Initial oxidation of zirconium: oxide-film growth kinetics and mechanisms

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1 Max-Planck-Institut für Intelligente Systeme (ehemals Max-Planck-Institut für Metallforschung) Stuttgart Initial oxidation of zirconium: oxide-film growth kinetics and mechanisms Georgijs Bakradze Dissertation an der Universität Stuttgart Bericht Nr. 238 November 2011

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3 Initial oxidation of zirconium: oxide-film growth kinetics and mechanisms von der Fakultät Chemie der Universität Stuttgart zur Erlangung der Würde eines Doktors der Naturwissenschaften (Dr. rer. nat.) genehmigte Abhandlung vorgelegt von Georgijs Bakradze aus Riga/Lettland Hauptberichter: Mitberichter: Prüfungsvorsitzender: Prof. Dr. Ir. E. J. Mittemeijer Prof. Dr. J. Bill Prof. Dr. E. Roduner Tag der Einreichung: Tag der mündlichen Prüfung: MAX-PLANCK-INSTITUT FÜR INTELLIGENTE SYSTEME (ehemals MAX-PLANCK-INSTITUT FÜR METALLFORSCHUNG) INSTITUT FÜR MATERIALWISSENSCHAFT DER UNIVERSITÄT STUTTGART Stuttgart 2011

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5 Wovon man nicht sprechen kann, darüber muss man schweigen. 1 L. J. J. Wittgenstein ( ) 1 (Germ.) What we cannot speak about we must pass over in silence.

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7 Contents Contents General introduction Initial oxidation of metals Focus of the thesis Zirconium and zirconium oxide (zirconia) Zirconium Zirconium oxide (zirconia) Short overview on zirconium oxidation studies Methods of characterization Angle-resolved X-ray photoelectron spectroscopy Real-time in-situ spectroscopic ellipsometry Scanning tunneling microscopy Time-of-flight secondary ion mass-spectrometry References The different initial oxidation kinetics of Zr(0001) and Zr(10-10) surfaces Introduction Experimental Data evaluation AR-XPS data RISE data Results and discussion Oxide-film constitution... 38

8 2.4.2 Oxide film thickness: comparison of AR-XPS and RISE analyses Oxide-film growth kinetics Conclusions References Valence-band and chemical-state analyses of Zr and O in thermally-grown thin zirconium-oxide films: an XPS study Introduction Experimental procedure and spectra evaluation Results and discussion The oxide-film valence band spectra The local chemical states of O and Zr in the oxide films Conclusions References An STM study of the initial oxidation of single-crystalline Zr surfaces Introduction Experimental Results and discussion Oxide-film microstructure at T = K Evolution of the oxide microstructure at 450 K Conclusions References Atomic transport mechanisms in thin oxide films grown on zirconium by thermal oxidation, as-derived from 18 O-tracer experiments...83

9 5.1 Introduction Experimental Specimen preparation Thermal oxidation In-situ deposition of an Al capping layer XPS, ToF-SIMS and HR-TEM analysis The oxide-film microstructure O-tracer depth distributions: identification of governing transport mechanisms Proposed oxidation mechanism Conclusions References Summary Zusammenfassung List of used abbreviations List of publications Acknowledgements...117

