ACTAUNIVERSITATISUPSALIENSISUPPSALA2006

Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 228

Electronic Propertiesof Metal Oxide Films Studiedby Core Level Spectroscopy

JAN HINNERK RICHTER

ISSN 1651-6214ISBN 91-554-6673-7urn:nbn:se:uu:diva-7180

To Elina

List of papers

This thesis is based on the following papers. Reprints were made with per-mission from the publishers.

I. Electronic structure of lithium-doped anatase TiO2 prepared in ultrahigh vacuum J. H. Richter, A. Henningsson, P. G. Karlsson, M. P. Andersson, P. Uvdal, H. Siegbahn, and A. Sandell Phys. Rev. B 71, 235418 (2005)

II. Phase separation and charge localization in UHV-lithiated ana-tase TiO2 nanoparticlesJ. H. Richter, A. Henningsson, B. Sanyal, P. G. Karlsson, M. P. Andersson, P. Uvdal, H. Siegbahn, O. Eriksson, and A. Sandell Phys. Rev. B 71, 235419 (2005)

III. Li insertion in sol-gel prepared Mn doped TiO2 studied by elec-tron spectroscopy under ultra-high vacuum conditionsJ. H. Richter, P. G. Karlsson, G. Westin, J. Blomquist, P. Uvdal, H. Siegbahn, and A. Sandell Submitted to J. Phys. Chem. B

IV. Threshold effects in the O 1s x-ray absorption spectrum of TiO2J. H. Richter, B. Sanyal, P. G. Karlsson, A. Henningsson, M. P. Andersson, P. Uvdal, H. Siegbahn, and A. Sandell In manuscript

V. Ultra-high vacuum metal organic chemical vapor deposition of ultrathin ZrO2 films on Si(100) and Si(111) studied by electron spectroscopy P. G. Karlsson, J. H. Richter, J. Blomquist, P. Uvdal, T. M. Grehk, and A. Sandell Submitted to Surf. Sci.

VI. Band alignment of ultrathin ZrO2 films on Si(100): Film-thickness-dependent band alignmentA. Sandell, P. G. Karlsson, J. H. Richter, J. Blomquist, P. Uvdal, and T. M. Grehk Appl. Phys. Lett. 88, 132905 (2006)

VII. Band alignment at the ZrO2/Si(100) interface studied by photo-electron and x-ray absorption spectroscopy J. H. Richter, P. G. Karlsson, B. Sanyal, J. Blomquist, P. Uvdal, and A. Sandell In manuscript

VIII. Combinatorial chemical vapour deposition of an ultrathin ZrO2-TiO2 film on Si(100)-(2x1) in ultra-high vacuum J. H. Richter, P. G. Karlsson, and A. Sandell In manuscript

IX. Electronic structure investigation of (Zn,Co)O room tempera-ture ferromagnets O. Karis, J. H. Richter, S. Valizadeh, A. Surpi, J. Hunter Dunn, M. Adell, P. Svedlindh, V. Stançiu, P. Warnicke, A. Sandell, A. Pers-son, P. Eckhold, L. Nyholm, G. Westin, and B. Sanyal In manuscript

Conference contributions

Modelling electrochemistry in ultra-high vacuum: Li insertion into a thin anatase TiO2 film (oral presentation) ECOSS 22, Sept. 2003, Prague

Phase separation in UHV lithiated anatase titaniumdioxide nanoparti-cles (oral presentation) ACSIN 8/ICTF 13, June 2005, Uppsala

Comments on my participation Scientific work is always based on teamwork. I had the main responsibil-

ity for experiments, data analysis and manuscripts in papers I, II, III, IV, VII and VIII. In papers V and VI I was involved in the experimental work and discussion of the results. In paper IX my contribution was limited to experi-mental work.

Contents

Introduction.....................................................................................................9

1 Fundamental concepts................................................................................111.1 Light and synchrotron radiation .........................................................111.2 Matter and electronic structure...........................................................12

1.2.1 Some information about the electronic structure of solids .........141.3 Photon-electron interaction ................................................................141.4 In situ and ex situ................................................................................151.5 Ultra-high vacuum (UHV) .................................................................16

2 Film preparation.........................................................................................172.1 In situ chemical vapour deposition (CVD).........................................17

2.1.1 Combinatorial CVD....................................................................182.2 In situ vapour exposure (lithium insertion) ........................................182.3 Ex situ preparation techniques............................................................19

2.3.1 Sol-gel preparation......................................................................192.3.2 Electrochemical preparation .......................................................202.3.3 Electrodeposition ........................................................................20

3 Characterisation techniques .......................................................................213.1 Core level electron spectroscopy........................................................213.2 Photoelectron spectroscopy (PES/XPS).............................................22

3.2.1 Determination of the work function............................................253.2.2 Chemical shifts ...........................................................................263.2.4 PES line shape ............................................................................28

3.3 Surface sensitivity and depth distribution ..........................................293.3.1 Determination of film thickness .................................................30

3.4 X-ray absorption spectroscopy (XAS) ...............................................313.4.1 Crystal field splitting ..................................................................323.4.2 XAS delineation process.............................................................34

3.5 X-ray magnetic circular dichroism (XMCD) .....................................363.5.1 The XMCD sum rules.................................................................37

3.6 Core hole decay processes..................................................................383.7 Resonant photoelectron spectroscopy (RPES) ...................................403.8 Theoretical approach: DFT ................................................................42

4 Summary of the results ..............................................................................444.1 Lithium insertion into TiO2 systems (Papers I-III).............................44

4.1.1 Common effects of lithium insertion ..........................................444.1.2 Paper I: Electronic Structure of ultra-high Vacuum Lithium inserted anatase TiO2 ...........................................................................464.1.3 Paper II: Phase Separation and Charge Localisation in UHV lithiated anatase TiO2 Nanoparticles....................................................494.1.4 Paper III: Li insertion in sol-gel prepared Mn doped TiO2 studied by electron spectroscopy under ultra-high vacuum conditions............50

4.2 Core hole effects in XAS....................................................................524.2.1 Paper IV: Threshold effects in the O 1s x-ray absorption spectrum of TiO2..................................................................................52

4.3 Investigation of UHV-CVD deposited ZrO2 on Si.............................544.3.1 Paper V: Ultra-high vacuum metal organic chemical vapor deposition of ultrathin ZrO2 films on Si(100) and Si(111) studied by electron spectroscopy......................................................................544.3.2 Paper VI: Growth of ultrathin ZrO2 films on Si(100): Film-thickness-dependent band alignment ...................................................564.3.3 Paper VII: Band alignment at the ZrO2/Si(100) interface studied by photoelectron and x-ray absorption spectroscopy ..........................584.3.4 Paper VIII: Combinatorial chemical vapour deposition of an ultrathin ZrO2-TiO2 film on Si(100)-(2x1) in ultra-high vacuum........60

4.4 Magnetism..........................................................................................624.4.1 Paper IX: Electronic structure investigation of (Zn,Co)O room temperature ferromagnets ....................................................................62

5 Populärvetenskaplig sammanfattning ........................................................645.1 Inledning.............................................................................................645.2 Resultaten av min forskning...............................................................65

6 Acknowledgements....................................................................................67

7 Bibliography ..............................................................................................68

9

Introduction

Metal oxides have many interesting properties that allow their use in various important applications. Examples for these applications are batteries, solar cells and photocatalytic devices. For example anatase TiO2 can accommo-date small ions, such as H+ or Li+, and displays electrochromic properties upon ion insertion. Due to its ability to effectively and reversibly store a substantial amount of lithium TiO2 is an interesting candidate for cathode material in rechargeable lithium ion batteries [1]. Owing to the accompanying colour change upon ion insertion it is also considered for electrochromic devices [2,3,4]. Nanoporous TiO2 with its tremendous internal surface area has been successfully used in solar cells [2,5] (dye-sensitised solar cell (DSSC)). Additionally TiO2 has been shown to have photocatalytic properties under ultra violet (UV) light. Recently some metal oxide systems have also been found to exhibit ferromagnetic properties upon doping [6,7] and are thus ex-amined for their potential application in spintronics (dilute magnetic semi-conductors (DMS)). A field of major current economic interest is the hunt for high k materials in order to replace SiO2 as the gate oxide in metal oxide semiconductor field effect transistors (MOSFET). With the rapid down-scaling of electrical components the SiO2 layer is simply becoming too thin (currently 12 Å [8]) to ensure decent device performance, as direct quantum mechanical tunnelling through the oxide gate reaches unacceptable levels. Some metal oxides, such as ZrO2, have a far higher dielectric constant than SiO2, which means that a thicker oxide layer, preventing tunnelling while retaining the same properties as a thinner SiO2 layer, could be used. Accord-ingly tremendous effort is put into the study of high k metal oxide systems [9,10,11,12]. Combining different materials usually leads to the formation of some form of interface layer a few nm thick with characteristics distinctly different from those of the bulk materials [13,14,15,16]. Since the metal oxide thickness is already in the low nm regime it is obvious that studies of the interface region are becoming increasingly important.

The properties of various metal oxide systems have been studied using core level electron spectroscopy, especially photoelectron spectroscopy and x-ray absorption spectroscopy. Core level spectroscopy is a family of power-ful experimental techniques that reveal element specific information of the sample. Among the properties studied are the electronic and geometric struc-ture, but even information about the chemical state or the magnetic proper-ties of specific elements can be gained.

10

In this dissertation metal oxide systems (mainly grown in situ by ultra-high vacuum chemical vapour deposition) are studied by core level electron spectroscopy under ultra-high vacuum conditions. The first three papers (I-III) treat effects of lithium insertion in different TiO2 systems. In Paper IV we investigate the effects the creation of an O 1s core hole has on the unoc-cupied DOS and that in turn by final state effects on the experimental spectra of TiO2. The following four papers focus on the growth of ZrO2 (V-VII) and combinatorial growth of ZrO2-TiO2 (VIII) on silicon. In Paper IX a dilute magnetic semiconductor material is studied.

11

1 Fundamental concepts

In this chapter the most basic concepts encountered in this dissertation are introduced in often highly simplified pictures and it is tried to explain these concepts to the general public in (hopefully) simple terms.

1.1 Light and synchrotron radiation As the nature of light was first studied two approaches to explain the proper-ties of light were developed. One was a particle approach that assumes that a beam of light consists of small particles (photons), the other a wave ap-proach, where light is thought to be an electromagnetic wave. It was found that the properties of light could only be explained by a combination of the two approaches, thus light is now described by a particle-wave dualism, where it is both wave and particle. The energy (E) of one individual photon is given by E=h . Here h is the Planck constant and is the frequency of the wave.

In our experiments we use light generated by the MAX-lab synchrotron radiation facility in Lund. In a synchrotron facility electrons are accelerated to near light speed and forced to travel in a ring (The MAX II ring has 90 meter circumference [17]). Electrons bent from their path emit a continuum of light along their propagation direction. This light is extremely focused and highly polarised, with the polarisation vector lying in the ring plane. With special experimental conditions a circular polarisation of the light can be obtained.

12

Experimental end stations are usually equipped with a monochromator that allows us to choose light with well-defined photon energies over a broad range of energies. End stations are usually designed for specific types of experiments. Most of our experiments were performed at beamline D1011 [18]. Only combinatorial CVD was performed at beamline I311 [19]. At I311 the size of the light spot is very small, which allows us to better study the varying composition of the sample surface found in this experiment.

Figure 1. Beamline D1011 at MAX-lab.

1.2 Matter and electronic structure In a highly simplified picture atoms consist of a core of protons and neutrons and electrons orbiting the core. Each electron is characterised by a set of four quantum numbers, which can be thought of as describing the electron’s dis-tance from the core, the form of its orbital and its spin. No two electrons can have the same combination of quantum numbers, i.e. all electrons surround-ing the same atom differ in some way.

