Template controlled octahedral rotations and the metal-insulator transition in SrRuO 3

Transcription

1 Master s Thesis Template controlled octahedral rotations and the metal-insulator transition in SrRuO 3 P. C. Spruijtenburg Enschede, April 1, 2011 Applied Physics University of Twente Faculty of Science and Technology Inorganic Materials Science Graduation committee: Prof. dr. ing. G. Rijnders Dr. ir. G. Koster Prof dr. ir. A. Brinkman J.E. Kleibeuker M.Sc. B. Kuiper M.Sc.

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3 Abstract Material properties are directly linked to their atomic structure. In the case of material properties of thin films, attention up until now has mainly focused on the effect of induced strain. Recently the role of oxygen octahedra in complex perovskite oxides has garnered interest. The oxygen octahedra in SrRuO 3 have been linked to its conducting properties and the thickness dependent metal-insulator transition (MIT) that occurs for very thin films of SrRuO 3. By growing intermediate layers between SrRuO 3 and the substrate, it is postulated that octahedral rotations, and thus the MIT in SrRuO 3 can be influenced. The intermediate material of infinite-layer SrCuO 2 is proposed and investigated as a way to decouple octahedral rotations from the substrate to the film. Insulating behaviour for thick films of SrRuO 3 up to 15 monolayers thick is shown when grown on SrCuO 2 as an intermediate layer. This can be linked to possible Cu interdiffusion leading to a ruthenium poor film, and/or a change in unit cell symmetry and volume.

4 Table of contents 1 Introduction Outline General theoretical aspects Perovskites Octahedral rotations Strain and template mediated octahedral rotations in perovskites Strain mediated octahedral rotations Template mediated octahedral rotations Metal-Insulator transitions Mott-transition Mott-Hubbard transition SrTiO 3 /SrRuO Core-level XPS satellites in Mott-Hubbard type systems Hypothesis Goals Experimental methods In or ex situ? Heteroepitaxy Pulsed laser deposition Growth modes Reflection high energy electron diffraction Materials properties of LaAlO 3, DyScO 3, SrRuO 3, and SrTiO Atomic Force Microscopy X-ray Photoelectron Spectroscopy Measurements and Setup Simulations Transport measurements XRD SrTiO 3 substrate treatment i

5 Table of contents 5 SrCuO Structure Growth XRD Photoelectron spectroscopy Measurements Simulations Discussion Conclusions The metal-insulator transition in SrRuO Heteroepitaxial structures Growth SrRuO LaAlO DyScO Fitting the Ru 3d photoelectron spectrum Thickness series SrCuO Surface effects SrTiO SrCuO DyScO LaAlO Sandwiching structures of SrRuO SrRuO 3 and LaAlO The screened peak in all sandwich intermediate materials Transport measurements Insulating behaviour of SrCuO 2 /SrRuO 3 /SrCuO 2 sandwich structures Reciprocal space-mapping Discussion Conclusions Discussion and Conclusions Recommendations References 56 Appendices 58 A Atomic structure files for calculations of XPS Ruthenium 3d charge transfer spectra 59 B Uncapped SrRuO 3 61 ii