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11 Chapter 1 General introduction 1.1 Initial oxidation of metals The chemical interaction of a metal surface with oxygen gas is a typical example of a heterogeneous chemical reaction, which can be described by the following chemical equation: x Me(s) + y/2 O 2 (g) Me x O y (s). (1.1) The oxidation reaction typically has a large thermodynamic driving force, owing to a negative standard free-energy change, G. Consequently, most oxidation reactions run spontaneously in the forward direction under ambient conditions: i.e. they do not require any activation. As a result, a native oxide film inevitably forms on most metal surfaces under ambient conditions. However, the sign and magnitude of G only states the feasibility (i.e. the thermodynamic driving force) for the oxidation reaction. It does not bear any information on the rate of oxide formation, as governed by the reaction kinetics at the phase boundaries (i.e. the gas/oxide and oxide/metal interface), as well as by the transport rate of the reactants (metal, oxygen and electron species) through the initial oxide film [1]. The native oxide layer on a metal surface typically acts as a diffusion barrier between the reactants, thereby inhibiting further oxidation of the underlying metallic compound. In practice, this knowledge is utilized to enhance the corrosion resistance of metallic coating systems under operation at higher temperatures by the deliberate (pre-)formation of a protective (i.e. adherent and stable) α-al 2 O 3 or α-cr 2 O 3 overlayer [2-3]. The thin surficial oxide on a metal has a direct influence on the chemical and physical properties of metallic components (e.g. corrosion resistance, adhesion activity, thermal stability, catalytic properties, tribologic properties, wear resistance, electrical properties etc.), as applied in state-of-the art technologies, such as optical, adhesive and corrosion-resistance coating systems, functionalized surfaces, microelectronics, heterogeneous catalysis and sensor devices [4]. Clearly, most of the above-mentioned properties are determined by the microstructure (e.g. thickness, morphology, chemical constitution, crystallographic and defect structure) of the surficial oxide. In order to control the oxide-film microstructure and, thereby, the material's properties during fabrication and subsequent operation, a profound

12 10 Chapter 1 knowledge on the relationships between the oxide microstructure and the oxidation conditions (e.g. temperature, oxygen partial pressure, surface purity) is required. In addition, fundamental understanding of the transport phenomena in thin growing oxide layers is needed to improve e.g. the corrosion resistance of metal and alloy surfaces. Up to date, the growth kinetics, the microstructural evolution, as well as the mechanical and transport properties, of oxide layers grown on metals and alloys by thermal oxidation have been investigated mainly at elevated oxidation temperatures (say at T > 500 K) [5]. At such elevated temperatures, solid-state transport through the growing oxide film (under influence of chemical potential gradients) becomes pronounced and, consequently, relatively thick (typically up to several hundreds of micrometers) oxide scales develop. The microstructure of these oxide scales can be well-characterized by conventional analysis techniques, such as gravimetry, interferometry, light or electron microscopy and XRD analysis. Contrarily, our knowledge on the thermodynamics and kinetics of the oxidation process at low temperatures (of, say, below 500 K) is far from complete and still suffers from the lack of reliable quantitative experimental data. This can be mainly attributed to the fact that the oxide films developing on metallic surfaces at low temperatures are typically only very thin (thickness < 10 nm). Delicate and expensive ultra-high vacuum (UHV) systems for surface preparation, controlled oxidation and surface analysis are mandatory to process and characterize the thin oxide-film microstructure in-situ. Additional experimental difficulties arise due to elaborate in-vacuo surface-preparation procedures. Indeed, first studies on the controlled oxidation of metals at low temperatures (under well-defined conditions) were reported in parallel with the fast development of UHV technologies, starting from the late 1960s. Already long before it was recognized that the electrochemical and mechanical properties of metal and alloy surfaces depend, to a great extent, on the crystallographic, defect and electronic structure of the native oxide. Wagner was among the first to develop a theoretical description for the hightemperature oxidation of metals and alloys on the basis of the thermally-activated diffusion of reactant species through the developing oxide scale under influence of the acting chemical potential gradients [6]. His theoretical framework in turn promoted the postulation of theories on the low-temperature oxidation behaviour of metal surfaces (i.e. in the absence of thermally activated diffusion), as governed by the surface field setup by chemisorbed oxygen species on the oxidizing surface (cf. Mott-Cabrera [7], Fromhold [2] and Fromm [3]). Despite these important advances in the field of low-temperature oxidation, the processes that occur and the changes that take place at the metal surface, in the developing oxide film, and at the