The four quantum numbers are n, l, m and s. The principle quantum num-ber n describes the distance the electron has from the core. It also divides the electronic structure into shells, where each shell can contain 2n2 electrons. For historic reasons these shells are named K (n=1), L (n=2), M (n=3), etc. A change in this number translates into the greatest change of energy for the electron. n can have a value of any positive integer, the lower the closer to the core.

The subsidiary quantum number l characterizes the shape of the atomic orbital. It has a significant effect on the binding energy, yet much smaller than n. l can run from 0 to n-1.

The magnetic quantum number m represents the relative orientation in space the orbital is pointing in. Barring the presence of an external magnetic field this number does not affect the energy. m can run from –l to +l.

The spin quantum number s can be either +½ (spin up) or -½ (spin down) and it characterizes the intrinsic angular momentum of the electron [20]. Elec-

13

trons experience a magnetic field from the protons in the core. A combina-tion of the l and s quantum numbers, the so-called spin-orbit coupling, leads to a small splitting of energies for different spin states. This split is propor-tional to the number of proton to the power of 4 and thus highly dependent on the core charge.

Electronic orbitals are named based on the quantum numbers n, l and s, for example 2p1/2. The first part of the name is a number equal to the value of n and the second part is a letter according to l. The letters are for historic reasons: s (l=0), p (l=1), d (l=2) and f (l=3). The last part is a combination of the l and s quantum numbers and refers to the total orbital momentum of the electron.

Figure 2. Schematic of atomic orbital binding energies.

Figure 2 shows a schematic illustration of the binding energy of the dif-ferent atomic orbitals. Here every line represents two electrons (one spin up, the other spin down). These electrons are usually degenerate in energy, i.e. they have exactly the same energy. The binding energy increases down-wards, thus lower lying levels are more tightly bound. Here we can see how increasing n and l quantum numbers affect the binding energies of electrons. It is noteworthy that the filling of levels progresses in a more complicated fashion for the higher orbitals, as can be seen by the 4s orbital having a higher binding energy than the 3d level. This means that the 4s level will be filled before the M shell is completely filled.

An important distinction is made between valence and core electrons. All chemistry occurs for the least tightly bound electrons, which are called va-lence electrons. The deeper lying orbitals are called core orbitals. They do not participate in interaction with their neighbours and their energy is very well defined and characteristic for a given element. That means every ele-ment has a set of core level binding energies that are unique for this element. This is of major importance for core level electron spectroscopy as this en-ables us to identify elements by analysing the binding energies of their elec-trons.

14

1.2.1 Some information about the electronic structure of solids In solids the highest electronic states are better described by a band structure than by atomic like orbitals. The extensive overlap of the many molecular orbitals in a solid leads to many levels with infinitesimally small energy separation that can thus be considered as a continuum. The two bands we usually consider are the valence and the conduction bands. The valence band consists of the highest occupied molecular orbitals, i.e. the weakest bound electrons, whereas the conduction band consists of the lowest unoccupied molecular orbitals.

Apart from many other aspects metals, semiconductors and insulators dif-fer in their electronic structure. Metals have an overlap between the valence and the conduction band, i.e. there is virtually no energy difference between the highest occupied and the lowest unoccupied state. Semiconductors and insulators on the other hand have well separated valence and conduction bands, which leads to a band gap. That is to say they can only absorb pho-tons with an energy above a certain limit. In the case of insulators this band gap is very large. If two materials are brought into contact with each other it is of interest to examine the relative position of the bands to each other. If the bands are close together electrons can easily be transported between the two materials, which is desirable in for example dye sensitised solar cells. On the other hand a large offset between the bands will inhibit electron transport, which is used in insulators.

In order to be able to compare experimental results a common reference has to be found. In electron spectroscopy the Fermi level is often chosen as the zero of energy, which all other energies are presented relative to. The Fermi level is defined as the level that at 0 Kelvin temperature has the prob-ability ½ of being occupied [21]. This is located in the middle between the valence and the conduction band. For a metal with its overlapping conduc-tion and valence band this basically coincides with the valence band edge. In a pure semiconductor the Fermi level is positioned in the middle of the band gap.

1.3 Photon-electron interaction If material is illuminated with x-ray light there are two main possibilities for photon-electron interaction.

An electron can be excited into a higher lying unoccupied valence orbital upon absorption of a photon. This can occur if the photon energy exactly matches the energy difference between the occupied and the unoccupied level and the transition of an electron between those two levels is allowed. If the dipole selection rules apply, which is the case for most free atoms and many solids, a transition must change the principal quantum number n by at

15

least 1 and the subsidiary quantum number l by exactly 1. Note that a negative change of n indicates a decay process into an unoccupied orbital rather than an excitation.

A less restrictive possibility of photon-electron interaction is photoioniza-tion, which was described by Einstein in 1905 [22]. Here the photon energy is sufficient to ionise the atom, i.e. the electron is completely removed from the atom and can then leave the material. In order for this to occur the photon energy must be larger than the electron’s binding energy. The number of emitted electrons is proportional to the intensity (the number of photons) of the incoming radiation, whereas the electron’s kinetic energy depends on the energy of the individual light quanta (the photon energy).

An important property that has to be considered in photoelectron spec-troscopy is the differential cross-section for photoexcitation. This quantity describes how probable it is for an electron to interact with an incoming photon, absorb its energy and enter an excited state. The differential cross-section is dependent on a number of factors [23], such as the experimental conditions, the type of electronic orbital, the energy of the photon in ques-tion and the transition the electron is to make.

1.4 In situ and ex situIn this dissertation we study samples prepared by many different methods. An important distinction exists between in situ and ex situ preparation of samples. In situ and ex situ mean in place and out of place respectively. Insitu thus means that all preparations and measurements are performed in the same place, i.e. the sample is never subjected to a different environment. For our studies on in situ grown metal oxide films this means that we need to modify the experimental end station according to our needs. For example the precursor liquid(s) have to be connected to the vacuum system and a dosage system installed. In an ex situ process the sample is prepared out of place and has to be transported by some means to where the measurement is per-formed. Thus it can be exposed to different environments, such as normal atmosphere or a sealed Argon containing module. In situ experimentation often simplifies the interpretation of experimental results, as all findings are a consequence of the experimental set up.

16

1.5 Ultra-high vacuum (UHV) We perform our experiments in an ultra-high vacuum environment. Vacuum is what remains, when everything is removed from inside a sealed volume (i.e. basically nothing). Ultra-high vacuum is like normal vacuum, just much more so, which means it contains more nothing and less everything else.

Since our experiments are electron detection based, we need vacuum so that the electrons can reach the analyser without being influenced, scattered or absorbed outside the sample. To perform electron spectroscopy experi-ments a base pressure of about 10-4 mbar is required.

Additionally our experiments are extremely surface sensitive, so the pu-rity of the surface is very important to us. A clean surface will become gradually contaminated by adsorption of residual gases. The rate of this con-tamination can be as high as 1 monolayer per second at a pressure of 10-6

mbar. This is completely unacceptable for our experiments therefore we need to prepare the experimental station so that the vacuum is as good as possible. To this end drastic measures are taken. The entire experimental station is baked for more than 16 hours at a temperature of at least 120°C. Successful baking leads to a base pressure below 10-10 mbar.

17

2 Film preparation

This chapter describes the preparation techniques used to produce the films examined in this study. The main preparation technique employed through-out this dissertation is in situ ultra-high vacuum metal organic chemical va-pour deposition. Additionally samples prepared by various ex situ techniqueshave been studied either as a comparison to the UHV model system or in order to characterise novel materials.

2.1 In situ chemical vapour deposition (CVD) CVD is a process that employs the chemical reaction of one or more precur-sor in gas phase on the sample surface to deposit a solid film. The chemical reaction is usually thermally activated with rather high activation tempera-tures (>900°C). This activation temperature can be lowered considerably by the use of metal organic precursors [24]. CVD offers attractive properties, such as uniform thickness over large areas and conformal step coverage of uneven surfaces [25,26].

To deposit our metal oxide films we perform single source precursor UHV-CVD. Single source meaning that the precursor molecule contains both the metal and the oxygen needed to form the oxide, thus no additional substances (e.g. oxygen gas) are needed for the deposition process. This improves the quality and purity of the film and facilitates the control of growth conditions. Deposition is achieved by setting the sample to the de-sired temperature and directing the precursor to the sample via a stainless steel tube positioned close (a few millimetres distance) to the sample sur-face. The precursor partial pressure is controlled by a leak valve. As precur-sors we have used titanium(IV) isopropoxide [Ti(OCH(CH3)2)4 (TTIP)] for TiO2 and zirconium(IV) tert-butoxide [Zr(OC(CH3)3)4 (ZTB)] for ZrO2.

18

2.1.1 Combinatorial CVD In paper VIII combinatorial CVD is used to grow a film of graded stoichiometry. To this end both precursors were directed to the sample si-multaneously via two parallel stainless steel tubes. Each stainless steel tube is connected via a leak valve to an isolated vacuum system containing their respective precursor, but with no connection to the other precursor. Thus contamination of the vacuum systems could be excluded from the very be-ginning. A photograph of the UHV part of this dosage system is shown in Figure 3.

Figure 3. The double dosage device

Deposition is performed in two steps. As the preparatory step we begin by setting the sample to temperature but placing it well out of line of the deposition geometry to avoid sample contamination. The first leak valve is then opened until the desired partial pressure for this precursor is reached and has stabilized. It is then assumed that the partial pressure does not change significantly under the course of the experiment. The second leak valve is then opened until a satisfactory total pressure is achieved. In the second deposition step we then move the sample directly in front of the stainless steel tubes with a few millimetres distance. After the desired depo-sition time we quickly remove the sample from the deposition geometry.

2.2 In situ vapour exposure (lithium insertion) Lithium vapour is produced by electrically heating a mixture of lithium chromate and a reducing agent (SAES’ St101 (Zr %84-Al %16)). In addition to reducing lithium back to its metallic state the reducing agent also serves to remove reactive gases from the device [27]. The resulting metallic lithium vapour is then simply directed to the sample. In this process metallic lithium is inserted into the sample as opposed to electrochemical insertion, where lithium ions are used.

19

2.3 Ex situ preparation techniques

2.3.1 Sol-gel preparation In Paper III two sol-gel prepared films, a pure TiO2 film and an Mn modified Mn:TiO2 film, were studied.

Sol-gel preparation employs the transition of a colloidal liquid (the sol, containing solid precursor material) from liquid into solid (gel) phase. The resulting porous gel is then chemically purified and fired at high tempera-tures to form high purity oxide materials.

The sol-gel derived films were prepared with an all-alkoxide route using the commercial Ti(OPri)4 (Aldrich) precursor and a novel Mn-alkoxide pre-cursor, Mn19O12(moe)14(moeH)14 (moeH = 2-methoxyethanol), prepared by metathesis of anhydrous MnCl2 and 2Na or 2Kmoe in toluene-moeH solvent [28]. This novel, reactive Mn-precursor provides facile removal of all organic components in the gel formed by hydrolysis by the moisture in the air on spin-coating. The purely inorganic gel-films provide an advantage in the homogeneous doping and more controlled conversion of the gel to oxide on heat-treatment, since there are no organic groups left that cause uncontrolled local heating and reduction of ions in the gel and subsequent re-oxidation by air-oxygen. This difference has been shown for the corresponding Co:TiO2system [29] where gels with only minor amounts of organic groups showed phase impurities while films free of organic became phase pure anatase.

On heating the corresponding powders of the Ti and Ti-Mn-gel films to 600°C, water was removed endothermically up to 200°C yielding essentially an oxide film, which showed a rather sharp exothermic peak in the differen-tial scanning calorimeter at ca 490°C for the Mn doped films. This peak might be associated with crystallisation of the anatase phase, which was the sole phase at 600°C for these samples.