6 Chapter 1 Introduction In recent years, techniques such as pulsed laser deposition and molecular beam epitaxy have helped to gain precise control over thin film heterostructures. Thin films with a thickness of only one unit cell can be fabricated. This control allows for the creation and investigation of novel materials and their properties. The subsequent characterisation of these thin films is an integral part of materials science. Many different properties, such as crystal structure, conductivity, or magnetic properties often need to be taken into account to form a complete picture of the way materials work. Techniques such as atomic force microscopy, x-ray photoelectron spectroscopy, x-ray diffraction, and many others each play their part in this characterisation and in the discovery of new physics in materials. Conductive oxides are essential components in most composite oxide heterostructures, where they are often used as electrode materials. The itinerant ferromagnetic SrRuO 3 is a popular choice for electrodes since it is one of the more conductive metallic oxides with good thermal properties[1]. SrRuO 3 is unusual among the correlated complex oxides, because it is a metallic itinerant ferromagnet. It exhibits a transition to a ferromagnetic state at T c = 160 K that was shown to be dominated by transverse fluctuations of robust local moments of size 1.6µ B [2]. At high temperatures, it is a so called bad metal, but at low temperatures behaves like a Fermi liquid. SrRuO 3 further has an orthorhombically distorted cubic perovskite crystal structure at room temperature. The structural and chemical similarities between perovskites mean they are well suited for use in oxide heterostructures. SrRuO 3 can thus easily be combined with other perovskites in heterostructures. In bulk, the crystal structure of SrRuO 3 has been shown to change from orthorhombic to tetragonal to cubic symmetry at 547 and 677 C, respectively[3]. A ferroelectric field effect on the conductivity of SrRuO 3 has also been observed[4]. Correlation effects present in SrRuO 3 have been show to be dependent on orthorhombic distortions by Ca doping of SrRuO 3 [5, 6]. Correlation effects appear to also be dependent on cation off-stiochiometry, with ruthenium poor films exhibiting less electron correlation[7]. Its conductive properties are what make SrRuO 3 an attractive electrode material. However, it has been shown for SrRuO 3 on SrTiO 3 that for extremely small thickness (3 monolayers), SrRuO 3 becomes insulating and no longer exhibits ferromagnetism[2]. This thickness dependent metal to insulator transition could give important clues as to the nature of the mechanism for conduction in SrRuO 3. The perovskite structure of SrRuO 3 means that it contains so-called oxygen octahedra, consisting of a ruthenium atom surrounded by 6 oxygen atoms at the ends of each axis. In perovskites, the rotations of these octahedra can vary depending on the ionic radii of the cations. The variation of these rotations has been shown to be at least dependent on strain[8]. At small thicknesses however, it could be possible that strain is not the most dominant driving force. Besides strain, it could be possible that octahedra rotations in the film are influenced by the rotations already present in the layer on which the thin film is grown. The rotations of these octahedra could play a significant role in the conduction mechanism in SrRuO 3 and could explain why the metal-insulator transition appears for small thicknesses. This research will try to influence the way oxygen octahedra interact between the SrRuO 3 thin film and the SrTiO 3 substrate, by growing layers between the substrate 1

7 Chapter 1: Introduction and the thin film. 1.1 Outline The outline of this report will be as follows. Chapter 2 will explore the general theoretical aspects of this research. Chapter 3 will concern the hypothesis and goals that will be the mainstay of this research. Experimental aspects such as measurement techniques and the particulars of doing simulations on x-ray photoelectron spectroscopy are discussed in chapter 4. Results of the experiments performed are presented in chapters 5 and 6. The last chapter, chapter 7, concludes this report with a discussion of all results and recommends possible further avenues of exploration and research on this topic. 2

8 Chapter 2 General theoretical aspects 2.1 Perovskites The general perovskite crystal structure can be defined as a structure with chemical formula ABX 3, where A and B are two cations of different sizes, and X is an anion that bonds to both A and B[9]. In the ideal cubic symmetry, the B cations are surrounded by an octahedron of X anions in 6-fold coordination. In the cubic unit cell, the A-atom will sit at the corners of the unit cell at position (0, 0, 0), the B-atom in the body centre at position (1/2, 1/2, 1/2), and the X anion can be found at face-centre positions (0, 1/2, 1/2), (1/2, 0, 1/2). In the case of perovskite complex oxides the chemical formula is reduced to ABO 3. Because of the formation of octahedra made of X anions around the B cations, complex oxides with a perovskite structure will have BO 6 oxygen octahedra. Figure 2.1 shows the cubic unit cell of a perovskite oxide. Figure 2.1: The ABO 3 perovskite unit cell. (Image from [10]) The perovskite structure is adopted in a lot of complex oxide structures. The complex oxides SrRuO 3 as well as SrTiO 3, DyScO 3 and LaAlO 3 (all used in this research) incorporate a perovskite structure and are thus known as perovskites Octahedral rotations While the ideal case of cubic symmetry can be found in e.g. SrTiO 3 at room temperature, more often the structure is modified by cation displacements as in BaTiO 3 or by the rotation of octahedra as found in CaTiO 3, SrRuO 3, LaAlO 3, or by a combination of both as in NaNbO 3 [11]. Octahedral tilting distortions present in many perovskites were first examined by Goldschmidt in