13 General introduction 11 metal/oxide and oxide/gas interfaces during the initial and subsequent stages of oxide-film growth are still only partially understood. Clearly, the initial oxidation of a bare metal surface is a complex physicochemical process, which involves a series of competing and overlapping processes (see Fig. 1.1), such as [2]: (i) oxygen molecule impingement on the metal surface and (ii) subsequent physisorption (i.e. O 2(g) O 2(phys) and/or O 2(g) 2O (phys) ), (iii) chemisorption of oxygen species (i.e. O (phys) + e O (chem) and/or O 2(phys) + e O 2 (chem) ); (iv) place exchange between atoms in the metal subsurface and chemisorbed oxygen species, incorporation of anions and cations into the growing oxide film and (v) oxide nucleation and continued oxide-film growth. Fig Schematic illustration of some competing physical and chemical processes for the initial stages of dry thermal oxidation of a bare metallic substrate as described by Eq. 1.1: (i) impingement and (ii) physisorption of oxygen molecules on the surface, (iii) dissociative oxygen chemisorption, (iv) incorporation of anions and cations into the growing oxide film. Note: only transport processes in the direction parallel to the surface normal are shown. Oxide nucleation, lateral growth and coalescence of oxide nuclei, followed by continued oxide-film growth (thickening), generally involve both volume and short-circuit transport of the reactants (metal cations, oxygen anions and their vacancies, as well as electrons and their holes) through the developing oxide film. The aforementioned theoretical descriptions of the oxidation process typically assume that the initial oxide layer covers the metal surface uniformly and that further oxide-film growth proceeds in a uniform way (i.e. by a layer-by-layer thickening). Evidently, these oversimplified assumptions are serious deficiencies in the theoretical treatments [2-3, 7] of, particularly, the initial oxidation stages [8]. The situation becomes even more complicated by the fact that the coupled (charged) currents of ions/vacancies and electrons typically act in the presence of steep gradients in the defect concentrations, intrinsic stress-level and/or the electric-field across the developing oxide film. Finally, it is noted that the oxidation behavior not only depends on the developing

14 12 Chapter 1 oxide-film microstructure, but also on e.g. the crystallographic orientation of the parent metal substrate, the presence of any (segregated) impurities, and in some cases on the rate of dissolution of oxygen into the metal (especially for elevated oxidation temperatures and low oxygen partial pressures). 1.2 Focus of the thesis The present PhD work addresses the growth kinetics (Chapters 2), chemical composition (Chapters 2, 3), morphology (Chapter 4) and transport properties (Chapter 5) of zirconiumoxide films, as grown by the dry, thermal oxidation of single-crystalline Zr surfaces at low oxidation temperatures (for details see Section 1.4). To this end, oxide films with thicknesses in the range of 1-10 nm were grown on bare (i.e. without a native oxide) Zr(0001) and Zr(10 1 0) surfaces by controlled exposure to O 2 (g) in the temperature range of T = K at an oxygen partial pressure of po 2 = Pa in an especially-designed UHV system. This study reveals, for the first time, the effect of the metal substrate orientation on the low-temperature growth kinetics and microstructural evolution of the initial oxide overgrowths on Zr single-crystalline surfaces with basal and prism orientations. Furthermore, two-stage tracer oxidation experiments using 16 O and 18 O isotopes were performed to reveal the governing atomic transport mechanism(s) in the thin oxide films developing on both Zr single-crystalline surfaces at 450 K. 1.3 Zirconium and zirconium oxide (zirconia) Zirconium Zirconium (Z = 40, [Kr]4d 2 5s 2 ) belongs to the IV group of the periodic table and, together with titanium and hafnium, forms a small IVb subgroup of metals, which all to a large extent exhibit very similar physical and chemical properties. Metallic Zr is extracted from minerals, such as zircon (ZrSiO 4 ) and the natural form of zirconium oxide baddeleyite (ZrO 2 ). Typical natural impurities in zirconium are oxygen, iron and hafnium. Zirconium has a silver-grayish colour with a characteristic metallic blister. Under atmospheric pressure zirconium has two allotropic modifications: the hexagonal α- phase with a Mg-type crystal structure (c/a = 1.59, a = Å, stable at T < 1139 K) and the BCC cubic β-phase with a α-fe-type crystal structure (stable at T > 1139 K) [9]. The relatively low α-to-β transition temperature in combination with the high chemical reactivity

