Section III.A.5.1.2.. ȃ Normalization: the following linear normalization allows extraction of the imaginary part of the O scattering factor because absorption is proportional to the imaginary part of the scattering factor f 2 : f E aI E b 2 ( ) = ( )+ , where I(E) is the absorption spectrum as a function of the X-Ray energy E, and a and b are constants. a and b are calculated using two tabulated values for the O scattering factor respectively below and above the absorption edge (here 520 eV and 550 eV) [CXRO]:

Figure

  Figure III.23 shows the O scattering factors retrieved from measurement of absorption spectra of three different samples in T.E.Y. and T.F.Y. modes. The samples are Si and Si 0.7 Ge 0.3 films oxidized by means of dry RTO.While the T.E.Y. mode probes the top few nanometers of the oxide (between 1 nm and 5 nm), the T.F.Y. mode probed depth is larger than 10 nm and thus probes all O atoms of the oxide, including the O atoms at the interface with the Si(Ge) layer.

Figure

  Figure III.23: Imaginary part of the O scattering factor around its K-absorption edge measured in (a) T.E.Y. and (b) T.F.Y. modes. The O scattering factor of the Si sample oxidized at 1100°C and measured in T.E.Y. mode is also plotted in (b) for comparison. In order to discuss further about this pre-edge "bump", the O scattering factors of the Si sample oxidized at 1100°C are compared with the ones of O atoms in the bulk of the oxide and at the interface with Si measured by Muller at al. [Muller99]. Muller et al. used a microscopy technique (Electron Energy Loss Spectroscopy: EELS) to probe the O scattering factor at the atomic scale as a function of the position with respect to the Si-oxide interface. The O scattering factors spectra published by Muller et al. were normalized following the procedure of SectionIII.5.1.1. and shifted in the x axis so that the one of the bulk of the oxide coincides at 533.6 eV with the one that we measured (such a shift is expected because of different experimental set-up and thus different absolute energy measurements).

  T.F.Y. and f 2-T.E.Y. are imaginary parts of O scattering factors of the Si sample acquired in T.F.Y. mode and T.E.Y. mode respectively, f 2-interface-Muller is the O scattering factor at the interface published by Muller et al., and z a coefficient.

  LeGoues et al. and Kilpatrick et al. showed by XPS that the SiGe-oxide interface has more Si atoms in intermediate oxidation states than the Si-oxide interface[LeGoues89'][Kilpatrick08]. In other words, there are more Si sub-oxides at the SiGe-oxide interface. The sharp absorption peak at 535 eV arises from O-O scattering, so that we can expect O atoms in Si sub-oxides to have scattering factors close to the one at the interface published by Muller et al.. Our results are thus consistent with a higher quantity of Si sub-oxides at the SiGe-oxide interface observed byLeGoues et al. and Kilpatrick et al.

Figure

  Figure III.24: Imaginary part of O scattering factors of the Si sample around the K-absorption edge and measured in T.E.Y. and T.F.Y. modes, as well as the ones measured by Muller et al. in the bulk of thermal oxide on Si and at the Si-thermal oxide interface [Muller99].

  Figure III.26 as a function of the oxidation temperature. The effects of the initial Ge concentration, of the process type and of the oxidation temperature are summarized below and indicated by arrows on Figure III.26:

Figure

  Figure III.25: XRR spectra of thermal oxide-SiGe samples. Dots are experimental XRR spectra acquired at 500 eV and lines are results of fits.

  As shown Figure III.29, the diffusion barrier modulation ΔE increases with the oxide density. There seems to be a linear increase of the diffusion barrier for the case of SiGe (see trend line on Figure II.29), in agreement with Cui et al. who assumed that the diffusion barrier modulation is proportional to the density [Cui09]. Due to the fact that the density variation due to the temperature is included into the activation energy E and not in ΔE, we cannot compare densities of oxides grown in Si with the ones of oxides grown on SiGe on Figure III.29.Equation III.17

Figure

  Figure III.29: Evolution of the modulation of the diffusion barrier of O 2 through thermal SiO 2 as a function of the oxide density.

Figure

  Figure III.30: Oxidation temperature-time charts for oxidation of Si 0.7 Ge 0.3 . Isolines of oxide thicknesses for both dry RTO and ISSG oxidation are displayed on chart.

Figure

  Figure III.32: Schematic illustration of the target specifications of the example of how the charts of Figure II.30 and Figure III.31 can be used (see text)

Figure

  Figure III.34: GRE as a function of the Ge concentration at the oxidation interface. Dots are experimental data and the line is the fitted exponential function. ȃ Ge injection below the SiGe-oxide interface: As shown in Section IV.A.3., there is no significant Ge loss during oxidation and all Ge atoms are repelled below the oxide. The flux of Ge atoms injected at the oxidation interface towards the SiGe film is calculated at each step based on the consumption of Si atoms by the oxidation reaction.

Figure

  Figure III.35: Oxide thickness as a function of the oxidation time for dry RTO of Si at 1000°C. Lines are simulations with starting oxide thicknesses indicated in the legend and squares are experimental data.

Figure

  Figure III.36: Experimental and simulated oxide thicknesses as a function of the oxidation time of dry RTO of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 at (a) 900°C, (b) 1000°C and (c) 1100°C. And experimental and simulated oxide thicknesses as a function of the oxidation time of ISSG oxidation of Si at (a) 850°C, (b) 950°C and (c) 1050°C. Dots are experimental data and lines are simulations.

Figure

  Figure III.37: Experimental SIMS and simulated Ge concentration profiles.

Figure

  Figure III.39: Simulated Ge concentration at the interface after it reached the plateau, and plotted as a function of the initial Ge concentration and of the oxidation temperature. (a) is for dry RTO and (b) is for ISSG oxidation.