9 Chapter 2: General theoretical aspects According to Goldschmidt, the degree of distortion in ABO3 perovskites can be determined by observing the tolerance factor t: t = R A + R O 2(RB + R O ) (2.1) where R A, R B, and R O are the ionic radii of the A-atom, B-atom and Oxygen-atom, respectively[12]. The ideal cubic perovskite has a tolerance factor of 1.0. When the atomic radius of the A-atom is smaller than ideal, the BO 6 octahedra will rotate and deform to fill the space. The rotation of the BO 6 oxygen octahedra are of importance while they are often of greater effect on the lattice parameters of materials than cation displacements. Therefore a classification of these rotations is needed. The glazer notation is an often used framework to describe these rotations and was developed by A.M. Glazer[11]. A set of three letters defines the rotations about the [001], [010], and [001] axes in the pseudo cubic unit cell, consecutively. Figure 2.2 shows these three axes. α, β, and γ are defined as the angles of rotation around these axes. Repeating letters in the notation indicate that the magnitude of rotation around the respective axes is equal. Rotations that are all unequal about all axes would thus be written as abc (α β γ), while all equal rotations would be written as aaa (α = β = γ). When rotating about one axis, the rotations of octahedra in the directions along the axis perpendicular to the axis of rotation become fixed to the opposite rotation. In this way, the positive tilt about [100] of magnitude a in the first octahedron makes a negative tilt about [100] of equal magnitude in the nearest-neighbour octahedra along [010] and [001]. The degree of freedom remaining is then the rotation of the octahedron along the axis of rotation. This is indicated by +,- or 0 for an in-phase, out-of-phase, or no rotation respectively. A full notation would look like a + b + c. Figure 2.2: Rotation axes for perovskite oxygen octahedra used in glazer notation. (Image from [8]) This notation is non-abelian in nature, meaning that the final rotation depends on the order in which the rotations are performed about the three axes. In practice this effect is limited and becomes of a second-order nature when rotations within 15 are considered. The materials with oxygen octahedra used in this work and their rotation systems can be found in table Strain and template mediated octahedral rotations in perovskites For perovskites, the lattice parameters are often of approximately the same magnitude. The lattice mismatch between these materials is small, and make them well suited for use in heterostructures. The effects of in-plane epitaxial strain has been well established for the ground-state structure and polarisation of ferroelectric perovskites[13]. However these calculations do not take into account any doubling of the unit cell due to oxygen octahedra rotations. 4

10 Chapter 2: General theoretical aspects Recently, the coupling between epitaxial strain and octahedral rotations of LaNiO 3 has been explored by determining the oxygen positions from half-order Bragg peaks in X-ray diffraction[8]. From this it is possible to infer the behaviour of octahedral rotations from a perspective that is purely based on strain Strain mediated octahedral rotations When growing an epitaxial film on a substrate, a film will experience strain when the in-plane lattice parameters of the film are unequal to the parameters of the substrate. In order to accommodate this strain, the unit cell of the grown film will deform. Because of volume continuity considerations, the out-of-plane lattice parameter of the unit cell will increase (decrease) for a compressive (tensile) biaxial strain. The repercussions for a perovskite structure and the rotations of its oxygen octahedra can then be understood through a simple geometrical argument. For the case of compressive strain the in-plane lattice parameters will decrease. Looking at the in-plane octahedron from the [001] axis, figure 2.3(a) shows that the octahedron can (at least partially) avoid a decrease in its B-O bond length by rotating about the [001] c-axis. The effect for the octahedron when viewed in the (100) or (010) plane can be seen in figure 2.3(b). The apical lattice parameter increases due to volume continuity. The octahedron will try, as it did for the in-plane rotation, to keep its bond lengths equal. Assuming the octahedron already had a rotation about the a or b axis, it can do this by rotating toward the [001] c-axis. From these arguments it thus follows that for compressive strain the α and β rotations will decrease, while γ rotations will increase. This behaviour is reversed symmetrically when going to the case of tensile strain, so α and β rotations will increase, while rotations around γ decrease. b c a a (a) (b) Figure 2.3: Effects from geometrical considerations of compressive strain on octahedral rotations in the (a) equatorial plane and the (b) apical plane. Big dots represent the B-site atoms. Smaller dots represent O-site atoms. This rotation behaviour is nicely demonstrated in figure 2.4. For increasing compressive strain (more negative strain percentage), the γ rotations increase while the α(/β) rotations decrease. For c/a=1 the unit cell is nearly cubic and a first-order phase transition occurs. A change in bond length can also be observed, with the equatorial Ni-O bond length showing a large susceptibility to changing of the lattice parameter Template mediated octahedral rotations The growth of the thin film depends strongly on the structure of the surface layer of the substrate. In the case of perovskites growing on a (001) oriented perovskite substrate, depending on the termination the film either sees a BO 2 or AO plane. With the presence of octahedral rotations, the surface morphology at an atomic scale is not the same for different rotations. In fact, the introduction of octahedral rotations can be seen as a doubling of the unit cell compared to the ideal cubic symmetry. With rotations, the oxygen atoms are being shifted when viewed in-plane and rotated along a single axis, as can be seen in a simplified 2d illustration in figure