20

2.3.2 Electrochemical preparation In Paper I we study a sol-gel prepared TiO2 film that has been inserted elec-trochemically with lithium.

Nanoporous TiO2 films for electrochemical lithium insertion are prepared from colloidal dispersion of TiO2, which is doctor bladed onto a conducting glass substrate. Sintering at 450°C leads to a less than 3 m thick nanopor-ous TiO2 film consisting of nanoparticles with about 5-10 nm radius.

Insertion of lithium is performed in an Ar-filled glove box using a three-electrode set up with Ag/AgCl in saturated LiCl in propylene carbonate as the reference electrode (RE), a lithium ribbon as the counter electrode (CE) and the TiO2 film as the working electrode (WE). To avoid contamination the CE is isolated from the electrolyte. The electrolyte solution consists of lithiumtriflate (LiCF3SO3) in propylene carbonate.

2.3.3 Electrodeposition In Paper IX the magnetic properties of electrodeposited (Co, Zn)O are inves-tigated.

Electrodeposition is based on material deposition on a substrate from a solution under either controlled current or controlled voltage conditions. This is a versatile technique that has been successfully used to deposit a wide range of compounds including metals, semiconductors, mixed metal oxides and magnetic nanolayered structures.

The samples studied in this thesis were produced by means of cathodic electrodeposition of Zn1-xCoxO at an electrolyte temperature of 348-353 K. Oxygen was continuously supplied to the electrolyte solution as a source for the synthesis of oxide. The electrodeposition was carried out in a solution containing ZnCl2 and CoCl26H2O salts at concentrations of 5*10-3 mol/l and of 5.6*10-3 mol/l, respectively. A supporting electrolyte of KCl (0.11 mol/l) was used. The pH of the solution was 2.5. The cathode was a gold-coated silicon wafer and the anode a filament of platinum. The reference electrode was silver/silver chloride. Depositions were performed in two steps at –1.2 and –0.9 eV.

21

3 Characterisation techniques

3.1 Core level electron spectroscopy As mentioned earlier (section 1.2) the electronic structure of atoms can be divided into two subgroups: valence electrons and core level electrons.

Valence electrons form the chemical bonds between atoms. Thus they are shared between different atoms. In solids valence electronic levels are best described by bands.

Core level orbitals are localised on a specific atomic site with discrete binding energies. Core level binding energies are element specific, i.e. there are large differences in the core level position for different elements. Thus core level spectroscopy can reveal important element specific properties of the studied sample, such as oxidation state, magnetism and the electronic and geometric structure.

22

3.2 Photoelectron spectroscopy (PES/XPS)

Photoelectron spectroscopy is an experimental method that probes the occu-pied electronic levels. It is based on the photoelectric effect described by Einstein in 1905 [22]. The sample surface is illuminated with monochromatic light, i.e. light with a very well defined photon energy. The photons interact with the material and electrons are ejected. The kinetic energy of these elec-trons is analysed. A schematic of the PES process is shown in Figure 4. A core electron (black circle) absorbs a photon and is raised in energy above the local vacuum level.

Figure 4. Schematic of the PES process.

The kinetic energy just outside the sample surface is given by:

vacb

outkin EhE

Here h is the energy of the incoming photon, outkinE is the kinetic energy of

the electron just outside the sample surface and vacbE is the binding energy

referred to the vacuum level (the ionisation potential) [30]. The binding en-ergy relative to the Fermi level is given by:

sfb

vacb EE

vacbE and f

bE are the binding energies referred to the vacuum level and the Fermi level respectively and s is the sample work function (the energy required for an electron located at the Fermi level to actually reach the vac-uum).

23

For materials in electric contact the Fermi levels line up, i.e. if electric contact is made, the sample and spectrometer share the same Fermi level [31].The binding energy for solid-state measurements is thus usually referred to the Fermi level, as it is common for most materials.

Figure 5. Schematic energy level diagram for a sample in electric contact with the spectrometer.

The kinetic energy detected in the analyser differs somewhat from the ki-netic energy just outside the sample surface due to the different sample and analyser work function. This is illustrated in Figure 5. The binding energy relative to the Fermi level is thus related to the measured kinetic energy in the spectrometer by:

spspkin

fb -E-hE

With fbE being the binding energy relative to the Fermi level, h the photon

energy spkinE the detected electron kinetic energy in the spectrometer and

sp the work function of the spectrometer. Obviously the kinetic energy of the detected electrons linearly depends on the photon energy.

24

After an actual measurement we receive a spectrum with the number of electrons plotted against the binding energy. Figure 6 depicts an overview spectrum of a thin combinatorial ZrO2-TiO2 film on a Si substrate. The spec-trum contains the spectral contributions of zirconium, oxygen, titanium and silicon.

Figure 6. PES overview spectrum of a thin TiZrxOy film on Si.

25

3.2.1 Determination of the work function

As mentioned above the binding energies for conducting samples are meas-ured relative to the energy of the Fermi level EF. Often it is also important to study the work function of the sample. For example the energy levels of insulators are best referenced to the local vacuum level, which requires knowledge of the local work function. The work function can be determined by clever application of photoelectron spectroscopy. To this end a negative bias voltage is placed on the sample. This bias ensures that all electrons emitted from the sample are detected by the electron energy analyser. These detected electrons include zero kinetic energy electrons that lost energy due to inelastic scattering within the sample but had just enough energy to leave the sample. These give rise to the secondary electron cut-off feature. Meas-uring the total span of detected electrons up to the Fermi level and carefully calibrating the photon energy we can determine the sample work function. The work function is then given by the photon energy minus the total energy span of detected electrons.

26

3.2.2 Chemical shifts The determination of electron energies described above is only a simplified picture. In a more complete model all molecular orbitals of the entire sample have to be considered. Thus the electron binding energy obtained in PES is the difference between the total final state energy (after photoemission) and the total initial state energy (prior to photoemission) [32].

totinitial

totfinalb EEE

A more detailed chemical understanding can be gained via PES. As men-tioned before the core levels are localized and do not participate in chemical reactions. This leads in principle to each element having a characteristic series of well defined core level binding energies. However core level bind-ing energies are not completely independent of their chemical surrounding. A change in environment can induce small shifts in the total energies of the initial and final states in the order of some eV. This leads to an observable shift in core level spectra, which is called the chemical shift. If two different chemical species (A and B) of the same element are present in a sample their binding energies will differ according to the relative difference of their final and initial state energies. That is the binding energy for the two species will be given by:

(A)E(A)E(A)E totalinitial

totalfinalb

(B)E(B)E(B)E totalinitial

totalfinalb

Thus their binding energy difference Eb will be given by:

(B))E(B)E(-(A))E(A)E(E totalinitial

totalfinal

totalinitial

totalfinalb

27

This aspect makes PES an element specific technique (for the core levels) with which not only the elements contributing to a spectrum, but also their respective chemical state can be identified. PES is therefore also known as electron spectroscopy for chemical analysis (ESCA) [33]. As an example the Ti 2p3/2 photoemission feature of a lithium modified LixTiO2 film is shown in Figure 7. Here we observe two peaks originating from different Ti spe-cies, namely Ti 3+ and Ti4+. The total shift between the two species is about 1.6 eV.

Figure 7. PES spectrum of the Ti 2p3/2 core level.

28

3.2.4 PES line shape As mentioned above electronic energies are very well defined. Thus we would expect spectral features from PES to be very thin lines. The peaks we observe do however have a certain width. A number of effects influence the line width. There are experimental effects, such as the resolution of the ana-lyser or the energy width of the monochromatic light (Gaussian broadening). Also physical effects can lead to a broadening of the line. These include the temperature (thermal broadening - Gaussian) and the lifetime of the excited state (lifetime broadening, due to the Heisenberg uncertainty principle - Lor-entzian).An example for line broadening is given here in the form of the Coster-Kronig effect. The Coster-Kronig effect describes an auger decay process in which the core hole is filled by an electron from a higher subshell of the same shell. This is only possible if the energy difference between the two subshells is sufficient to excite or remove another electron. The Coster-Kronig effect leads to a reduction of the lifetime of the excited state and thus to a broadening of the observed spectral, feature, which is a direct conse-quence of the Heisenberg uncertainty principle mentioned above.

In the case of TiO2 the spin-orbit split of about 5.7eV for the Ti 2p1/2 and Ti 2p3/2 levels allows for a transition from the Ti 2p3/2 into the Ti 2p1/2 level upon simultaneous emission of a valence electron. The broadening effect can be observed in Figure 8, where the Ti 2p1/2 feature at 465eV exhibits a broader peak than the Ti 2p3/2 feature at 459.3eV.

Figure 8. PES spectrum of the Ti 2p core level.

29

3.3 Surface sensitivity and depth distribution Electron spectroscopy techniques are very surface sensitive. In order to be detected by our analyser electrons have to be emitted from the sample. An electron that absorbs a photon will gain a certain amount of kinetic energy. On the way from its atomic position out of the sample surface it can be either scattered elastically, inelastically or not at all. Electrons that are not scattered or scattered elastically will be detected at their original kinetic energy, giv-ing rise to the characteristic peak. Electrons that are scattered inelastically lose some of their kinetic energy and with it most meaningful information. These electrons contribute to a continuous background. Accordingly we can only gain information from depths that electrons can escape from without being inelastically scattered. For a given material this depth depends on the so-called inelastic mean free path, which describes the average distance an electron can travel in the material without being inelastically scattered. It is similar for most materials and is dependent on the kinetic energy of the elec-tron, which can be seen in the so-called universal curve (Figure 9) [34,35]. It is observed that the mean free path is very short (5 Å-35 Å), which is the rea-son for the surface sensitivity of electron spectroscopy. Due to the variation of the mean free path length with kinetic energy, as seen in Figure 9, a depth analysis of the sample is possible. By choosing different suitable photon energies, electrons originating from various depths of the sample can be collected. This can be very useful to determine the distribution of the ele-ments throughout the sample.

Figure 9. Sketch of inelastic mean free path as a function of kinetic energy.

30

3.3.1 Determination of film thickness In our experiments we often perform a growth series of a metal oxide on silicon and study the properties of that system with regard to the film thick-ness of the metal oxide overlayer. It is thus very important to be able to de-termine the thickness of said overlayer. Applying our knowledge about the inelastic mean free path mentioned in the previous section 3.3 allows us to estimate the film thickness for samples where the silicon signal of the buried bulk silicon is still detectable. Let the bulk silicon signal intensity of the clean silicon surface be I0. Deposition of material on top of this silicon sur-face will lead to an attenuation of signal intensity as more and more elec-trons are inelastically scattered the further they have to travel inside the solid. The measured bulk silicon intensity I for a given situation is then re-lated to the original bulk silicon intensity I0 via:

d

0

eII

Here I and I0 are the signal intensities of buried and surface bulk silicon respectively, d is the metal oxide film thickness and is the inelastic mean free path. The thickness of the oxide overlayer can thus be determined from:

0IIln*d

The intensities I and I0 are direct experimental properties, whereas is de-termined by the universal mean free path curve, the material properties and the experimental conditions.

31

3.4 X-ray absorption spectroscopy (XAS) X-ray absorption spectroscopy is an experimental technique that probes the unoccupied states of the sample. This is achieved by exciting an electron from a core orbital into an unoccupied valence orbital following the dipole selection rules (mentioned in section 1.3). The photon energy is scanned over an interval that includes an allowed transition from a core level to an unoccupied state. A core electron can thus be excited to an unoccupied level, leaving the atom within a highly excited unstable state. To illustrate the dif-ference to PES a schematic for the creation of the core hole is shown in Fig-ure 10.