  « D'une certaine façon, être vivant ne s'atteint que par ce triple combat: contre les forces de gravité en nous -la paresse, la fatigue, la quête du repos ; contre l'instinct de répétition -le déjà-fait, le connu, le sécurisant ; et enfin contre les séductions du continu -tous les développements durables, le réformisme ou ce goût très fréole de la variation plaisante, du pianotement des écarts autour d'une mélodie amusante. » Alain Damasio, La Horde du Contrevent Chapter IV -Strain and Crystal Quality of SGOI Films Formed by the Condensation Technique 137/230 Chapter IV -Strain and Crystal Quality of SGOI Films Fabricated by the Condensation Technique and Crystal Defects of SGOI Layers IV.A.5. Characterization of SGOI Composition and Crystal Quality by MEIS IV.A.6. Conclusion of Section IV.A.

  Figure IV.1: Condensation model.

Figure

  FigureIV.2: (a) The full line is the Ge concentration of the SGOI layer as a function of its thickness for the both C1 and C2 and the symbol is the SGOI thickness and Ge concentration target point. (b) The full line is the oxide thickness as a function of the Ge concentration of the SGOI layer for C1 and C2 and the symbols are target oxide thicknesses.

Figure

  Figure IV.3: Evolution of the melting temperature of SiGe during condensation and of the condensation temperature which is decreased stepwise as soon as the melting temperature is only 50°C above the condensation temperature.

  Figure IV.4: Schematic illustration of the condensation-oxide removal-condensation sequence for the fabrication of high-Ge content SGOI films. Here x<y<z.

  Figure IV.5 shows WDXRF intensities of a reference sample, i.e. after SiGe epitaxy, and of the same sample after condensation by means of dry RTO at 900°C and at 1100°C. As peak intensities only are compared, no absolute energy calibration is required.

Figure

  FigureIV.5: (a) Ge K-α and (b) L-α lines WDXRF spectra for a sample before condensation (Reference) and after condensation by means of dryRTO at 900°C and 1100°C.

FigureFigureFigureFigure

  Figure IV.6: (a) HRTEM image of a post epitaxy sample with 30% Ge concentration, (b) in-plane and (c) out-of-plane deformation maps, and (d) depth profile of the out-of-plane deformation measured by GPA.

Figure

  Figure IV.10: Compressive in-plane strain of SGOI layers measured by Raman Spectroscopy and GPA and plotted as a function of the Ge concentration.

  Figure IV.11: Compressive in-plane strain of SGOI layers measured by Raman Spectroscopy as a function of the Ge concentration. Dots are experimental data and dotted lines are guides for the eye. IV.A.4.1.3. Comparison with Literature Takagi et al. used a condensation process based on dry furnace oxidation to fabricate SGOI films with Ge concentrations up to 100% [Takagi10]. They also found that the compressive strain of the SGOI layer increases with the Ge content until a peak strain value. Above the Ge concentration that corresponds to this peak value, the compressive in-plane strain decreases. Figure IV.12 shows the compressive in-plane strain values published by Takagi et al. and obtained with their best condensation process, and values from this work. We observe that the "bell shapes" of the evolution of the compressive in-plane strain as a function of the Ge concentration are very similar, with a maximum compressive strain of 2.2% at 60% Ge concentration for Takagi et al. and 2.1% at 63% Ge concentration for this work.

Figure

  Figure IV.13 plots the Ge concentration of SGOI films fabricated by condensation as a function of its thickness, for both this study and Takagi et al. one. Samples that are fully strained are indicated on Figure IV.13, and the SGOI57 sample, which is close to the onset of strain relaxation and likely slightly relaxed, is also indicated. The most common equilibrium models are Matthews and Blakeslee one and People and Bean one, and are also plotted on Figure IV.13 for comparison [Matthews74][People85]. With fully strained SGOI layers above Matthews and Blakeslee critical thickness, both our study and Takagi et al. disagree with this model. Our study seems to be below People and Bean model, while Takagi et al. had fully strained SGOI films above People and Bean critical thickness. In summary, neither Takagi et al. data nor our study agrees with the most accepted equilibrium models of epitaxial SiGe-Si systems.

Figure

  Figure IV.13: Ge concentration as a function of the SGOI thickness. This study experimental data and Takagi et al. experimental data [Takagi10] are plotted. The critical thickness of Matthews and Blakeslee model and the one of People and Bean model are also plotted as dashed lines [Matthews74][People85].

Figure

  Figure IV.15: Schematic illustration of the condensation process based on low temperature ISSG oxidation to fabricate SGOI films with a Ge concentration of 60%. The thermal oxide is not shown for clarity considerations.

Figure

  Figure IV.17: 5 μm x 5 μm AFM images of the surface of (a) SGOI36, (b) SGOI45, (c) SGOI57, (d) SGOI63 and (e) SGOI77 films.

FigureFigureFigure

  Figure IV.18: (a) HRTEM image of the Reference sample, and (b) and (c) numerical Moiré patterns of (111) planes. (a) (b) (c)

Figure

  Figure IV.20: (a) HRTEM image of the SGOI57 sample, and (b) and (c) numerical Moiré patterns of (111) planes. Triangles indicate where missing half planes are observed and the insets of (b) and (c) are magnifications of the area where some of these missing half-planes are observed.

Figure

  Figure IV.20: Simulated MEIS spectra of a thermal oxide-SGOI stack where the contribution of each element of each layer is shown.

Figure

  Figure IV.21: MEIS 2D map measured in channeling conditions with an incident angle of 45° and around a central scattering angle of 90°C. The arrows indicate the main peaks: (1) Ge peak of the SiGe layer, (2) Si peak of the thermal oxide, (3) Si peak of the SiGe layer appears lower in intensity (blue) because of the channeling conditions, and (4) Si peak of the BOx layer.

Figure

  Figure IV.23: Ge concentration profiles extracted from random MEIS spectra. Profiles are plotted as a function of depth taking the BOx-SGOI interface as zero and the x axis towards the surface. (a) is for the "Reference" sample, (b) for the "Low Ge" one and (c) for the "High Ge" one.

  from 45°F igure IV.24: Calculated angular tolerance and channeling angle deviation from 45° for a fully strained SiGe layer and plotted as a function of the Ge concentration.