15 General introduction 13 of metallic Zr make it very difficult to prepare pure Zr single crystals. Pure Zr is ductile and can be readily processed by conventional metal forming techniques like forging, cold-rolling or drawing and welding under inert atmosphere [10]. However, the presence of oxygen, nitrogen, carbon and hydrogen strongly affects the mechanical and chemical properties of Zr [5]. Fig The Zr-O phase-diagram [11]. Note an extensive region of solid solutions of O in α-zr even at low temperatures. Zirconium and its alloys are stable in water, air and acids (except hydrofluoric and concentrated sulphuric acid) at room temperature. Due to its good corrosion resistance at low temperatures, Zr and its alloys are increasingly being used in applications for chemical processing equipment, oceanic instruments and marine hardware [12-13]. Some Zr alloys are used in biomedical applications (e.g. for hip joint implants) due to their wear resistance and compatibility with biological tissues [4]. Zr has a high affinity for oxygen ( G = 1037 kj/mol at standard conditions [14]) and actively absorbs gases (like O 2 and H 2 ) [5] at elevated temperatures T > 500 K. In addition, the O solubility in α-zirconium is as high as 30 at.% even at moderate temperatures (see Fig. 1.2). Consequently, Zr is very reactive to the residual gases in vacuum (especially to CO

16 14 Chapter 1 which dissolves readily). These properties insure the use of Zr-alloys as a getter material in a new generation of getter pumps for UHV applications [15]. For many decades, zirconium finds his primary use in the nuclear industry 1 for inreactor components, especially in the cladding of the fuel rods, due to zirconium's low neutron scattering cross-section and passivating oxidation behavior at low operating temperatures < K [12-13]. The integrity of the passivating oxide film on the cladding elements is crucial in order to assure the safe reactor operation. Hence thorough information is required concerning the structure, composition, kinetics and mechanism of the growth of oxide films on Zr and its alloys Zirconium oxide (zirconia) Zirconium-oxide consists in three crystalline forms: the monoclinic α-phase at low temperatures, the tetragonal β-phase above 1400 K and the cubic γ-phase above 2600 K [14]. Crystalline zirconium-oxide generally exists in the monoclinic phase at room temperature. However, the cubic phase can be stabilized at lower temperatures by the enhanced formation of vacancies in the anion sublattice, as induced by the addition of e.g. ZrN, CaO, Y 2 O 3, MgO [17]. The existence of a metastable, amorphous Zr-oxide phase has also been reported for the thermal oxidation of Zr substrates for short oxidation times at low temperatures (up to about 573 K) (see Ref. [18]) and can be rationalized on a thermodynamic basis [19]. In nanomaterials, the tetragonal ZrO 2 phase can be stabilized due to its lower interfacial energy (as compared to monoclinic phase) [20-21]. Zirconia finds diverse applications in jewellery, microelectronics, fuel cells and oxygen sensors [22-23]. Due to its high melting point (2953 K), chemical durability and high hardness, zirconium dioxide has long been used for refractory containers and as an abrasive medium. Zirconia-based materials have similar thermal expansion coefficients as some hightemperature-resistant metallic alloys and are therefore widely applied in thermal barrier coating systems for jet-engines [24]. In recent years, zirconia has also been considered as a promising candidate for tunnel barrier applications in next-generation metal-oxidesemiconductor field-effect transistor (MOS-FET) devices owing to its high dielectric constant (κ 25), its large conduction band offset with Si, as well as its predicted stability with respect to solid-state reactions with Si [25-26]. 1 Unlike other metal alloys, zirconium alloys applied in nuclear applications are rather pure (>97-98% Zr) and can almost be considered as single-component systems [4]. However, a strong influence of foreign elements on the diffusion in Zr has been reported in [16].