Figure 10. Schematic of core hole creation.

The difference is that the excited electron occupies a high lying valence or-bital and is thus still located on the atomic site. The highly excited atom can then relax by either radiative decay emitting a photon or an Auger process emitting an electron (electrons emitted by these processes are called secon-dary electrons). These autoionization processes are schematically shown in Figure 16, but will be explained later (section 3.6). Either photon or electron detection based methods can be used to study the sample, but the focus here is, of course, set on electron detection. The number of secondary electrons is directly proportional to the number of absorbed x-rays and the x-ray absorp-tion spectrum is thus measured by recording the number of detected secon-dary electrons as a function of photon energy. X-ray absorption is highly photon energy dependent and an XAS peak will be observed if the photon energy exactly matches the energy difference between the core excited final state and the ground state (recall 3.2.2). This means that XAS actually probes the unoccupied states in the presence of a core hole at a specific atomic site. The implications of the presence of the core hole for XAS are discussed in Paper IV.

32

XAS spectra are extremely sensitive to the local symmetry of the absorb-ing species. Thus XAS spectroscopy is very useful for determining the crys-tal structure of ordered materials, as fingerprint spectra for nearly all crystal phases are known. Other applications of XAS include the determination of magnetic properties explained in section 3.5, the possibility to determine the orientation in space of surface adsorbates and an improved understanding of the electronic structure.

3.4.1 Crystal field splitting In this thesis XAS on transition metal oxides with partly filled d bands is performed. For d electrons a further splitting of otherwise degenerate energy levels occurs due to the orbitals’ orientation in space with regard to the local electron density. In a simplified picture the energy of an electron will in-crease if its wave function is directed towards a neighbouring atom and thus into a region in space with higher electron density. Since this is dependent on the crystal structure of the material this effect is called crystal field split.

Figure 11. Illustration of the geometry of 3d orbitals in octahedral symmetry.

As an illustration Figure 11 shows the form and orientation of the 3d orbi-tals for an atom in octahedral symmetry with the nearest neighbours shown as small circles positioned on each axis. Naturally the electron density sur-rounding the atom is highest in the direction of the nearest neighbours. The wavefunction for the eg states points towards nearest neighbour atoms, whereas it points diagonally for the t2g states.

33

This leads to an increase in energy for the 2zd and the 22 yx

d orbitals relative to the energy of the xyd , the xzd and the yzd orbitals. This energy splitting is shown schematically in Figure 12, where, again, each line repre-sents two electrons and the binding energy increases downwards (compare with section 1.2).

Figure 12. Schematic of the crystal field effect.

34

3.4.2 XAS delineation process Understanding the exact shape of XAS spectra is often a complicated matter, especially if there is more than one species present contributing to the ob-served XAS spectrum (comparable to the chemical shift in PES, see section 3.2.2). Nevertheless XAS spectra can at times be understood in a simple picture as a direct superposition from two or more individual contributions. For example in moderately lithium doped anatase TiO2 the Ti L-edge XAS spectrum can be envisioned to be comprised of contributions from an anatase like phase with low lithium content and a lithium titanate phase with high lithium content. We have identified fingerprint spectra for these two phases and have been able to interpret recorded XAS spectra as a superposition of these two. The relative contributions to the XAS spectra for all intermediate situations could be determined and a better understanding of the changes in geometric structure gained. If an unknown XAS spectrum is encountered and a certain atomic species is suspected in the sample, then this delineation process can also be employed to both identify the relative amount of that species in the sample and the origin of the remaining contribution (assuming only two contributing species are present in the sample).

Figure 13. Ti L-edge (2p 3d) XAS spectrum for a pure nanostructured anatase TiO2 film (A) and spectra obtained after deposition of 0.33 Li/Ti (B) and 0.45 Li/Ti (C). The intensity of spectrum A is chosen to illustrate its contribution to spectrum B. Spectrum B-A is the difference spectrum between B and A.

35

As an example the delineation procedure for lithium doped TiO2 is shown in Figure 13. Here the spectra A and C are characteristic of the anatase like and the lithium titanate phase respectively. Spectrum B is a spectrum of a situation with moderate lithium content. If this moderate situation can be viewed as a superposition of the two extreme cases, then the subtraction of the relative contribution of one of the phases should yield a spectrum identi-cal to the other phase. This is tried by subtracting an appropriate amount of spectrum A from spectrum B. This resulting spectrum should then only con-tain information from the lithium titanate phase and as seen by the similarity of the difference spectrum B-A to spectrum C (the titanate phase) this holds true. Thus it is possible to determine the relative contributions of the two phases if the spectral shape of the phases is known or identify the spectral shape of one contribution if the other is known and the assumption of a sim-ple superposition of two phases can be justified.

36

3.5 X-ray magnetic circular dichroism (XMCD) XMCD is an x-ray absorption technique for studying the magnetic properties of a material. Here circular polarised x-rays, carrying an angular momentum of for right helicity and – for left helicity, are used. Upon absorption the spin of an electron is not changed, i.e. an electron can only be excited into an empty state with equal spin. Typically the x-ray absorption from a spin orbit split p state into a partially occupied d band, split by exchange interaction, is examined. The XMCD mechanism will be explained here according to this transition and a highly simplified illustration is given in Figure 14 (left hand). The arrows correspond to the electron spin, the 2p and 3d levels are shown and the asymmetric partial occupation of the 3d states is indicated by the white area.

Figure 14. Schematic of the XMCD process (left hand) and the resulting XMCD difference spectrum of a magnetic material (Fe)(right hand).

In magnetic materials an exchange splitting of the d band removes the de-generacy for spin up and spin down states. Thus the majority of the occupied d band states will have a certain spin (majority spin), whereas the majority of the unoccupied states will be of opposite spin (minority spin). We thus excite electrons from the p3/2 and p1/2 states, which are split in energy due to spin orbit coupling, into the empty states of the d band, which are asymmetric with regard to spin orientation. With circular polarised light the probability of exciting an electron in a certain spin state depends on the helicity of the light. Using right helicity the probabilities of exciting a spin up or spin down electron are 62.5% and 37.5% for the p3/2 state and 25% and 75% for the p1/2state [36]. The opposite is true for left helicity. The empty d states then serve as a detector for the excited core electrons. The asymmetric distribution of empty spin states causes a spin selective excitation process, which allows us to gain information on the spin states of the occupied d states. An XMCD difference spectrum, where XAS spectra taken at opposite helicity have been subtracted from each other, is shown as the solid line in Figure 14 (right hand). We will now consider an experiment where the majority of unoccu-

37

pied d states are of spin down character. We will first observe two absorp-tion edges (the L3 and L2 edges) due to the spin orbit split p states. The in-tensities of these edges will be dependent on the available empty d states (i.e. dominated by spin down contributions in this case), their spin orientation and the helicity of the incident light. At the L3 edge the probability of excit-ing a spin down electron is higher for left helicity than for right helicity light, thus the observed feature will be higher in intensity for left hand polarized light. At the L2 edge however the probability for exciting a spin down elec-tron is lower for left helicity than for right helicity. Thus this feature will be lower in intensity for left helicity.

3.5.1 The XMCD sum rules Information on the magnetic properties (spin (ms) and orbital (ml) moment) can be extracted from the XMCD spectra with help of the magneto-optical sum rules [37,38] by the following relations [39]:

Zs T27

r2q3pCm

3r2qCml

6qp9

2mm

s

l

Here p and q are values taken directly from the integrated XMCD difference spectrum and are illustrated in Figure 14. r is a value from the integrated average spectrum, shown in Figure 15. The constant C is dependent on the number of empty states and the degree of polarization of the x-rays. <TZ> is the expectation value of the intra-atomic magnetic dipole operator. Note that the ratio ml/ms can be directly determined from the experimental spectra, which means that it can be determined to a higher degree of accuracy and without any theoretical assistance.

Figure 15. XMCD average spectrum for a magnetic material (Fe).

38

3.6 Core hole decay processes Excitation from the core levels creates a core hole and leaves the atom in an unstable, highly excited state. This state will rapidly decay by the transition of an electron from a higher level into the unoccupied core hole. To obey energy conservation the excess energy gained by the electron must be re-leased. This is possible by either emission of a photon (radiative decay) fol-lowing the dipole selection rules or by transferring the energy to another electron. Since our experiments are electron detection based the focus is here set on the decay path involving electrons shown schematically in Figure 16.

To a first approximation the electron filling the core hole gains energy equal to the energy difference between the two orbitals involved in the tran-sition. This energy is transferred to another electron, which gains enough energy to leave the sample and can then be detected. This is a two-electron process whose probability is governed by a Coulomb operator involving the electronic orbital overlap. Since the energy difference between orbitals for a given element is constant, the ejected electron will always have gained the same kinetic energy, depending only on the transition involved, i.e. the proc-ess is independent of the exciting photon energy. Thus this technique is ele-ment specific, as orbital energy differences are characteristic for each ele-ment.

Figure 16. Schematic of the possible decay pathways for core hole decay.

39

The Auger process sketched on the right hand side of Figure 16 occurs af-ter the creation of a core hole. A non-resonant excitation process, where the excited electron is completely removed from the atom, leaves the atom core ionised. In Auger decay the core hole is then filled from one of the occupied higher lying levels and another electron ejected. This leaves the atom in a doubly ionised state.

On the left hand side of Figure 16 the autoionization process is shown schematically. In a resonant excitation process a core electron is lifted to an unoccupied valence state. Decay can occur via two possible autoionization pathways both leaving the atom singly ionised. The first possible decay path is participator decay. Here the excited electron participates in the decay process. An electron from a higher lying level decays into the core hole, while another electron is emitted. Note that electrons are often indistinguish-able and it is not normally possible to identify which electron decays into the core hole. The other possible pathway is spectator decay. The same mecha-nism applies here except that the excited electron remains in its state and does not participate in the decay process. Filling of the core hole and emis-sion is performed by two other electrons. Note that the kinetic energy gained by the emitted electron is lower in this case as the final state is an excited state.

40

3.7 Resonant photoelectron spectroscopy (RPES) In RPES the excitation energy is scanned over an absorption edge, i.e. a pho-ton energy range that includes energies suitable for excitation from a core level to an unoccupied valence level. Unlike normal XAS the intensity of a photoemission peak is monitored. This technique is also known as constant initial state (CIS) spectroscopy.

This process can be used to enhance spectral features and elucidate their origin. For example in lithiated TiO2 we observe a band gap state at about 1eV binding energy, that is assumed to be mainly of Ti 3d character local-ised at Ti3+ sites. Using RPES we scan the photon energy over the Ti L edge (2p 3d). Neglecting the small partial population of the 3d level hybridised with O 2p states, the electronic configuration of the two species can in a simplified picture be envisioned as 2p63d0 for Ti4+ and 2p63d1 for Ti3+. In this picture lithium donates an electron considered to be of localised 3d character to Ti3+.

We will now discuss the different possible contribution to the RPES spec-trum:

First we will receive a contribution to the Ti 3d band gap state from direct photoemission of the Ti3+ 3d electron:

2p63d1 2p63d0

Here we end up with a final state configuration of 2p63d0 and an electron for detection.

In RPES we are interested in the resonant enhancement of this feature by excitation of an electron and subsequent decay into the same final state. Both Ti species can absorb photons at the Ti L edge and their photoexcitation process is as follows:

2p63d0 2p53d1 for the Ti4+ species and2p63d1 2p53d2 for the Ti3+ species

In this case the only decay that can occur is a 3d electron filling the 2p core hole. For autoionization to occur the energy from that transition must be transferred to an electron that can then leave the sample. In this case this is only possible for the 3d electrons and since one 3d electron has to fill the core hole, autoionization can only occur at the Ti3+ sites. The Ti4+ species lacks the second 3d electron that could absorb the energy. The Ti3+ species however can decay from its excited state via participator decay where the combined excitation and decay process can be viewed as a two-step process:

2p63d1 2p53d2 2p63d0

41

These are the same initial and final states as for photoionization, thus this process enhances the spectral Ti 3d feature.