Figure

  Figure IV.25: Ge concentration and χ min profiles extracted from random and channeling MEIS spectra. Profiles are plotted as a function of depth taking the BOx-SGOI interface as zero and the x axis towards the surface. (a) is for the "Reference" sample, (b) for the "Low Ge" one and (c) for the "High Ge" one.

  Figure IV.26 summarizes the evolution of this crystal quality during condensation.

Figure

  Figure IV.26: Schematic illustration that summarizes the evolution of the amount of atoms displaced from lattice sites. (a) is for an epitaxial SiGe layer on SOI, (b) for a low Ge content SGOI film, and (c) for a high Ge content SGOI film. Key changes of the crystal quality from the epitaxial sample to the low Ge content SGOI film and finally to the high Ge content SGOI fims are indicated on the right side.

Figure

  Figure IV.27: Integration of SGOI channels into pMOS active areas showing the STI etch process step after condensation. Intermediary process steps such as SiN film removal are not shown for clarity considerations. SGOI strain profiles measured at the edge of active areas by Nano Beam Electron Diffraction (NBED) [Coo-per16] and Dark Field Electron Hololography (DFEH -a strain characterization technique developed at CEMES laboratory [Hÿtch08]) were compared with FEM simulations based on the elastic theory [Boureau16]. As shown in Figure IV.28, the distance over which the SGOI film is partially relaxed after STI etch is much longer than the one predicted by the elastic theory. Usuda et al. also showed by Raman spectroscopy that relaxation of a SiGe film grown on Si agrees with the elastic theory while the SGOI films relaxation distance is longer [Usuda13].

FigureFigure

  Figure IV.29: Schematic view of specimens preparation.

Figure

  FigureIV.30: (a) Ge2p spectrum of "SGOI 900°C", (b) Ge3d spectrum of "SGOI 900°C annealed", (c) Ge2p spectrum of "SGOI 1100°C", and (d) Ge3d spectrum of "SGOI 1100°C".

Figure

  Figure IV.31: Si2p spectra of the (a) "SOI", (b) of the "SGOI 900°C annealed", and (c) of the "SGOI 1100°C".

  Three specimens were studied and are illustrated FigureIV.32:

FigureFigure

  Figure IV.32: Schematic representation of (a) "Reference epi", (b) "SGOI 1100°C" and (c) "SGOI 900°C" samples.

  , ρ, b O , b Si , b Ge , M O , M Si , M Ge and x are the Avogadro number, the mass density of either Si 1-x Ge x or SiO 2 , the scattering lengths of O, Si and Ge, the molar masses of O, Si and Ge and the Ge content.

Figure

  Figure IV.34: Si, Ge and O arbitrary concentration levels and Cl absolute concentration measured by ToF-SIMS in the "Reference epi" sample.

  β = , , and where u α is the displacement of the lattice point along x α .

.Figure A. 1 :

 1 Figure A.1: Schematic representation of stress components on surfaces of an infinitesimal cube. From [Sun09].

Figure B. 1 :

 1 Figure B.1: Schematic illustration of SIMS principles.

  Figure C.1: Illustration of the process on which the XRF technique is based. The process follows the sequence A-B-C.With A: A core electron of the K-shell is emitted into vacuum due to the incoming X-Ray beam. B: An electron of the L-shell filled in the hole of the K-shell. C: the energy set free by process B induces emission of a X-Ray photon with a specific wavelength corresponding to the Kα spectral line.

  Figure D.1: Illustration of Bragg's conditions for diffraction by a family of planes {hkl}.

  Figure D.2: JVX7300 diffraction setup.

Figure D. 3 :

 3 Figure D.3: Experimental and fitted rocking curves of an epitaxial SiGe layer on a bulk Si substrate. The Si peak is cut on the experimental rocking curve.

  Figure F.1: Schematic view of a Transmission Electron Microscope.

Figure I. 1 :

 1 Figure I.1: Illustration schématique d'un MOSFET de type p.

  t et τ sont l'épaisseur d'oxyde, le temps d'oxydation et une constante, et où B/A, la constante linéaire, et B, la constance parabolique, dépendent des paramètres physiques du modèle et du procédé d'oxydation.Nous pouvons remarquer que lorsquet A B  2 4, la loi cinétique est correctement décrite après approximation par l'équation parabolique suivante:

  Trois concentrations en Ge (0%, 10% et 30%) ont été utilisées afin d'identifier l'influence de la composition du film SiGe sur la cinétique d'oxydation. Les films SiGe contraints en compression par le substrat, tel qu'illustré Figure I.2, et déposés par épitaxie, présentent des épaisseurs de l'ordre de 20 nm. Pour chacun des procédés d'oxydation décrit en partie I.2., différentes températures et durées d'oxydation ont été utilisées. L'épaisseur des oxydes thermiques ainsi formés et la composition des films SiGe en fonction de la profondeur ont été déterminés par ellipsométrie et par spectrométrie de masse à ionisation secondaire (SIMS : Secondary Ion Mass Spectroscopy) respectivement.

Figure II. 2 :

 2 Figure II.2: Epaisseur d'oxyde en fonction du temps d'oxydation pour (a) l'oxydation sèche en four à 900°C, (b) l'oxydation sèche par RTP à 900°C et (c) pour l'oxydation ISSG à 950°C de couches Si, Si 0.9 Ge 0.1 et Si 0.7 Ge 0.3 . Les points sont les épaisseurs expérimentales.

Figure II. 3 :

 3 Figure II.3: Carré de l'épaisseur d'oxyde en fonction du temps d'oxydation pour (a) l'oxydation sèche en four à 900°C, (b) l'oxydation sèche par RTP à 900°C et pour (c) l'oxydation ISSG à 950°C de couches Si, Si 0.9 Ge 0.1 et Si 0.7 Ge 0.3 . Les points sont les épaisseurs expérimentales et les lignes sont les résultats des ajutements des lois paraboliques aux points expérimentaux.