17 General introduction Short overview on zirconium oxidation studies Up to date, the oxidation behaviour of zirconium and its alloys has been extensively investigated under various oxidation conditions (e.g. oxidation temperature and time, partial oxygen pressure, oxidizing atmospheres; cf. Refs. [27-32]) and by using a broad range of (surface-)analytical techniques. Particularly, the effect of the oxidation temperature and partial oxygen pressure on the oxide-film growth kinetics and the developing oxide-film microstructure has been studied thoroughly, but unfortunately unsystematically. Many of the earlier oxidation studies on Zr pertain to the high-temperature oxidation regime and often suffer from the fact that only a single analytical technique has been applied to characterize the developing oxide microstructure. Several theoretical and experimental studies have focused on the very initial stages of oxygen interaction with the bare Zr(0001) surface (i.e. for oxygen exposures < 50 L) [33-40]. It was found that, at 90 K, 293 K and 473 K the adsorbed oxygen atoms preferentially occupy octahedral subsurface sites in the Zr(0001) substrate for low oxygen coverages (< 0.5 ML) [33, 40]. The oxygen atoms may penetrate into the subsurface octahedral sites, thereby leaving free sites on the top surface layer and thus enabling further adsorption and incorporation of oxygen. Initial exposure of the bare Zr(0001) surface up to about 1 Langmuir (see footnote 1 ) of O 2 (g) and subsequent flash-annealing to about 473 K yields a well-ordered (2 2) O-adsorbate overlayer structure by low-energy electron diffraction (LEED) [38, 41-43]. On the prism plane the stable surface phase is Zr (10 10) -O(2 4) [44]. The oxidation kinetics of the Zr(0001) surface for more prolonged oxygen exposures have been studied by Auger electron spectroscopy [40]. It was found that oxide-layer growth proceeds by a layer-by-layer growth mechanism at 90 K, whereas at T 293 K oxidation occurs by initial formation of oxide islands, which predominantly grow inwardly into the metal substrate. Oxidation of the Zr(0001) surface at room temperature results in the formation of a ultra-thin disordered (amorphous), overall non-stoichiometric oxide film [38, 41]. Investigations by X-ray Photoelectron Spectroscopy (XPS) have revealed that the oxide film is constituted of a non-stoichiometric suboxidic layer (ZrO x with x < 2, which contains several Zr δ+ valence states with δ < 4) at the interface with the parent Zr metal and a (near- )stoichiometric ZrO 2 oxide adjacent to the surface [4, 30-31, 45]. 1 1 Langmuir = Torr s.