11 Chapter 2: General theoretical aspects QUANTIFYING OCTAHEDRAL ROTATIONS IN STRAINED PHYSICAL REVIEW B 82, Ni O bond length (Å) angle (degrees) c/a γ α apical equatorial strain (%) FIG. 5. Color online a The DFT calculated rotation angles, b Ni-O bond lengths, and c the c/a axial ratio as a function of strain open symbols. A first-order transition occurs at approximately c/a=1. LNO/STO and LNO/LAO, respectively. In contrast, the outof-plane is highly sensitive to the substrate-induced strain; this bond angle decreases from in LNO/LAO to in LNO/STO. The insensitivity of the in-plane can be understood as both the or and angles contribute to the in-plane, and we find empirically that the average of and is robust to strain. Table I summarizes the bonding data for the two samples. III. DENSITY-FUNCTIONAL CALCULATIONS To investigate the origin of these effects and to determine the rotational behavior for intermediate strain values, we perform DFT calculations within the local-spin density approximation LSDA plus Hubbard U method U eff =U J =3 ev as implemented in the Vienna ab initio simulation package VASP We follow the Dudarev approach 34 and include an effective Hubbard term U eff =U J of 3 ev to accurately treat the correlated Ni 3d orbitals. The core and a b c valence electrons are treated with the projector-augmente wave method 37 with the following valence-electron configu rations: 5p 6 5d 1 6s 2 La, 3p 6 3d 9 4s 1 Ni, and 2s 2 2p 4 O. Th Brillouin-zone integrations are performed with a Gaussia smearing of 0.05 ev over a Monkhorst-Pac k-point mesh 38 centered at, and a 450 ev plane-wave cu off. For structural relaxations we also used a Gaussian broad ening technique of 0.05 ev and relaxed the ions until th Hellmann-Feynman forces are less than 1 mev Å 1. In a calculations, ferromagnetic spin order is enforced. We fir optimize the internal degrees of freedom for bulk rhomboh dral LNO 10-atom unit cell at the experimental lattice p rameters to confirm that we accurately describe the electron and atomic structure. Consistent with the experimental dat our optimized structure is indeed metallic and consists o equal and alternating phase rotations of the NiO 6 octahed about each Cartesian axis with = = =5.76, in goo agreement with the 1.5 K structural measurement of The structural optimization for the two films was carrie out using the experimentally determined lattice parameter The substrates are not explicitly simulated in our calculation but the symmetry reduction imposed by enforcing equal in plane lattice parameters on the bulk R3 c structure is include by optimizing the internal degrees of freedom within monoclinic 20-atom unit cell space group C2/c. The ex cellent consistency between our experimentally and theoret cally obtained structures is apparent in Fig. 4 a. We no that the DFT ground state indicates an additional energy low ering distortion through small antiparallel displacements o 0.01 Å for the La atoms along the 110 directions. Suc displacements would produce weak but measurable Figure 2.4: Behaviour of LaNiO 3 under strain. The (a) density functional theory (DFT) calculated rotation angles, (b) Ni-O bond lengths, and the (c) c/a axial ratio as a function of strain (open symbols). A first-order transition occurs at approximately c/a=1[8] and 3 2 peaks, which however were not observed in th diffraction measurements, and likely persists in the DFT ca culations due to the common exchange-correlation function underestimation of the equilibrium volume. We extend our analysis to intermediate strain states Figure 2.5: Conceptual illustrations of 2d oxygen octahedra with a rotation, showing alternating off-centre oxygen atoms at the interface. Big dots represent the B-site atoms. Smaller dots represent O-site atoms. investigate the region between the experimental values. I this section, we carry out structural relaxations about th theoretical LSDA+U equilibrium volume rather than at th experimental lattice parameters reported previously. We ex plore biaxial strain states by fixing the in-plane lattice con stants at each strain value, and performing full structural op This off-centre behaviour of oxygen atoms at the surface could lead to a continuation of the rotations of octahedra in subsequently grown layers by the coupling of the octahedra to the atoms. Figure 2.6 shows this coupling of two materials octahedra with the interface in the middle. Figure 2.6(a) has no rotations and thus no off-centre oxygen-atoms. The symmetry of off-centre atoms in figure 2.6(b) causes the octahedra of the second layer of octahedra to mimic the exact same rotations. timizations of the internal atomic coordinates and c-ax parameter about the LSDA+U reference structure. In Fi 5 a, we show that over the strain range investigated th α FIG. 6. Color online The calculated band structure of LaNiO 3 on a LaAlO 3, b bulk LaNiO 3, and LaNiO 3 on c SrTiO 3 alon high-symmetry lines in the Brillouin given as 0,0,0 M 1/2,1/2,0 X 1/2,0,0 0,0,0. For each strain state, LaNiO 3 remai metallic. The role of strain on the electronic structure is to partially split the orbital degeneracy between the majority-spin d x 2 y 2 and d 3z 2 orbitals light blue points as indicated at the point. The Fermi level is denoted by the solid red line at 0 ev α (a) (b) Figure 2.6: Coupling of two materials oxygen octahedra across an interface (indicated by the dashed line). Big dots represent the B-site atoms. Smaller dots represent O-site atoms. Octahedral rotations by substrate induced strain are certainly dominant in thick epitaxial films. The behaviour of ultra-thin films has not been studied in much detail but could be expected to show some of the same behaviour. One could expect that the long range ordering caused by strain however is not as prevalent in thinner films. Thus the mechanism of template mediated 6