Figure

  Figure II.5: GRE en fonction de la concentration initiale en Ge pour (a) l'oxydation sèche en four et par RTP et pour (b) l'oxydation ISSG. L'influence de (1) la concentration initiale en Ge, de (2) la manière dont l'énergie thermique est apportée au système et de (3) la température est indiquée par des flèches.

Figure II. 6 :

 6 Figure II.6: Profils en concentration en Ge après oxydation sèche par RTP à (a) 900°C, (b) 1000°C and (c) 1100°C. L'axe des x est systématiquement décalé afin de positionner l'interface oxyde-SiGe à la profondeur 0.II.1.7. Dépendance de la cinétique d'oxydation à la concentration en Ge à l'interface d'oxydationNous pouvons observer en Figure II.7 que le GRE augmente avec la concentration en Ge à l'interface oxyde-SiGe pour chacun des trois procédés d'oxydation. Pour l'oxydation sèche par RTP, les influences de la concentration initiale en Ge et de la température discutées en partie II.5. sont ainsi réduits à une unique influence de la concentration en Ge à l'interface oxyde-SiGe.

Figure

  Figure II.7: GRE en fonction de la concentration en Ge à l'interface oxyde-SiGe. II.1.8. Conclusion de la partie II.1.

Figure II. 8 :

 8 Figure II.8: Evolution de la densité des oxydes crûs sur Si et SiGe en fonction de la concentration initiale en Ge et des paramètres du procédé d'oxydation. Les influences de la concentration initiale en Ge, du type de procédé et de la température sont indiquées par des flèches (voir texte).

Figure

  Figure II.10: Evolution du GRE en fonction de la densité de l'oxyde.

Figure

  Figure II.11: Evolution de la modulation de la barrière à la diffusion d'O 2 dans SiO 2 en fonction de la densité du SiO 2 .

  Ce chapitre a permis d'apporter davantage de compréhension concernant les mécanismes et la cinétiqueII.2.6. Conclusion de la partie II.2. Oxidation of SiGe for Advanced CMOS FD-SOI Technologiesd'oxydation du SiGe qui doivent être finement maîtrisés pour une utilisation industrielle du procédé de condensation.

  et défauts cristallins des films SGOI III.1.1. Protocole expérimental Deux types de procédé de condensation, dont les profiles en température sont représentés en Figure III.1, sont discutés dans cette partie III.1. : ȃ L'oxydation sèche par RTP à haute température : la température de condensation est réduite par paliers de 100°C, en commençant à 1100°C, afin que cette dernière demeure toujours inférieure à la température de fusion du SiGe qui, elle, diminue au cours du procédé du fait de l'augmentation de la concentration en Ge. Ce type de procédé conduit à l'élaboration de films SGOI relativement homogènes en composition grâce à une forte diffusion du SiGe (voir partie II.1.6.). ȃ L'oxydation ISSG à basse température : l'oxydation ISSG à 850°C d'un film Si 0.7 Ge 0.3 conduit à la formation sous l'oxyde thermique d'une couche enrichie à 60% de Ge par effet « pile-up » (voir partie II.1.6.). La faible température du procédé ne permet pas la diffusion à l'interface Si 0.7 Ge 0.3 -Si. La condensation est arrêtée lorsque la couche enrichie à 60% de Ge constitue la totalité du film SGOI tel qu'illustré en Figure III.2.Les temps des procédés de condensation ont été calibrés grâce aux résultats de cinétique d'oxydation présentés en Chapitre II.

Figure

  FigureIII.1: Evolution de la température de fusion du SiGe au cours du procédé de condensation et des températures des deux procédés de condensation décrits dans le texte. Pour l'oxydation sèche par RTP à haute température, la température de condensation est réduite de 100°C lorsque l'écart entre cette dernière et la température de fusion du SiGe n'est plus que de 50°C.

Figure

  Figure III.4: Evolution de la contrainte des films SGOI fabriqués par un procédé de condensation reposant sur l'oxydation sèche par RTP à haute température (déjà présentées en Figure III.3), ainsi que les données publiées par Takagi et al. sur les contraintes de films SGOI fabriqués sur substrats SOI et sSOI par un procédé de condensation reposant sur une oxydation sèche en four à haute température [Takagi10].Cependant, il est à noter que les épaisseurs des films SGOI fabriqués par Takagi et al. sont très différentes de celles des films SGOI de cette étude. Or, on pourrait logiquement s'attendre à une forte influence de l'épaisseur, comme c'est le cas pour les films SiGe épitaxiés sur Si. Cette observation suggère des mécanismes de relâchement de la contrainte propres aux films SGOI.III.1.4. Discussion autour des mécanismes de relâchement de la déformationCertains des films SGOI de cette étude et de ceux fabriqués par Takagi et al. sont totalement contraints alors que les contraintes des films SiGe de même épaisseur et de même composition sont partiellement relâchées d'après les modèles de référence de Matthews et dePeople [Matthews74][People85]. Le relâchement plastique des contraintes des films SGOI n'obéit donc pas aux mécanismes de relâchement des contraintes des films SiGe sur Si.

FigureFigure

  Figure III.5: Evolution de la contrainte des films SGOI fabriqués par condensation à haute température (par oxydation sèche par RTP) et à basse température (par oxydation ISSG).III.1.5. Révélation de défauts cristallinsLes images de microscopie à force atomique (AFM : Atomic Force Microscopy), dont certaines sont représentées en Figure III.6, révèlent la présence de motifs particuliers en surface des films SGOI présentant des concentrations en Ge supérieures ou égales à 63% et fabriqués par condensation à haute température. Ces motifs sont constitués de lignes orientées dans les directions <110> et sont révélatrices de dislocations dans les plans {111} du cristal de SiGe[Gunji11]. Les films SGOI dont les contraintes sont relâchées avec certitude (dont les concentrations en Ge sont 63% et 77%) présentent donc des dislocations.