The contribution from the direct photoemission process to the Ti 3d state can be considered constant within the chosen photon energy interval. The contribution from the Ti3+ autoionization process however resonates with the photon energy, i.e. the amount of absorbed photons and with it the number of detected electrons is highly dependent on the incident photon energy. Thus RPES can be seen as a measure for the probability of x-ray absorption at the Ti3+ site.

The resulting spectra of the example discussed above yield a two-dimensional map as shown in Figure 17. Here the resonant Ti3+ related states can be observed in the band gap between 1.5 eV and 3.5 eV binding energy.

Figure 17. An RPES spectrum of the band gap and valence region for lithiated TiO2.

42

3.8 Theoretical approach: DFT It is often interesting to compare experimental results with theoretical com-putations. Calculations of electronic properties of solids are very difficult to perform. Already three interacting particles are impossible to calculate ex-actly and in a solid the number of electrons is in the order of 1023. For this reason, all theoretical approaches have to make use of approximations.

A very powerful technique is density functional theory (DFT), which was introduced by Hohenberg and Kohn in 1964 [40] and further improved by Kohn and Sham in 1965 [41]. Instead of solving the many-electron wavefunc-tion they suggested to replace it with the electron density as the fundamental variable. In this approach the ground state is defined by that electron density, which minimises the total energy. All other ground state properties follow as a direct consequence from the electron density. Every electron is here envi-sioned to be moving in an average effective potential generated by all other electrons and cores. The ground state properties of an electron can then be calculated as a function of a single variable (the electron density) solving the energy functional. The total energy of a system is a function of the total ki-netic energy of non-interacting electrons, the Coulomb attraction (electron-ion interaction), the Coulomb repulsion (electron-electron interaction) and an exchange-correlation energy functional. Unfortunately the exact form of the exchange-correlation functional is not known. The most common approach is to use the local density approximation (LDA), where the exchange-correlation energy density of the real, spatially inhomogeneous system is approximated by that of a homogeneous electron gas with a density equal to the local density.

In Paper I we present calculations on the electronic structure of lithiated TiO2. DFT within the LDA framework falls short of yielding reasonable results for compounds containing transition-metals (partially filled d-band) and rare-earth metals (partially filled f-band). So a further complication (i.e. improvement) of the theory is in order. To correctly reproduce experimental results from these types of materials the LDA+U (local density approxima-tion + coulomb interaction parameter U) is used. In this method the electrons are separated into two subsystems. Localised d- and f- electrons, which are strongly correlated and demand the introduction of a Coulomb interaction parameter in the exchange-correlation functional and delocalised s- and p- electrons that can be described within the simpler LDA framework.

In Paper IV the implications of the presence of a core hole for XAS are studied. DFT helps to understand the energy alignment of XAS and PES spectra recorded. For core hole related studies the Z+1 approximation is often used. An additional proton increases the core charge by one in this approximation. It is assumed that the density of states of core ionised sys-tems is well represented by the system with the additional core charge.

43

44

4 Summary of the results

Here I present a summary of the publications presented in the second part of this dissertation. The first three papers (I-III) treat effects of lithium insertion in different TiO2 systems. In Paper IV we investigate the effects the creation of an O 1s core hole has on the unoccupied DOS and that in turn by final state effects on the experimental spectra of TiO2. The following four papers focus on the growth of ZrO2 on silicon. In Paper IX a dilute magnetic semi-conductor is studied.

4.1 Lithium insertion into TiO2 systems (Papers I-III) In these three papers lithium insertion into TiO2 systems is studied. Electro-chemical preparation processes are an important part of many applications and electrochemically prepared Li:TiO2 systems have been studied exten-sively. These systems always carry a substantial amount of surface contami-nation that cannot be avoided. Our idea was to develop a model system of the electrochemical lithium insertion process, which leads to similar inser-tion properties but without the accompanying contamination. To this end a pure UHV in situ process, where an ultra-thin in situ grown TiO2 film is exposed to lithium vapour, was chosen. This ultra-thin CVD grown anatase TiO2 film is extensively studied in papers I and II. We compare this UHV insertion process to results obtained from electrochemical insertion in Paper I. In Paper III this insertion process is tested on an ex situ sol-gel prepared Mn modified Mn:TiO2 film.

4.1.1 Common effects of lithium insertionHere we describe the results obtained from the lithium insertion studies that are common for all three TiO2 systems and both insertion methods. Under UHV conditions lithium is inserted into the TiO2 systems by exposure to atomic lithium vapour. The reaction taking place is:

TiO2+xLi LixTiO2

A very similar reaction occurs in the electrochemical experiment. Here lith-ium insertion is accomplished by an electrolyte solution containing lithium

45

ions. Electrons entering the film through the back contact counter the charge of the lithium ions diffusing into the film maintaining charge neutrality.

The reaction can be written as:

TiO2+xLi++xe- LixTiO2

In all cases lithium penetrates the film and is evenly distributed throughout the anatase part of the film. Charge is transferred from lithium to titanium, which is therefore reduced from its 4+ to its 3+ oxidation state. Using PES we have followed the relative contribution of Ti4+ and Ti3+ (and Ti2+) under lithium insertion. This behaviour is similar for both methods and shown for the UHV model system in Figure 18.

Figure 18. Evolution of the relative contribution of the Ti4+, Ti3+ and Ti2+ species to the Ti 2p core level photoemission spectra under lithium insertion.

In this figure the solid line represents the theoretical Ti3+ contribution assum-ing a homogeneous distribution of lithium throughout the anatase part of the film. For x values up to 0.35 (where x denotes the mole fraction of Li in TiO2) the behaviour of the Ti3+ contribution agrees well with the theoretical model, only for longer exposure times is a plateau reached. We thus con-clude that lithium can be inserted with a homogeneous distribution through-out the anatase part of the film up to an x value of about 0.4. Further interca-lation leads to lithium agglomeration on the sample surface with concomi-tant reduction of titanium to Ti2+ [42]. For electrochemical insertion the maximum amount of lithium insertion into anatase TiO2 varies with tech-nique, electrolyte and temperature, but does not typically exceed x=0.5 [43,44,45,46].

46

Upon lithium insertion we observe the population of a state of mainly Ti 3d character in the band gap. On a side note the population of this band gap state is the origin of the electrochromaticity in TiO2

[2,5].Two distinct LixTiO2 phases are identified. For small amounts of inserted

lithium a lithium poor phase (x<0.02), with a crystal structure very close to that of the pristine anatase film is observed [47]. Further lithiation leads to the formation of a lithium rich phase (x=0.5), denoted lithium titanate (Li0.5TiO2). The transformation from the anatase to the lithium titanate struc-ture is described by an orthorhombic distortion of the atomic positions [48].

4.1.2 Paper I: Electronic Structure of ultra-high Vacuum Lithium inserted anatase TiO2

In this paper we compare the electrochemical and UHV insertion processes. The focus is laid on the changes in electronic structure.

4.1.2.1 Comparison of UHV and electrochemical insertion A direct comparison of core level PES spectra obtained from the two inser-tion methods has been performed. The Ti 2p spectra (shown in Figure 19) are very similar in both cases, which shows that the use of the UHV ap-proach as a model system for electrochemical insertion is appropriate.

Figure 19. Ti 2p core level photoemission spectra for increasing amounts of lithium inserted by UHV evaporation (a) and electrochemistry (a).

As seen in Figure 20, the O 1s spectra for different amounts of inserted lithium in the two films clearly show the advantage of the UHV model sys-tem with respect to contamination. The UHV spectra shown in 20(a) display a single sharp O 1s peak, whereas contamination contributions in the form of a high binding energy shoulder are clearly visible for the electrochemical insertion system in 20(b). For high lithium insertion these even dominate the spectra for electrochemical insertion.

47

Figure 20. O 1s core level photoemission spectra for increasing amounts of lithium inserted by UHV evaporation (a) and electrochemistry (b).

Also shifts of the sample work function from UHV are found to agree rea-sonably well with measured cell voltages from electrochemical insertion further supporting the applicability of our model system.

4.1.2.2 Charge transfer The amount of charge transferred from Li to the emerging Ti 3d spectral feature has been estimated and is shown as the electron transfer ratio as a function of lithium content in the film in Figure 21.

Figure 21. Estimated charge transfer ratio from lithium to titanium as a function of x in LixTiO2. The as-measured values for the two highest points (open circles) have been corrected for the presence of the Ti2+ species.

48

To determine this ratio the integrated intensity of the Ti 3d feature and the Ti3+ contribution to the Ti 3p peak have been compared. The relative amount of electrons contributing to spectral features can be estimated by dividing the area of those features with the appropriate differential photoionization cross-section of the atomic species giving rise to that feature. Here it has to be taken into account that the electronic configuration of Ti3+ is 3p64s03d1. The cross-section for atomic titanium with two 3d electrons does not reflect the true differential cross-section of the Ti3+ 3d electrons. Good results have been achieved with an approximation using the differential cross-section of atomic Scandium. Scandium has an electronic configuration of 3p64s23d1,i.e. the same as Ti3+ except for two additional 4s electrons, whose influence on the 3d electrons is assumed negligible. The differential cross section of the Ti 3p core level is expected to be independent of the valence electron configuration. Applying this approximation, the electron transfer ratio from lithium to the Ti 3d band gap state is determined to 0.85, which is in good agreement with theoretical results [49]. The key point in this analysis is knowledge of titanium’s oxidation state. This is seen in Figure 21, where the as-measured transfer ratio for the last two points (empty circles) could be corrected for the presence of Ti2+ formed by lithium agglomeration at high x-values.

4.1.2.3 Valence band evolution and Ti 3d – O 2p hybridisation The evolution of the valence band under lithium insertion has been followed. No continuous rigid band shift behaviour for x 0.02 is observed. Instead we observe an initial jump of the binding energy of the valence band edge of about 0.2eV for x=0 0.02, after which it stays relatively constant, and pro-gressing changes in the shape of the valence band under the entire range of lithium insertion. This is a first indication of the localised character of the occupied Ti 3d states.

O 1s XAS spectra reveal two sharp peaks related to empty O 2p states strongly hybridised with Ti 3d states. These peaks are assigned to t2g (peak 1) and eg (peak 2) states according to the crystal field splitting model (see section 3.4.1). Upon lithium insertion both the intensity of peak 1 relative to peak 2 and the magnitude of the crystal field split decrease. Electron dona-tion mainly occurs into the Ti 3dyz orbital with t2g symmetry, which leads to a reduction of empty t2g states and thus a decrease of the associated peak. The reduced crystal field effect is explained by a decreased O 2p – Ti 3d hybridisation.

49

4.1.3 Paper II: Phase Separation and Charge Localisation in UHV lithiated anatase TiO2 Nanoparticles Paper II continues the study of the UHV system. We concentrate on the for-mation of the two different LixTiO2 phases. The electronic properties of the lithium titanate phase are investigated in detail.

4.1.3.1 Deconvolution of XAS spectra The shapes of the XAS spectra originating from the anatase and the lithium titanate phase have been identified. According to the procedure from section 3.4.2 (see Figure 13) it is demonstrated how XAS spectra of lithiated TiO2can be understood as a superposition of spectral contributions from the two different phases. In combination with PES measurements we can determine the lithium titanate stoichiometry to Li0.5TiO2, which is also predicted in other studies [48,50].