12 Chapter 2: General theoretical aspects octahedral rotations could play a role in the ultra-thin regime. 2.3 Metal-Insulator transitions The change in transport properties, from metal-like behaviour and conduction to insulator-like behaviour and no conduction, is what is known as a metal-insulator-transition (MIT). The transport properties are often dependent on such factors as oxygen pressure during growth, temperature during growth, and doping of the material. A proper definition of metallic and insulating behaviour is only possible at zero temperature T = 0. For an insulator the static electrical conductivity σ DC vanishes, whereas one finds a nonzero σ DC (T = 0) in a metal[14]. σ DC (T = 0) = 0 (2.2) Mott-transition The mechanism for the MIT in transition metal oxides can often be understood through the Motttransition[15]. This mechanism describes the process from metal to non-metal as a result from delocalised electrons. These delocalised electrons result when 1 2 K bt of energy is available per site so that the Boltzmann distribution is such that a significant portion of the electrons will be separated from their original site, becoming conduction electrons. This transition is sudden and occurs when N 1 3 a H Mott-Hubbard transition General band theories do not take into account the interactions of electrons with each other. Mott- Hubbard type transitions however are caused by this interaction. As in the Mott-mechanism, electrons are considered in their normal orbitals, before becoming conducting and hopping between lattice sites. However in this case the electrons are effected by the Coulomb repulsion of the electrons at other sites. The hamiltonian for a 1d chain of hydrogen atoms looks like: N H = t (c i,j c j,σ) + U n i n i (2.3) <i,j>,σ The first term is the regular tight-binding model of nearest-neighbour interaction with hopping integral t. The last term describes the coulomb interaction energy when electrons try to hop lattice site with interaction strength U. In order to distinguish Mott-Hubbard type insulator transitions the following definition is helpful: For a Mott-Hubbard insulator the electron-electron interaction leads to the formation of a gap in the spectrum for single charge excitations. The correlations force a quantum phase transition from a correlated metal to a paramagnetic Mott-Hubbard insulator, in which the local magnetic moments do not display long-range order.[14] The on-(lattice-)site Coulomb repulsion U splits the LDA bands into two sets of so-called Hubbard bands. One can envisage the lower Hubbard band as consisting of all states with one electron on every lattice site and the upper Hubbard band as those states where two electrons are on the same lattice site. Since it costs an energy U to have two electrons on the same lattice sites, the latter states are completely empty and the former completely filled with a gap of size U in-between. In the case of strongly correlated metals, as in some transition metal oxides, heavy quasiparticles are formed at the fermi energy with an effective mass or inverse weight of m m 0 = 1 Z >> 1[16]. These quasiparticles form what is also known as a coherent peak. i=1 7