Figure

  Figure III.7: (a) Représentation de l'empilement dont le spectre MEIS simulé est représenté en (b). La contribution de chaque couche est indiquée en (b).Lorsque le faisceau d'ions parallèle est orienté dans une direction cristallographique particulière d'une couche cristalline de l'échantillon, ce dernier peut théoriquement traverser la totalité de la couche sans rétrodiffusion. C'est l'effet de canalisation, illustré en Figure III.8, qui du point de vue de la mesure diminue le nombre d'ions rétrodiffusés par les atomes du cristal. Le facteur de réduction du nombre d'ions rétrodiffusés par effet de canalisation est appelé χ min . Ce facteur est le pourcentage d'atomes déplacés des sites du cristal, qui peuvent être des atomes en sites interstitiels ou situés au coeur de dislocations par exemple. Le χ min est égal à 100% pour un matériau totalement amorphe et à 5% pour un cristal de Si d'excellente qualité. En résumé, ce facteur est un indicateur de la qualité cristalline.

Figure

  Figure III.8: (a) Représentation du principe de la canalisation dans un cristal de SiGe et (b) spectres MEIS simulés du même empilement qu'en Figure III.8, en conditions normales (non canalisées) et en conditions de canalisation.

Figure•

  Figure III.9: Spectres expérimentaux (notés « Exp ») et spectres simulés (notés « Simu ») par ajustement aux spectres expérimentaux pour les échantillons (a) « Référence », (b) « Faible Ge » et (c) « Fort Ge ». Les spectres obtenus en conditions de canalisation sont aussi représentés et identifiés dans la légendeIII.2.4. Qualité cristalline des films SGOINous soulignons tout d'abord que la composition des films SiGe et SGOI est relativement homogène en profondeur (cette discussion, reposant sur l'analyse des spectres MEIS hors conditions de canalisation seulement, n'est pas détaillée ici). La Figure III.10 résume l'évolution au cours de la condensation du profil en profondeur du nombre d'atomes déplacés des sites cristallins, représenté par le facteur χ min , depuis l'état avant condensation (« Référence »), en passant par un film SGOI faiblement contraint du fait d'une faible concentration en Ge (« Faible Ge ») et jusqu'au film SGOI fortement contraint et à forte concentration en Ge (« Fort Ge »). Nous avons choisi de tracer les profils en profondeur du χ min en prenant comme origine de l'axe des abscisses l'interface avec l'oxyde enterré et en allant vers la surface de l'échantillon pour les valeurs positives des x.

Figure

  Figure III.10: Illustration de l'évolution du nombre d'atomes déplacés des sites atomiques du cristal. (a) correspond à un film SiGe épitaxié sur SOI (« Référence »), (b) à un film SGOI de faible concentration en Ge (« Faible Ge ») et (c) à un film de forte concentration en Ge (« Fort Ge »). Les changements majeurs de la qualité cristalline, représentée par le facteur χ min , au cours de la condensation sont indiqués point par point à droite de chaque sous-figure. L'axe des abscisses correspond à la profondeur dans le film en partant de l'interface avec l'oxyde enterré (BOx) et en allant vers la surface. L'oxyde thermique n'est pas représenté sur les schémas des empilements sous chaque sous-figure.

  Afin de bénéficier d'un gain maximal sur la mobilité des trous dans le canal des transistors de type p, les films SGOI fabriqués par condensation doivent présenter des contraintes importantes et une excellente qualité cristalline.

Si BOx Si 1-x Ge x BOx SiGe SiO 2 BOx Si 1-y Ge y SiO 2

Table II .

 II 1: Comparison of some technical specifications of furnace oxidation and RTO. Dry oxidation: the atmopshere is pure O 2 or O 2 diluted in N 2 . ȃ Wet oxidation: H 2 O is introduced into the chamber. it is only used in furnace oxidation systems. ȃ In-Situ Steam Generation (ISSG): H 2 and O 2 are introduced into the chamber and react with each other inside it. ISSG is only used in RTO systems.We summarize TableII.2 the main thermal oxidation processes presented above.

	Oxidation	Number of processed	Oxidation	Typical working	Temperature ramp
	Technique	wafers in a single run	plateau duration	temperature range	up and ramp down
	Furnace oxidation	Batch (100 wafers)	From tens of mi-	From 800°C to	From 10°C/min to
			nutes to hours	1200°C	50°C/min
	Rapid Thermal	Single wafer	From a few	From 800°C to	From 10°C/second
	Oxidation (RTO)		seconds to a few	1200°C	to 50°C/second
			minutes
	Then, we have to distinguish oxidation processes depending on the gases introduced into the chamber. The
	three oxidation atmospheres used in this study are:
	ȃ Table II.2: Technology system along with the oxidizing gases intro-
		duced into the chamber used in the semiconductor industry.

Dry Oxidation: O 2 Wet Oxidation: H 2 O ISSG: H 2 + O 2 Furnace Oxidation Rapid Thermal Oxidation (RTO)

Ge-Ge Si-Ge Si-Si Si T2g

		2 000					520.9
		1 500
	Intensity (cnt)	500 1 000	283.3 296.8	418.6 408.9	457.6	489.7	505.8
		0
			300	400 Raman shift (cm -1 )		500

Table III .

 III 1: Standard Gibbs energies of formation of various phases of the Si-Ge-O ternary system at 1000 K. Adapted from [Barin77].

	Si

SiO 2 Ge GeO 2 O 2 Standard Gibbs energy of formation (kJ.mol -1 )

	-30.39	-981.78 -44.12 641.31	-220.93

  Costello84] and Medium Energy Ion Scattering (MEIS) experiments based on 16 O-18 O sequential oxidation

	III.2: Experimental oxidation rate in dry
	environment and at various temperatures. Solid
	lines are the oxidation rate predicted by the Deal
	and Grove model and dotted lines are experimental
	data. From [Massoud85].
	III.A.1.3.2. Revised Models of Oxidation of Si
	Nuclear reaction resonance [Rosencher79][Rochet84][Trimaille89], Secondary Ion Mass Spectroscopy
	(SIMS) [

  [Kazor95].