4.1.3.2 Electronic structure DFT calculations of the lithium titanate phaseDFT has been used to calculate the electronic structure of the lithium titanate phase. The unit cell is assumed to be a distorted anatase phase (space group Imma no. 74) with two different Ti sites with Ti-Li bond length of 2.82Å (Ti1) and 2.76Å (Ti2).

DFT calculations within the LDA+U framework show a correlation driven separation of Ti 3d states from the conduction band, leading to a band gap state. Since it is necessary to take electron-electron interaction into ac-count to reconcile the calculations with experimental results the Ti 3d state is considered to be of localised character. Even though the Ti-Li bond length of the two Ti sites are similar the calculations yield very different electronic structures at the two sites. The band gap state is predominantly associated with the Ti1 atom, whereas the states with highest binding energy in the valence band are dominated by contributions from the Ti2 atom. Thus the lithium titanate phase has two inequivalent Ti species with oxidation states +3 (Ti1) and 4+ (Ti2) as a direct consequence of the localised character of the Ti 3d state and the consequential electron correlation.

4.1.3.3. RPES and XAS The properties of the Ti 3d state of the lithium titanate phase have been stud-ied in detail comparing XAS spectra and RPES spectra of the band gap and valence band. According to the calculations Ti L edge XAS spectra should contain contributions from two different Ti sites with oxidation states 4+ and 3+. We show that these contributions can be identified by RPES spectra (see Figure 17 section 3.7). The RPES spectrum reflects Ti 2p 3d excitations on Ti3+ sites in the band gap and mainly on Ti4+ sites in the valence band.

50

4.1.4 Paper III: Li insertion in sol-gel prepared Mn doped TiO2studied by electron spectroscopy under ultra-high vacuum conditionsIn this paper we have studied lithiation of an ex situ prepared Mn modified Mn:TiO2 film and demonstrate that UHV lithium insertion is feasible even for ex situ prepared samples.

4.1.4.1 Characterisation of the sol-gel prepared Mn:TiO2 film The manganese content of this film is estimated by PES to be 9.4 0.4%, which is in reasonable agreement with the predicted 7% from the preparation method. Ti is identified to be in its 4+ oxidation state and Ti L-edge XAS spectra show that the as prepared film consists of anatase TiO2. Mn L-edge spectra are found to be a superposition of two contributions originating from Mn3+ and Mn2+. The relative amount of the two species contained in the film is determined to be 0.6 Mn3+ to 0.4 Mn2+. This is in agreement with the be-haviour of lithium insertion discussed later. Comparison with a second sol-gel prepared pure TiO2 film allowed us to determine the Mn modified va-lence band edge position. To this end the spectrum of the pure sol-gel film (dashed line) was subtracted from the Mn:TiO2 film (thin solid line). As shown in Figure 22 this eliminates most contributions of Ti and O and the difference spectrum (thick solid line) is then dominated by Mn 3d states and the valence band edge easily determined. We find that the band gap de-creases upon Mn modification by about 1 eV from 3.5 eV in pure TiO2 to 2.5 eV for the Mn:TiO2 film.

Figure 22. Photoemission spectra showing the valence band edge for the sol-gel prepared Mn:TiO2 (solid line) and TiO2 (dashed line) films. The thick solid line represents the difference spectrum, in which the valence band edge is indicated.

51

4.1.4.2 Lithiation of the Mn:TiO2 film Upon lithiation a preferential reduction of Mn3+ to Mn2+ over Ti4+ to Ti3+

occurs. This is observed by a drastic change in the Mn L-edge XAS spectra. After this initial deviant behaviour the effects of further lithium insertion agree very well with the previously established one for the fully UHV pre-pared Li:TiO2 system. This similarity is illustrated in Figure 23 where the evolution of the contribution of the Ti3+ component to the Ti 2p signal is shown for the Mn modified film (circles) and the model system (crosses) on a common scale.

Figure 23. The evolution of the relative contribution of the Ti3+ component to the Ti 2p signal as a function of relative lithium content. The circles represent values for insertion into Mn:TiO2, while the crosses represent values for insertion into the UHV-CVD grown TiO2 film. The solid line represents semi-empirical values, calcu-lated from the estimated film thickness and Li evaporation rate.

It can be seen that both processes are nearly identical distinguished only by an offset in lithium content for the Mn:TiO2 film due to the preferential re-duction of Mn3+. Lithium can be inserted up to a maximum value of about x=0.4 Li/Ti. The formation of the two different phases is observed and as in Paper II a delineation of the Ti L-edge XAS spectra was successfully per-formed and a good agreement between the PES and XAS determined x-value found.

52

4.2 Core hole effects in XASAs mentioned earlier the creation of a core hole can have a significant effect on the final state energy. It is thus important to study core hole effects in order to gain a more complete understanding of core level spectroscopy.

4.2.1 Paper IV: Threshold effects in the O 1s x-ray absorption spectrum of TiO2

Here theoretical calculations combined with experimental data are evaluated to further the understanding of the electronic levels of rutile and anatase TiO2. The effect of the creation of an O 1s core hole on the unoccupied states is estimated by ab initio calculations and the implications of the unoccupied DOS on electron spectroscopy experiments are investigated.

4.2.1.1 Core hole effect on the unoccupied DOS The creation of a core hole on an oxygen site according to the XAS process (see section 3.4) leads to a nearly complete filling of the O 2p shell. This has two effects on the final state DOS. First the filling of the shell decreases the unoccupied DOS and second an extensive dynamical effect induces a redis-tribution of the p-projected DOS so that a singularity is built up at threshold. As a result the leading edge of the O 1s XAS spectrum is subject to broaden-ing effects of the same type as found in the O 1s core level photoemission spectrum. Thus the leading peak rather than the leading edge of the O 1s XAS spectrum defines the unoccupied O p DOS threshold. This also leads to the leading edge of the O 1s XAS and the corresponding photoemission spectrum being nearly identical in shape. This phenomenon is expected to be a general property of transition metal oxides.

4.1.2.2 Unoccupied DOS effect on electron spectroscopy Interesting information can be gained by comparing the O 1s XAS and PES spectra on a common energy scale. Here the XAS spectrum is calibrated according to the true photon energy and the PES spectrum is referenced to the Fermi level. The PES spectrum thus shows the energy difference of the O 1s core level from the Fermi level, whereas the XAS spectrum reveals the distance to the unoccupied oxygen DOS. Assuming that the conduction band is aligned with the Fermi level, the energy position of the PES peak repre-sents the conduction band minimum (CBM) in the XAS spectrum.

53

Figure 24. O 1s XAS and PES spectra for (a) rutile and (b) anatase TiO2 on a com-mon energy scale.

As shown in Figure 24, we investigate the XAS and PES spectra for rutile and anatase TiO2. We observe that the positions of the leading XAS peak and the PES peak match for rutile and lightly lithium doped anatase, whereas the peak positions do not coincide for pure anatase. According to the picture drawn above this means that the CBM and the Fermi level coincide with the onset of the unoccupied O 2p DOS for the first two cases, whereas for pure anatase there is an energy difference of 0.3 eV. Comparing with the ab initio calculations of the unoccupied DOS we find the explanation for this behav-iour. For rutile the onset of the empty p and d partial density of states coin-cides, thus the Fermi level is located at the O p threshold. For pure anatase the situation is different. An offset of about 0.6 eV between the thresholds of the unoccupied Ti d and O p DOS is found in the DFT calculations. The smaller offset observed in the experimental spectrum is attributed to defects in the film. For the different anatase TiO2 films measured in this thesis the offset varies between 0.25 eV and 0.35 eV, depending on the number of defects in the film. It is suggested that defects and lithium insertion fill the empty Ti states from the CBM inducing a rigid shift effectively removing this offset. As observed earlier in papers I-III the Ti states giving rise to this effect are completely filled already at lithium concentrations of x 0.02.

These findings are also applicable in the following studies on ZrO2, al-though differences in the unoccupied DOS have to be taken into account. For example no offset between the unoccupied Zr d and O p DOS thresholds is found.

54

4.3 Investigation of UHV-CVD deposited ZrO2 on Si In the next four papers the properties of ZrO2 films and their influence on the substrate have been examined for UHV-CVD film growth experiments on Si. ZrO2 is a candidate for replacing SiO2 as the gate dielectric in MOSFETs, thus knowledge of the properties of ultrathin ZrO2 films on Si is very impor-tant. Of special interest are such qualities as film structure and stability, in-terface formation and the band alignment relative to Si. For instance there should be an offset of at least 1 eV between the valence and conduction bands of Si and the gate oxide in order to function properly [51].

4.3.1 Paper V: Ultra-high vacuum metal organic chemical vapor deposition of ultrathin ZrO2 films on Si(100) and Si(111) studied by electron spectroscopy In this paper film growth of ultrathin ZrO2 films by UHV-CVD of zirconium (IV) tert butoxide (ZTB) has been studied on the silicon surfaces Si(100)-(2x1) and Si(111)-(7x7). Deposition was performed at 400°C sample tem-perature, where the films are expected to grow in tetragonal phase. Special attention has been paid to the thickness dependent structure and the amount and origin of carbon contamination of the grown films.

In both cases growth of ZrO2 is observed with Zr in its 4+ oxidation state. Carbon contamination is due to decomposition of the precursor. Photoemis-sion spectra of the C 1s core level for increasing ZrO2 film thickness show the different behaviour of the Si(100) (Figure 25(a)) and Si(111) (figure 25(b)) surface.

55

Figure 25. C 1s photoemission core level spectra for increasing ZrO2 thickness on Si(100) (a) and Si(111) (b).

The peaks denoted C3 and C4 are assigned to two carbon species found within intact t-butoxy. The other peaks are assigned to species originating from t-butoxy fragmentation. It is now seen that the carbon contamination in the first stage of ZrO2 growth on the Si(100) surface is mainly due to highly fragmented precursor molecules. The decrease of the C 1s species signal intensity upon growth of further ZrO2 is consistent with dampening from an overlayer (see section 3.3.1). It is thus concluded that, due to the reactivity of the Si(100) surface, the precursor molecule are fragmented and the reaction products cannot desorb effectively and are thus incorporated into the first few Ångström of the film. Further deposition results in the growth of pure ZrO2. On the Si(111) surface however the initial contamination is due to intact t-butoxy. The decrease of this species between the first deposition points is unproportional to the additional film thickness, thus it is suggested that the remaining t-butoxy groups are not stable and can react with addi-tional ZTB forming desorbable species.

O 1s XAS data has been collected for the growth series on both surfaces and analysed with regard to the crystal structure of the ZrO2 films. The ex-pected tetragonal phase is formed for both surfaces, however beginning from very different thickness. For the Si(100) surface the tetragonal structure is observed already at a film thickness of only 11 Å, whereas the corresponding thickness on Si(111) is 51 Å.

56

4.3.2 Paper VI: Growth of ultrathin ZrO2 films on Si(100): Film-thickness-dependent band alignment In this paper the influence of the film thickness on the substrate and on the alignment of the ZrO2 valence and conduction bands relative to the Si bands is determined.

At the clean Si(100)-(2x1) surface band bending due to dangling bonds occurs, which leads to a Fermi level position in the middle of the band gap. It is shown that flatband conditions in the Si substrate are reached after deposition of a film of 25 Å thickness. This is associated with the comple-tion of the interface region and in agreement with the interface thickness estimated in Paper VII.

The band edge positions of the Si substrate and the ZrO2 film have been determined relative to the Fermi level employing XAS and PES core level spectroscopy. As discussed in Paper IV the O 1s PES peak, plotted on the same scale as the O 1s XAS spectrum, gives a measure of the offset between the Fermi level and the conduction band edge. This is illustrated in Figure 26, where the alignment of the O 1s PES peak relative to the Fermi level with the O 1s XAS spectrum is shown for a 19 Å thick film.