18 16 Chapter 1 Only a few studies have been reported on the oxidation of single-crystalline Zr surfaces other than Zr(0001). Noteworthy, the oxidation behaviour of the Zr (10 10) prism plane is of particular industrial interest, because the cold-rolled Zr samples are textured with {10 1 0}-planes parallel to the surface [34, 44]. Already in the early 50s, it was found that Zr (10 11) and Zr (1120) surfaces have a lower oxidation rates than Zr(0001) and Zr (10 10) surfaces [34]. For the oxidation of Zr at 773 K in steam [46], the oxidation rate of Zr surfaces of different crystallographic orientation increased in the order: (10 12) < (1120) < (10 10) < (10 11) <(0001). However, these findings are in contradiction with the results reported in Ref [47], stating that the oxidation rate of polycrystalline Zr reached its minimum when the c-axis was normal to the surface plane, whereas the maximum oxidation rate is attained when the c- axis is inclined to the surface plane of the sample by 20. The kinetics of oxidation of Zr(0001) and Zr (10 10) (1 4) surfaces were investigated by the Norton group [38, 43-44] in the K range, however the oxygen exposures in all experiments were very low and the oxidation kinetics of both substrates has not been done. Current scientific and technological interest concerns primarily the successive stages of the oxidation process associated with: (i) formation of a closed oxide film on the Zr metal surface and (ii) subsequent thickening of the thin (thickness < 100 nm) oxide-film by transport of reactant species (i.e. cations, anions and their vacancies, as well as electrons and electron holes). Surprisingly, the initial oxidation of Zr at intermediate temperatures (i.e. T = K), where oxide films can be grown with controllable thicknesses in the nanometer range, have been largely unaddressed up to date [34, 44]. Also comprehensive knowledge on the developing oxide-film microstructure and its atomic transport mechanisms during the oxidation process at intermediate temperatures is still lacking. This is also evidenced by the contradictory statements in the scientific literature on the oxidation mechanisms of Zr at low and intermediate temperatures. For example, it was postulated in Ref. [48] that the growth of an amorphous Zr-oxide film on Zr at low temperatures proceeds by the coupled currents of cations and electrons across the oxide-film under influence of the surface-charge field (as setup by chemisorbed oxygen species at the oxidizing surface). Also in Ref. [31], the transport of Zr 4+ ions to the film surface was proposed as the rate-limiting step for the oxidation process. However, according to Ref. [49], the rate-limiting step in the oxidation process is either the O 2 dissociation rate or the O transport rate through the growing ZrO 2 film. There are also contradictions regarding the rate-controlling step for oxygen dissolution into the Zr metal at elevated temperatures. As evidenced from O 18 -isotope tracer oxidation

19 General introduction 17 experiments, the inward migration of oxygen along oxide grain boundaries (GBs) plays a dominant role in the growth of thick, polycrystalline oxide scales on Zr [50-52] and its alloys [53] at elevated temperatures. According to Ref. [54], the O dissolution rate is governed by the transport rate of O through the interfacial suboxide layer and not by the rate of oxygen dissolution in the metal at the suboxide/metal interface (as postulated in Ref. [33]). Upon oxidation of zirconium alloys at higher oxidation temperatures (T > 600 K) oxygen shortcircuit transport along the crystallite GBs is commonly considered as the predominating transport mechanism [29, 51]. Recently, an investigation on the initial oxidation of weakly textured, polycrystalline Zr surfaces in the temperature range of T = K was carried out at Max Planck Institute for Metals Research (Stuttgart). In that study the native oxide on the Zr surface was removed by Ar + sputter-cleaning (SC) under UHV conditions [55-56], but without performing a final in-vacuo annealing step prior to oxidation (to restore the distorted crystallography at the ion-bombarded surface). Consequently, the surfaces were not in the crystallographically well-defined state and hence the effect of the substrate orientation on the oxidation process could not be established. 1 The present study, for the first time, presents a direct comparison of the initial oxidation of single-crystalline Zr surfaces with basal and prism orientations (i.e. Zr(0001) and Zr(10 1 0), respectively) performed under the same, well-controlled experimental conditions in the temperature range of K and at po 2 = Pa. To this end, well-defined, single-crystalline Zr surfaces were prepared by an elaborate cyclic treatment of alternating Ar + SC and in-vacuo annealing steps under UHV conditions. The relationships between the oxidation kinetics, the developing microstructure and the crystallographic orientation of the parent metal substrate were established by application of various (surface-)analytical techniques (see Section 1.5). Furthermore, two-stage tracer oxidation experiments using 16 O and 18 O isotopes were successfully employed to reveal the atomic transport mechanisms in the very thin (thickness < 10 nm) oxide films, as grown on the single-crystalline Zr surfaces by thermal oxidation at 450 K and po 2 = Pa. 1.5 Methods of characterization Thorough characterization of the growth kinetics and microstructural evolution of thin (< 10 nm) oxide overgrowths on bare Zr substrates requires a complex experimental approach by 1 To see the effect of SC on the surface morphology see Fig.4.1.