		Wet (fast) oxidation	Dry (slow) oxidation
		Interstitials
		SiO 2	SiO 2
	Oxidation of Si	Si Interstitials do not have	Si Interstitials diffuse away
		time to diffuse away and	from the oxidizing
		accumulate at the	interface
		oxidizing interface
	Oxidation of SiGe	SiGe SiO 2 Interstitials amount	SiGe SiO 2 Interstitials amount
		injected toward the	injected toward the
		substrate is lower	substrate is lower

  TableIII.2: Activation energies of the parabolic constant α of dry RTO and ISSG oxidation of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 (eV). III.17: Symbols are experimental parabolic constants α of dry RTO and ISSG oxidation of Si, Si 0.9 Ge 0.1 and Si 0.7 Ge 0.3 as a function of 1/k B T, and lines are results of fits with Arrhenius laws.

	1000
	)
	.s -1
	Parabolic constant α (Å 2	10 100		Dry RTO Si Dry RTO SiGe10% Dry RTO SiGe30% ISSG Si
				ISSG SiGe10%
	Figure 1	8	8.5 ISSG SiGe30% 9 1/k B T (eV -1 ) 9.5	10	10.5
					0.1	Si 0.7 Ge 0.3
	Dry RTO			2.29 ± 0.04	2.21 ± 0.05	2.06 ± 0.03
	ISSG Oxidation		0.95 ± 0.03	0.97 ± 0.03	0.99 ± 0.03

  TableIII.3: Total oxide thicknesses measured by ellipsometry and interfacial layer thicknesses calculated using Equation III.12 and Equation III.13. Samples were oxidized by means of dry RTO at the temperature indicated in the Table.

		1100°C	Si 0.7 Ge 0.3 1100°C	Si 0.7 Ge 0.3 900°C
	Total thickness	95.4 Å	138.8 Å	99.5 Å
	Interfacial layer thickness	26.5 Å	50 Å	46.3 Å

III.A.5.1.3. Methodology for Measurement of the Oxide Density by Fitting XRR Spectra

B.2. Chemical Bonding at the SGOI-BOx interface Inves- tigated by XPS IV.B.2.1. Introduction In

  . respectively. this section, we investigate the chemical bonds at the SGOI-BOx interface, which then influence the mechanical properties of the interface.Kim et al. and Suh et al. mentioned that Ge-O bonds are weaker than Si-O bonds [Kim17][Suh16]

IV.

  .2 and TableIV.3 respectively. Mass densities of all SiGe and SiO 2 layers measured by XRR were found to be in good agreement with theoretical ones (SiGe mass densities can be calculated according to Poisson theory, as described in Section I.2.2.3. and SiO 2 mass densities are discussed in details in SectionIII.A.5.)

	XRR	Reference epi	SGOI 1100°C	SGOI 900°C
	Mass density (g.cm -3 ) Thermal oxide		2.17	2.25
	SiGe top	3.24		3.78
	SOI/Si	2.27		2.23
	SGOI		2.97
	BOx	2.21	2.20	2.23
	Table IV.3: Neutron scattering length densities determined by fitting neutron reflectivity data.

  Schematic view of principles and apparatus of the AFM characterization technique.

	200/230							Study of Thermal Oxidation of SiGe for Advanced CMOS FD-SOI Technologies
								Feedback control
								PI
								Set point	z	piezo
	V	A		B		C		D
									Piezo tube
	lever Figure F.1: Optical	Sample	dither	
									Equation E.1

Les mécanismes et la cinétique de l'oxydation thermique du SiGe

	Introduction
	Chapitre I -La technique de condensation appliquée aux technologies CMOS
	FD-SOI
	Introduction au Chapitre I
	I.1. La technologie MOSFET, les films SGOI et la technique de condensation :
	limitations, potentiel, et problématiques
	I.2. Procédés d'oxydation
	Chapitre II -Introduction au Chapitre II
	II.1. Cinétique d'oxydation et évolution de la composition du film SiGe
	II.1.1. Le modèle de Deal et Grove
	II.1.2. Protocole expérimental
	II.1.3. Régime d'oxydation
	II.1.4. Méthode d'analyse: définition du Growth Rate Enhancement (GRE)
	II.1.5. Dépendance du GRE à la concentration initiale en Ge et à la température
	II.1.6. Evolution de la composition du film SiGe pendant l'oxydation
	II.1.7. Dépendance de la cinétique d'oxydation à la concentration en Ge à l'interface d'oxydation
	II.1.8. Conclusion de la partie II.1.
	II.2. Corrélation entre cinétique d'oxydation et densité de l'oxyde
	II.2.1. Protocole expérimental
	II.2.2. Densité des oxydes thermiques
	II.2.3. Dépendance de la densité de l'oxyde aux paramètres du procédé d'oxydation
	II.2.4. Dépendance de la cinétique d'oxydation à la densité de l'oxyde
	II.2.5. Dépendance de la diffusion de l'O 2 à la densité de l'oxyde
	II.2.6. Conclusion de la partie II.2.
	Conclusion du Chapitre II
	Chapitre III -

Déformation et qualité cristalline de films SGOI fabriqués par condensation

Introduction au Chapitre III III.1. Déformation et défauts cristallins des films SGOI III.1.1. Protocole expérimental III.1.2. Evolution de la déformation de films SGOI en fonction de la concentration en Ge III.1.3. Comparaison avec la littérature III.1.4. Discussion autour des mécanismes de relâchement de la déformation

Introduction au Chapitre I

  Dans ce premier chapitre, nous introduisons tout d'abord la brique élémentaire d'une majorité des systèmes électroniques : le MOSFET (Metal Oxide Semiconductor Field Effect Transistor). Ses limitations et l'intérêt que présente l'utilisation de canaux de type SGOI pour le MOSFET de type p (dont la conduction est assurée par les trous) sont soulignés. La technique de condensation et les deux problématiques de cette thèse, adressées dans les Chapitre II et Chapitre III respectivement, sont ensuite présentées. Enfin, les procédés d'oxydation thermique utilisés sont décrits.

I.