Figure 26. Bottom: O 1s core level photoemission spectrum after deposition of 19 Å ZrO2 on Si(100) relative to the Fermi level. Top: The corresponding O 1s XAS spec-trum (dotted) and the PES spectrum (solid line) shifted in order to match the leading O 1s XAS feature.

57

It has to be noted that the Fermi level does not coincide with the Si con-duction band, thus, for a determination of the conduction band offset be-tween ZrO2 and Si, the Fermi level offset with regard to the Si conduction band has to be taken into account. In table 1 the offsets between the respec-tive valence and conduction bands of ZrO2 and Si are displayed:

Table 1. ZrO2 to Si band offsets for increasing film thickness

Film thickness (Å) VB offset (eV) CB offset (eV)

4 3.37 0.98 11 3.39 1.01 19 3.48 0.98 27 3.47 0.94 37 3.53 0.89 51 3.61 0.77 74 3.65 0.73

From the values in this table it is evident that the band alignment is less symmetric for thicker films. Also the band offset is somewhat smaller than desired for gate oxide applications.

58

4.3.3 Paper VII: Band alignment at the ZrO2/Si(100) interface studied by photoelectron and x-ray absorption spectroscopy This paper is a continuation of the growth studies of ZrO2 on Si(100)-(2x1). A proper description of the Zr levels in ZrO2 on Si is introduced. ZrO2 is found to be an insulator under the entire film growth. Thus the Zr levels are best referenced to the local vacuum level.

From a comparison of the evolution of the relative intensity of O 1s and Zr 3d photoemission peaks taken from overview spectra it is concluded that an interface about 20 Å thick is formed. Further deposition leads to the growth of pure tetragonal ZrO2. This is in good agreement with the previous papers (V and VI).

The sample work function has been established as a function of film thickness. It is found that the work function decreases from 5.1 eV for the pure Si(100)-(2x1) surface to about 4 eV for a 62 Å thick ZrO2 film. It is observed how the Si 2p and Zr 3d core level binding energies increase with increasing film thickness. In the case of Si this behaviour is explained by the annihilation of the band gap states due to dangling bonds (see Paper VI) and is thus related to the Fermi level. For zirconium however, the situation is more complex. It is found that Zr is electronically decoupled from the sub-strate already from the first deposition point onwards and the zirconium lev-els are thus better referenced to the local vacuum level. This behaviour is shown in Figure 27, where the evolution of the Zr 3d5/2 core level binding energy relative to the substrate Fermi level (circles, left scale) and the local vacuum level (triangles, right scale) as a function of film thickness is dis-played.

Figure 27. The Zr 3d5/2 binding energy as a function of film thickness. Solid circles: Relative to the Fermi level (left scale); Solid triangles: Relative to the local vacuum level (right scale)

59

It is evident that the Zr binding energy is constant relative to the local vac-uum level throughout the entire experiment. Thus the zirconium electronic structure follows changes in the work function independent of the substrate. It is noteworthy that this is already the case from the first deposition point onwards, thus the formation of a mixed interface does not inhibit the insulat-ing properties of ZrO2 films.

In resemblance with the method presented in Paper IV the band gap can be studied by combining O 1s XAS and PES with ab initio calculations of the unoccupied DOS. The ab initio calculations show about the same dy-namical core hole effect with accompanying singularity formation on the O p DOS threshold. However no offset is found between the Zr and O unoccu-pied states. In this case the O 1s XAS spectrum calibrated according to the true photon energy and the O 1s PES spectrum referenced to the Fermi level are plotted on the same energy scale. Here it is found that the position of the PES O 1s relative to the leading O 1s XAS peak signifies the position of the Fermi level relative to the conduction band edge.

60

4.3.4 Paper VIII: Combinatorial chemical vapour deposition of an ultrathin ZrO2-TiO2 film on Si(100)-(2x1) in ultra-high vacuumCompound metal oxide systems can have advantageous properties compared with monometallic oxides. Investigating all possible combinations of these compound materials using pure sample studies requires an enormous amount of experimental work. It is thus highly desirable to produce samples with graded composition that can serve as a library for the material combination in question. In this paper combinatorial CVD of TiO2 and ZrO2 under UHV conditions in ultra-high vacuum is performed. A good lateral compositional spread over the entire film from TiO2 over TiZrxOy to ZrO2 is found. This behaviour is illustrated in Figure 28, where the relative amounts of Ti (circles) and Zr (triangles) detected at a given sample position are shown as a function of sample position.

Figure 28. Relative atomic concentration of Ti (empty circles) and Zr (solid trian-gles) plotted against the position on the sample.

Here we clearly observe the graded stoichiometry from a TiO2 side at 0 mm sample position to a ZrO2 side at 8 mm sample position. The changes this spread induces in the electronic structure are most beautifully reflected in the O 1s XAS data shown in Figure 29.

61

Figure 29. O 1s XAS spectra taken at the different sample positions.

In Figure 29 it can clearly be seen how the O 1s XAS spectra change from TiO2 dominated spectra on the TiO2 side to ZrO2 dominated spectra on the ZrO2 side.

The crystal structure of TiO2 changes from anatase for pure TiO2 to amorphous upon blending with ZrO2. ZrO2 is found to be in its tetragonal phase on the Zr rich side. A clear tendency for Ti surface segregation is ob-served. In line with Paper VII the Zr electronic levels are best described rela-tive to the local vacuum level. However the preferential surface location of Ti induces changes in the properties of the ZrO2 side. The isolated surface titanium also becomes electronically decoupled from the substrate and fol-lows the Zr levels and the local vacuum level. Modification of the local ZrO2work function leads to an improved, i.e. more symmetric, alignment of the bands relative to Si. Thus modified ZrO2 might be used as gate oxide mate-rial in MOSFETs.

62

4.4 Magnetism A field of major current interest is the search for diluted magnetic semicon-ductors (DMS). The exact identification of magnetic signals is not trivial and for many reported DMS there is a discord as to the origin of the magnetism. For a DMS the magnetic signal must not be due to the clustering of an atomic species that was magnetic to begin with. It is thus important to be able to identify the chemical state of the magnetic species. Core level elec-tron spectroscopy can accomplish just that by means of x-ray magnetic cir-cular dichroism (XMCD) (see section 3.5).

4.4.1 Paper IX: Electronic structure investigation of (Zn,Co)O room temperature ferromagnets In this paper the magnetic properties of (Co,Zn)O films produced by electro-deposition have been investigated. The films had been predicted to exhibit room temperature ferromagnetism by means of SQUID magnetometry. To gain insight into the nature of the atomic species giving rise to the observed RT magnetism the samples were studied by XMCD.

Figure 30. Co L-edge XMCD spectra and their difference spectrum (circles)

63

In Figure 30 the XMCD spectra for majority and minority charge carriers and the XMCD difference spectrum are displayed. The XMCD spectra re-veal weak magnetism confirming the SQUID results and the multiplet struc-ture of the L3 absorption edge shows that cobalt is present mainly as Co2+

with oxygen nearest neighbours. The XMCD difference spectrum allows to identify the chemical state of the species giving rise to magnetism. The mul-tiplet structure in the difference spectrum, consistent with Co in its 2+ oxida-tion state, is clearly visible, whereas the presence of metallic cobalt is not observed but cannot be fully excluded. It is thus concluded that at least a dominating contribution to the magnetic signal is due to substitutional Co with oxygen nearest neighbours. Thus it is demonstrated that XMCD, due to the elemental and chemical specificity of the magnetic signal, is ideal for studying DMS material.

64

5 Populärvetenskaplig sammanfattning

5.1 Inledning Metalloxidsystem ingår som viktiga komponenter i många praktiska applika-tioner. Titandioxid, TiO2, som har studerats mycket i första delen av den här avhandlingen har redan många användningsområden och det finns ett flertal potentiella nya applikationer. Två tänkbara användningsområden är till ex-empel energilagring och energiproduktion. TiO2 kan lagra laddning i form av joner och kan därmed användas i litiumjonbatterier. TiO2 används också i en typ av solcell, den så kallade nanoporösa färgämnessolcellen, där solenergi tas upp av ett färgämne som sitter på titandioxiden. TiO2 är också tänkt att kunna användas i ”smarta” fönster. Dessa typer av fönster minskar framför-allt uppvärmning av hus, vilket är ett stort problem världen över. Minskat behov av att kyla ner fastigheter under varma soliga dagar medför enorm energibesparing. Titandioxid är genomskinligt när det är rent, men får en djupblå färg med litiumjoner i sig. Tanken är därför att applicera titandioxid-film på fönsterglas och att ha ett system för att sätta in och ta bort litiumjo-ner. Mestadels skulle ett sådant fönster fungera som ett vanligt fönster men när solstrålningen blir för stark färgas fönstret blått. På så sätt strålar mindre solljus, och därmed mindre värme, in i rummet. Dessutom kan titandioxid i vissa fall sönderdela molekyler som sitter på dess yta (som t.ex. i en kataly-sator), vilket kan vara användbart i många sammanhang.

Ett annat viktigt forskningsområde är nya material för transistorer. I da-gens transistorer är ett kiseldioxidlager med bland annat en isolerande funk-tion viktigt för att kunna styra strömmen genom två delar av transistorn. Transistorerna byggs allt mindre och då måste också kiseldioxidlagret bli tunnare för att behålla de egenskaper som transistorn behöver för att fungera. Tjockleken på detta kiseldioxidlager är nu cirka 12 Å. Om kiseldioxidlagret görs ännu tunnare går den isolerande funktionen förlorad. Det är därför nöd-vändigt att byta ut kiseldioxiden mot ett material som kan byggas tjockare men fortfarande ha samma egenskaper som det tunna kiseldioxidlagret. Po-tentiella material måste ha en hög dielektricitetskonstant. Dielektricitets-konstanten är ett mått på hur materialet påverkas av elektriska fält. I stort sett är i detta fall så hög som möjligt bäst. Zirkondioxid och igen titandioxid är

65

exempel på sådana material. I den senare delen av avhandlingen studeras dessa material med transistorapplikationen i baktankarna.

I den sista delen av avhandlingen studeras magnetiska halvledarmaterial. Halvledare är den viktigaste byggstenen inom elektroniken och lyckas man få en magnetisk halvledare, så är en helt ny typ av elektronik möjlig, som baseras på magnetiska egenskaper. Magnetiska halvledare är en helt ny typ av material som har upptäckts bara de senaste åren och den grundläggande förståelsen för dessa material är begränsad inom forskningsvärlden. Det är därför viktigt att studera sådana material både på en grundläggande och på en mer applicerad nivå. Studier som leder till bättre förståelse av de ovan-nämnda metalloxidsystemens egenskaper kan leda till förbättringar i existe-rande applikationer eller till helt nya användningsområden. Det är därför viktigt att få en så bra bild som möjligt över alla egenskaper som dessa mate-rial har.

5.2 Resultaten av min forskning Jag har använt mig av elektronspektroskopiexperiment för att studera mina material. I elektronspektroskopi belyses ett prov med energirikt ljus (rönt-genstrålning) och man mäter egenskaperna hos de elektroner som ljuset slår ut ur materialet. Dessa så kallade fotoelektroner tillåter en att förstå egen-skaper av det studerade materialet. Jag har utfört alla mina experiment i ult-rahögvakuum. Ultrahögvakuum är en artificiell miljö och detta måste man ta hänsyn till när vi drar paralleller med experiment gjorda i atmosfär. Den stora fördelen med ultrahögvakuum är dock att det inte finns något runt om-kring våra prover som kan påverka våra mätningar. Allt som ses är således en följd av de egenskaper som materialet har, vilket naturligtvis medför en större grad av förståelse för de resultat som fås.