20 18 Chapter 1 various in-situ, preferably non-destructive, surface-sensitive analytical techniques. In the present thesis, a combined experimental approach by in-situ angle-resolved X-ray photoelectron spectroscopy (AR-XPS), real-time in-situ spectroscopic ellipsometry (RISE), in-situ scanning tunnelling microscopy (STM), ex-situ High-resolution Transmission Electron Microscopy (HR-TEM) and ex-situ time-of-flight secondary mass-spectrometry (ToF-SIMS) has been applied to study the microstructural evolution and growth kinetics, as well as transport mechanism, during growth of thin oxide overgrowth on Zr single crystals Angle-resolved X-ray photoelectron spectroscopy XPS belongs to the vast family of electron spectroscopic techniques and is widely used for chemical characterization of near-surface regions in solid compounds. The working principle of XPS is based on the measurement of the energy spectrum of electrons emitted due to the outer photoeffect, i.e. due to the high-energy photon-induced photoemission of electrons from characteristic energy levels of the atoms in the sample [57]. Fig Schematic view of the AR-XPS principle. The sample is exposed to a flux of monochromatic photon radiation with energy hν. The kinetic energies, E k, of the ejected electrons are recorded at different detection angles θ, thus obtaining depth-resolved information of the investigated chemical species. Adopted from Refs. [57-58]. In an AR-XPS set-up the kinetic energy of the emitted photoelectrons, E k, can be recorded at various detection angles θ (with respect to the sample surface normal), as illustrated in Fig The sensitivity of XPS to the surface composition of the sample originates from the fact that emitted photoelectrons with kinetic energies below 500 ev are

21 General introduction 19 easily inelastically scattered in the solid. For example, for the electrons with energies of ev in inorganic solids, the inelastic means free paths (IMFP) is generally less than 10 Å. Only the photoelectrons, which are emitted in the near-surface region of the solid, have a finite probability to escape from the solid into vacuum and reach the detector without kinetic energy (KE) loss [59]. The measured KE of these unscattered and elastically scattered electrons can be easily be converted into the respective binding energies (BE), E b, using the following equation: E b = hν E k Φ, (1.2) where Φ denotes the work function of the spectrometer (for conductive samples it equals the work function of the sample). If the effects of elastic scattering of the traversing photoelectrons are neglected, 95% of the unscattered and elastically scattered photoelectrons (for θ = 0) will originate from depths up to 3 times the IMFP below the sample surface. The effect of elastic scattering can be accounted for by using the so-called effective attenuation length (EAL, symbol λ eff ) instead of the IMFP. The information depth then varies with the photoelectron detection angle according to: 3λ eff cos θ. Angle-resolved XPS measurements thus principally allow the investigation of the depth distribution of various chemical species in very thin (thickness < 6 nm) oxide films. In Chapters 2 and 3 of this thesis, AR-XPS analysis of the bare and oxidized metal substrates has been applied to determine the thickness, chemical composition and constitution of the oxide overgrowths on Zr. One of the most challenging problems in the quantitative processing of the recorded XPS data is the deconvolution of a measured spectrum into its individual spectral contributions as originating from different the chemical states of the elements in the solid [59]. The accuracy of quantitative AR-XPS analysis strongly depends on the methods used to reconstruct and subtract the superimposed background intensity of inelastically scattered electrons to the measured spectrum [58]. For a measured binding energy region of a corelevel photoelectron line with a single chemical-state contribution, the inelastic background can be easily removed by subtracting some arbitrary background, such as a linear or a Shirley-type background [58]. However, if the recorded XPS core-level spectrum contains several (overlapping) peaks (as is typically the case for a recorded XPS spectrum of an oxidized metal, which consists of at least one metallic and one or more oxidic main peaks; see Fig. 2.3), each main peak provides its own background of inelastically scattered electrons and, consequently, the simplified methods for the background subtraction can cause significant errors in the determination of e.g. the oxide-film thickness and composition. In

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