1. La technologie MOSFET, les films SGOI et la technique de condensation : limitations, potentiel et problématiques

  Si la technologie MOSFET a évolué sous de nombreux aspects depuis le premier circuit intégré commercialisé en 1971, ses principes de fonctionnement sont restés identiques. Ce travail portant sur le MOSFET de type p, représenté Figure I.1, nous nous concentrerons désormais sur ce dernier. La conductivité du canal entre la source et le drain, qui représente l'information binaire (« 0 » ou « 1 ») des systèmes logiques, est contrôlée par la tension appliquée à la grille.

BOx Si 1-x Ge x BOx SiGe SiO 2 BOx Si 1-y Ge y SiO 2 Homogénéisation de la composition par inter-diffusion Oxydation sélective du Si Système de chauffage Si SiGe

  Schéma de principe du procédé de condensation. BOx est l'abréviation de «Buried Oxide» pour «oxyde enterré», a Si , a ˔ et a || sont les paramètres de maille du Si relaxé, du SiGe dans le plan et du SiGe hors-plan. Deux principales caractéristiques sont associées aux procédés d'oxydation thermique : le type d'équipement et la chimie d'oxydation. Les types d'équipement utilisés pendant cette thèse sont :ȃ Le four : les gaz introduits dans la chambre sont chauffés par des résistances et ce type d'équipement permet de traiter plusieurs dizaines de plaquettes à la fois. ȃ Les systèmes de recuit rapide (RTP : Rapid Thermal Processing). Une unique plaquette est chauffée par le rayonnement émis par un réseau de lampes halogènes au tungstène situé au-dessus de la plaquette. Ce type d'équipement permet des montées et descentes en température bien plus agressives que le four (de l'ordre de 50°C/s pour le RTP contre 50°C/minute pour le four).

	𝑎𝑎 ∥ = 𝑎𝑎 𝑆𝑆𝑆𝑆	𝑎𝑎 ∥ = 𝑎𝑎 𝑆𝑆𝑆𝑆
	𝑎𝑎 ⊥ > 𝑎𝑎 𝑆𝑆𝑆𝑆	𝑎𝑎 ⊥
	𝑎𝑎 𝑆𝑆𝑆𝑆
	𝑎𝑎 𝑆𝑆𝑆𝑆
	Figure I.2: Les deux atmosphères oxydantes utilisées pendant cette thèse sont : I.2. Procédés d'oxydation
	ȃ O 2 pur à 740 Torr, appelée oxydation sèche. ȃ 33% H 2 + 66% O 2 à une pression totale de 7 Torr, appelée oxydation ISSG (In-Situ Steam Generation). Il a été démontré que cette atmosphère conduit à la génération de radicaux d'oxygène O* qui sont l'espèce oxydante
	principale de cette atmosphère [Kee00].
	Les procédés d'oxydation étudiés et combinant un de ces types d'équipement et une de ces chimies sont :
	• Oxydation sèche en four
	• Oxydation sèche par RTP
	• Oxydation ISSG par RTP, abrégée par la suite en oxydation ISSG

L'utilisation de cette technique dans un environnement industriel nécessite de relever deux défis : ȃ Maîtriser les mécanismes et la cinétique d'oxydation. Il est important d'étudier plusieurs procédés d'oxydation à la fois industriels et pertinents au regard des besoins technologiques actuels. Cette problématique fait l'objet du Chapitre II. ȃ Atteindre les plus fortes contraintes et qualités cristallines pour les films SGOI. Dans cette optique, la compréhension de l'évolution des contraintes et de la qualité des films SGOI en fonctions des paramètres du procédé de condensation se révèle indispensable pour l'optimisation de ce dernier. Cette problématique est adressée au Chapitre III.

Chapitre II -

Les mécanismes et la cinétique de l'oxydation thermique du SiGe Introduction au Chapitre II II.1. Cinétique d'oxydation et évolution de la composition du film SiGe II

  .1.1. Le modèle de Deal et Grove La « consommation » d'un volume de Si conduit à la croissance de 2.25 volumes de SiO 2 en raison de la différence de densité entre les deux matériaux. Quant à la cinétique de l'oxydation du SiGe, elle est généralement décrite par le modèle de Deal et Grove [Deal65]. Le modèle, illustré en Figure II.1, décrit l'oxydation par l'enchaînement des trois étapes suivantes :ȃ L'espèce oxydante est absorbée en surface de l'oxyde depuis l'atmosphère. ȃ L'espèce oxydante diffuse à travers l'oxyde en direction de l'interface oxyde-SiGe. Ce processus obéit à la loi de Fick. ȃ L'espèce oxydante réagit à l'interface oxyde-SiGe avec les atomes de Si pour former du SiO 2 .

		Gaz	Oxyde	SiGe
			𝑥𝑥
	𝐶𝐶 *	𝐶𝐶	𝐶𝐶 0
			𝐶𝐶 𝑖𝑖

𝑥𝑥 Figure II.1: Le modèle de Deal et Grove de l'oxydation thermique du silicium. C est la concentration en espèces oxydantes à une profondeur donnée et x l'épaisseur d'oxyde. Adapté de [Deal65].

Remerciements

Si

III.1.2. Evolution de la déformation de films SGOI en fonction de la concentration en Ge

Les contraintes (que nous assimilons à la contrainte relative au paramètre de maille du SiGe non contraint et non à celui du Si) des films SGOI fabriqués par oxydation sèche par RTP à haute température sont tracées en fonction de la concentration en Ge en Figure III.3. La courbe « pseudomorphique SiGe » correspond à la contrainte théorique d'un film SiGe dont le paramètre de maille dans le plan est exactement celui du Si. La contrainte expérimentale des films SGOI suit la courbe théorique jusqu'à une concentration en Ge au moins égale à 45%. Le cas du film SGOI à 57% de concentration en Ge est délicat car sa contrainte est très légèrement inférieure à la contrainte théorique. La contrainte continue néanmoins d'augmenter jusqu'à un maximum de 2.1% à 63% de concentration en Ge, puis diminue au-delà de cette concentration.