I de tre första artiklarna har jag studerat litiuminsättning i olika titandiox-idsystem. Det visar sig att en stor mängd litium (ungefär en Li atom per 2.5 Ti atomer i filmen) kunde infogas i anatas titandioxid (anatas är en speciell kristallfas för titandioxid). Jag har tagit fram resultat som laddningsöverfö-ring och den elektroniska och geometriska konfigurationen. Kunskaper om laddningsöverföring är förstås särskilt viktigt för den potentiella batteriap-plikationen. Förståelse för materialordningen är också en viktig aspekt. Det är tänkbart att materialordningen förändras så att titandioxiden expanderar när litium sätts in. I en batteriapplikation skulle det leda till stora problem om man inte tog hänsyn till sådana effekter. Jämför t.ex. med en fylld vatten-flaska som sprängs i frysen när vattnet omvandlas till is.

I de därefter följande artiklarna har jag studerat uppväxt av zirkondioxid, hur den beter sig under uppväxten och hur dess elektroniska nivåer beskrivs jämfört med titandioxid. Jag har först studerat en zirkondioxidfilm på kisel. Jag har varit intresserad av hur tjockleken påverkar materialets egenskaper

66

och har studerat en uppväxtserie från bara några få atomer på kiselsubstratet till en tunn film. Avslutningsvis har möjligheten studerats att växa två mate-rial (zirkondioxid och titandioxid) samtidigt. Det har lyckats att skapa en fördelning över provet från ren titandioxid över en blandning av titan- och zirkondioxid till ren zirkondioxid. Jag ville se huruvida det överhuvudtaget är möjligt och vilka egenskaper dessa material skulle få beroende på graden av blandning (eller renhet). Från filmuppväxten av ren zirkondioxid på kisel har jag lärt mig att zirkondioxid är en elektrisk isolator redan från första början av filmuppväxten. I situationer där det endast finns enstaka zirkondi-oxidmolekyler på provytan är dessa elektriskt isolerade från kisel och zir-kondioxid förblir sedan en isolator under hela uppväxten. Det leder till nöd-vändigheten av en mer komplicerat beskrivning av de elektroniska nivåerna. I första studien av ren zirkondioxid har det tagits fram hur dess elektron-struktur ser ut jämfört med kisels elektronstruktur. Från jämförelsen upp-täcktes att det troligtvis kan läcka ström genom zirkondioxidlagret vilket gör det osannolikt att zirkondioxid kan användas rakt av. Men det visade sig att egenskaperna kunde förbättras om man blandar lite titandioxid i zirkondiox-iden. Detta betyder att det skulle kunna vara möjligt att använda zirkondiox-id med lite titandioxid inblandat som ersättning för kiseldioxid. Blandfilm-experimentet visade att det var möjligt att bygga filmen som jag hade tänkt mig, med en gradvis övergång från zirkondioxid till titandioxid. Jag observe-rade dock att ingen perfekt blandning åstadkommits, utan att titan tycker mera om att sätta sig på ytan än zirkon. Det leder till att fördelningen av atomer i djupled inte är jämn, med mer titan på ytan och mer zirkon djupare ner.

I sista artikeln har det studerats material som misstänktes kunna vara magnetiska halvledare. Sådana material är blandningar av olika substanser och det ska vara själva blandningen som gör att det är magnetiskt. Det ska inte vara så att blandningen är dålig och materialet är magnetiskt för att ett material var magnetiskt från början och har samlats på ett ställe. Ungefär jämförbar med chokladmjölk, där hela drycken ska smaka choklad och inte bara ett bottenlager av chokladpulver. Genom experiment visades att materi-alet verkligen var magnetiskt. Det verkligt intressanta var att det också kun-de visas att magnetismen inte berodde på en ansamling av det magnetiska materialet, metallisk kobolt i det här fallet. Det som var magnetiskt var istäl-let kobolt med syre som närmaste grannar, vilket bevisar att substanserna hade blandats ordentligt.

67

6 Acknowledgements

First of all I would like to thank Hans Siegbahn, for giving me the opportu-nity to work in his group, my supervisors Anders Sandell and Håkan Ren-smo, for their support and interesting discussions scientific or otherwise. My heartfelt gratitude goes to my co-workers through all this time, who made these five years a very enjoyable experience: Patrik, Erik, Emma and Maria. Thanks also to my short-term office room mates Ylvi and Pia. I would also like to thank Petter, for being a good friend during all my years in Sweden, IF Trikadien, my parents and life in general, for being good to me. And fi-nally, special thanks to my dearest Elina. Kiitos Pupu!

68

7 Bibliography

1 S. Y. Huang, L. Kavan, I. Exnar et al., J. Electrochem. Soc. 142, L142 (1995) 2 B. O’Regan and M. Grätzel, Nature 353, 737 (1991) 3 C. Bechinger, S. Ferrere, A. Zaban, J. Sprague and B. A. Gregg, Nature 383, 608 (1996) 4 R. van de Krol, A. Goossens and E. A. Meulenkamp, Appl. Phys. 90, 2335 (2001) 5 A. Hagfeldt and M. Grätzel, Chem. Rev. 95, 49 (1995) 6 Y. Matsumoto, et al., Science 291, 854 (2001)7 S. A. Chambers, et al., Appl. Phys. Lett. 79, 3467 (2001) 8 www.intel.com (2006) 9 G. D. Wilk, R. M. Wallace and J. M. Anthony, J. Appl. Phys. 89, 5243 (2001) 10 S. A. Campbell, D. C. Gilmer, X. Wang, M. Hsieh, H.-S. Kim, W. L. Gladfelter and J. Yan, IEEE Trans. Electron Devices ED-44 104 (1997) 11 J. P. Chang, Y.-S. Lin, S. Berger, A. Kepten, R. Bloom and S. Levy, J. Vac. Sci. Technol. B 19, 2137 (2001) 12 B.-O. Cho and J. P. Chang, Appl. Phys. Lett. 93, 745 (2003) 13 T. S. Leon, J. M. White and D. L. Kwong, Appl. Phys. Lett. 78, 368 (2001) 14 Y.-S. Lin, R. Puthenkovilakam, J. P. Chang, C. Bouldin, I. Levin, N. V. Nguyen, J. Ehrstein, Y. Sun, P. Pianetta, T. Conard, W. Vandervorst, V. Venturo and S. Sel-brede, J. Appl. Phys. 93, 5945 (2003) 15 K. Muraoka, Appl. Phys. Lett. 80, 4516 (2002) 16 D. J. Burleson, J. T. Roberts, W. L. Gladfelter, S. A. Campbell and R. C. Smith, Chem. Mater. 14, 1269 (2002) 17 The Max-lab homepage: www.maxlab.lu.se 18 J. N. Andersen, O. Björneholm, A. Sandell, R. Nyholm, J. Forsell, L. Thånell, A. Nilsson, N. Mårtensson, Synchrotron Radiat. News 4, 15 (1991) 19 R. Nyholm, J. N. Andersen, U. Johansson, B. N. Jensen and I. Lindau, Nucl. In-strum. Meth. A 467-468, 520 (2001) 20 P. Atkins, J. de Paula, Atkins’ Physical Chemistry, 7th edition (Oxford University Press 2002) 21 D. A. Neamen, Semiconductor Physic and Devices, (Irwin 1992) 22 A. Einstein, Annalen der Physik 17, 132 (1905) 23 J. J. Yeh and I. Lindau, Atomic and Nuclear Data Tables 32, 1 (1985) 24 K. L. Choy, Progress in Materials Science 48, 57 (2003) 25 M. Ritala, K. Kukli, A. Rahtu, P. I. Räisänen, M. Leskelä, T. Sajavaara and J. Keinonen, Science 288, 319 (2000) 26 A. C. Jones and P. R. Chalker, J. Phys. D: Appl. Phys. 36, R80 (2003) 27 The SAES Getters Group homepage: www.saesgetters.com/default.aspx?idPage=470 28 I. A. M. Pohl, L. G. Westin and M. Kritikos, Chem. Eur. J. 7, 3438 (2001)

69

29 G. Westin, K. Jansson, A. Pohl and M. Leideborg, J. Sol-Gel Sci. Technol. 31, 25 (2004) 30 T. L. Barr, Applications of Electron Spectroscopy to Heterogeneous Catalysis In: D. Briggs and M. P. Seah Practical Surface Analysis 2nd ed. Vol 1, John Wiley & Sons, Chichester, England (1996) 31 H. Siegbahn and L. Karlsson, Handbuch der Physik encyclopedia of physics: photoelectron spectroscopy, Springer, Berlin (1982) 32 N. Mårtensson and A. Nilsson, J. El. Spectr. Rel. Phen. 75, 209 (1995) 33 K. Siegbahn, Philosophical Transactions of the Royal Society of London A (Mathematical and Physical Sciences) 268, 33 (1970) 34 S. Hüfner, Photoelectron Spectroscopy, Springer, Berlin Heidelberg (1995) 35 D. R. Penn, Phys. Rev. B 13, 5248 (1976) 36 J. Stöhr, J. Elec. Spectros. Rel. Phenom. 75, 253 (1995) 37 P. Carra, B. Thole, M. Altarelli and X. Wang, Phys. Rev. Lett. 70, 694 (1993) 38 B. Thole, P. Carra, F. Sette and G. van der Laan, Phys. Rev. Lett. 68, 1943 (1992) 39 D. Arvanitis, M. Tischer, J-. Hunter Dunn, F. May, N. Mårtensson and K. Baber-schke, Lecture Notes in Physics 466, 145 (1996) 40 P. Hohenberg and W. Kohn, Phys. Rev. 136 B864 (1964) 41 W. Kohn and L. J. Sham, Phys. Rev. 140 A1133 (1965) 42 A. Henningsson, M. P. Andersson, P. Uvdal, H. Siegbahn and A. Sandell, Chem. Phys. Lett. 360, 85 (2002). 43 F. Bonino, L. Busani, M. Lazzari, M. Manstretta, B. Rivolta and B. Scrosati, J. Power Sources 6, 261 (1981). 44 B. Zachau Christiansen, K. West, T. Jacobson and S. Atlung, Solid State Ionics 28(30), 1176 (1995) 45 A. Henningsson, H. Rensmo, A. Sandell, H. Siegbahn, S. Södergren, H. Lind-ström and A. Hagfeldt, J. Chem. Phys. 118, 5607 (2003) 46 H. Lindström, S. Södergren, A. Solbrand, H. Rensmo, J. Hjelm, A. Hagfeldt and S. E. Lindquist, J. Phys. Chem. B 101, 7717 (1997). 47 M. Wagemaker, R. van de Krol, A. P. M. Kentgens, A. A. van Well, F. M. Mul-der, J. Am. Chem. Soc. 123, 11454 (2001) 48 R. J. Cava, D.W. Murphy, S. Zahurak, A. Santoro and R. S. Roth, J. Solid State Chem. 53, 64 (1984). 49 M. V. Koudriachova, S.W. de Leeuw and N. M. Harrison, Phys. Rev. B 69,054106 (2004). 50 R. van de Krol, A. Goossens and E. A. Meulenkamp, J. Electrochem. Soc. 146,3150 (1999). 51 G. D. Wilk, R. M. Wallace and J. M. Anthony, J. Appl. Phys. 89, 5243 (2001)

Acta Universitatis UpsaliensisDigital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 228

Editor: The Dean of the Faculty of Science and Technology

A doctoral dissertation from the Faculty of Science andTechnology, Uppsala University, is usually a summary of anumber of papers. A few copies of the complete dissertationare kept at major Swedish research libraries, while thesummary alone is distributed internationally through theseries Digital Comprehensive Summaries of UppsalaDissertations from the Faculty of Science and Technology.(Prior to January, 2005, the series was published under thetitle “Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology”.)

Distribution: publications.uu.seurn:nbn:se:uu:diva-7180

ACTAUNIVERSITATISUPSALIENSISUPPSALA2006