Figure III.3: Evolution de la contrainte des films SGOI fabriqués par un procédé de condensation reposant sur l'oxydation sèche par RTP à haute température.

La forme en « cloche » de cette évolution de la contrainte des films SGOI est très semblable à celle publiée par Takagi et al. pour le cas de films SGOI fabriqués par un procédé de condensation reposant sur l'oxydation sèche en four à haute température [Takagi10] (voir Figure III.4). Le type de procédé, four ou RTP, ne semble pas avoir d'effet sur l'évolution de la contrainte des films SGOI. Alors que la comparaison entre ces deux jeux de données pourrait laisser supposer que l'évolution de la contrainte est dépendante uniquement de la concentration en Ge, un troisième jeu de données, publié par Takagi et al. et portant sur la contrainte de films SGOI fabriqués sur substrats sSOI (substrats dont le SOI est contraint en tension) montre qu'il n'en est rien. En effet, la forme en « cloche » de l'évolution de la contrainte de ces films semble simplement décalée vers de plus fortes concentrations en Ge. Au vu de la comparaison de ces trois jeux de données, le paramètre semblant déterminer l'évolution de la contrainte, et en particulier l'existence d'une contrainte maximale, est la valeur de la contrainte elle-même, avec un maximum entre 2.2% et 2.5%.

III.1.3. Comparaison avec la littérature

Publications in Regular Journals

• Rozé, F., Gourhant, O., Blanquet, E., Bertin, F., Juhel, M., Abbate, F., Pribat, C. & Duru, R. (2017)

Talks in International Conferences

Publications in Conference Proceedings

KEY WORDS

The tremendous spread of electronic devices and networks into our day-to-day life has been enabled by the constant downscaling of transistors since the 60's. However, downsizing transistors has become increasingly difficult in the past few years and going to the nanometer scale brings new detrimental effects that have put power consumption and performances on quasi-plateaux for a few years. To overcome these limitations, high mobility channels based on new materials and new transistor architectures are being introduced. Ultrathin compressively strained SiGe-On-Insulator (SGOI) films benefit from the advantages of both the higher hole mobility of compressively strained SiGe as well as of the better electrostatic control of On-Insulator structures. The condensation technique is a CMOS-compatible technique that allows fabrication of such films with possibly high Ge content and high strain levels. The technique is based on Si-selective thermal oxidation of SiGe and concurrent SiGe diffusion between the thermal oxide and the buried oxide layer that acts as a diffusion barrier.

Two main challenges still need to be taken up for an efficient and optimized use of the condensation technique in an industrial environment: oxidation mechanisms and kinetics must be well controlled, and strain and crystal quality of the SGOI film must be as high as possible.

Firstly, this work bridges the gap between previous studies by covering various oxidation processes relevant to today's technological needs with a new and quantitative analysis methodology of oxidation kinetics. A correlation is established between the diffusivity of the oxidizing species that governs oxidation kinetics, the Ge concentration at the oxidation interface, and the oxide density measured by X-Ray Reflectivity on a synchrotron beamline.

Secondly, SGOI films with Ge concentrations up to 80% were fabricated by the condensation technique. The evolution of strain of SGOI films is discussed as a function of process parameters and strain energy levels. How the condensation technique alters the crystal quality, both at interfaces with oxides and in the bulk of the SiGe crystal, is evaluated by the Medium Energy Ion Scattering (MEIS) technique by using the channeling effect.

TITRE: Etude de l'oxydation thermique du SiGe pour application aux technologies CMOS FD-SOI avancées

RESUME

CMOS, canal à haute mobilité, SiGe, SiGe-On-Insulator (SGOI), contrainte, condensation, oxydation thermique, SiO 2

MOTS CLES

La réduction continue des dimensions des transistors depuis les années 60 est à l'origine de l'explosion des usages de l'électronique. Toutefois, la réduction des dimensions à l'échelle nanométrique s'accompagne de nouvelles difficultés qui tendent à limiter les gains des transistors en termes de performances et de consommation. Afin de surmonter ces obstacles et maintenir cette dynamique, des canaux à base de nouveaux matériaux à forte mobilité et de nouvelles architectures de transistors sont désormais utilisées ou à l'étude. L'intérêt de films SiGe contraint en compression sur isolant (SGOI: SiGe-On-Insulator) ultra-minces est double : ils bénéficient de la forte mobilité des trous du SiGe contraint en compression ainsi que du meilleur contrôle électrostatique des structures dites « sur isolant ». Des films SGOI présentant une forte concentration en Ge et une importante contrainte peuvent être fabriqués par une technique industrielle appelée condensation. Cette technique repose sur deux processus simultanés : l'oxydation thermique et sélective du SiGe (seul le Si est oxydé) et l'inter-diffusion du SiGe entre l'oxyde thermique et l'oxyde enterré qui se comporte comme une barrière à la diffusion.

L'utilisation de cette technique dans un environnement industriel nécessite de relever deux défis : maîtriser les mécanismes et la cinétique d'oxydation, et atteindre les plus fortes contraintes et qualités cristallines pour les films SGOI.

La cinétique de plusieurs procédés d'oxydation industriels et pertinents au regard des besoins technologiques actuels est étudiée à l'aide d'une nouvelle méthodologie d'analyse quantitative. Nous établissons une corrélation entre le coefficient de diffusion de l'espèce oxydante, qui détermine la cinétique d'oxydation, la concentration en Ge à l'interface d'oxydation, et la densité de l'oxyde mesurée par réflectivité de rayons X sur une ligne de synchrotron.

Puis, nous avons fabriqué des films SGOI présentant des concentrations en Ge jusqu'à 80%. Nous discutons l'évolution de la contrainte de ces films en fonction des paramètres du procédé et des niveaux de contrainte. Enfin, nous mettons en évidence les effets du procédé de condensation sur la qualité cristalline du film SiGe aux interfaces avec les oxydes grâce à l'effet de canalisation d'une technique de rétrodiffusion d'ions à moyenne énergie (MEIS : Medium Energy Ion Scattering)