13 Chapter 2: General theoretical aspects sents a single-hole state in the LHB, the of a Mott insulator has extra linewidth Figure 2.7: Schematic diagram of a half-filled Hubbard band system. The upper and lower Hubbard bands (LHB, UHB) are separated by the coulomb interaction energy U, while the coherent peak (CP) is visible in the centre.[17] FIG. 4. Schematic diagrams of three major configurations contributing to the core-level the impurity level at the core-hole site is attractive Coulomb potential Q. The upper Hubbard band (UHB), lower Hubbard band (LHB) and coherent peak or quasiparticle peak can all be seen in figure 2.7. Solving these systems requires dynamical mean field theory, as developed by Metzner and Vollhardt[18] SrTiO 3 /SrRuO The metal-insulator transition in SrRuO 3 comes about when going from 3 monolayers to 4 monolayers (ML) of SrRuO 3 on (001) oriented SrTiO 3 [2]. The 3 ML thick SrRuO 3 is not conducting, whereas 4 ML of SrRuO 3 and thicker returns to normal thin-film behaviour with regards to transport properties. The SrRuO 3 film also changes from an anti-ferromagnetic state to a ferromagnetic state when going from a 3 ML thick to a 4 ML film. A strong reduction in the magnetic moment is observed which for 3 ML and below lies in the plane of the film. Exchange-bias behaviour was also observed below the critical thickness which may point to induced anti-ferromagnetism in contact with ferromagnetic regions[2]. Knowing the structure of SrRuO 3, the most likely candidate for conduction comes from the overlap of the Ru 4d and O 2p orbitals. In the perovskite structure of the SrRuO 3, oxygen octahedra form Ru-O-Ru bonds with specific angles and lengths. Drawing the 4d and 2p orbitals in a simplified 2-dimensional picture in figure 2.8, it can be understood that the overlap of these orbitals depends on the rotations of these octahedra. When there is no rotation of the octahedra the 4d orbitals of the Ru atoms has no component parallel to the component of the O 2p orbital(figure 2.8(a)). With increasing rotation of the octahedra, the overlap of the orbitals also increases (figure 2.8(b)). α (a) (b) Figure 2.8: Ru 4d (red) and O 2p (green) orbitals with (a) no overlap and (b) overlap through octahedron rotation. 8

14 (Received 24 February 2004; published 16 September 2004) -ray photoemission spectra for the Mott-Hubbard systems are calculated by the n-field theory based on the exact diagonalization method. The spectra show a twoscreened and unscreened peaks. The screened peak is absent in a Mott insulator, but e main peak when the Chapter correlation 2: General strengththeoretical becomes weak aspects and the system turns metallic. pectral behavior is consistent with the experimental Ru 3d core-level spectra of various new mechanism ofrondinelli the core-level et. photoemission al. have shown[19] satellitethrough can be utilized Local to Spin reveal Density the Approximation + Hubbard phenomenon in various (LSDA+U) strongly calculations correlatedthat electron with systems, the inclusion especially of electron-electron in nanoscale interactions, bulk orthorhombic SrRuO 3 se-separated materials. is best described by a 0.6 ev on-site Hubbard term. However they are unable to reproduce the metal-insulator transition seen in experiments. They suggest that the behaviour srevlett may be dominated by extrinsic PACSfactors, numbers: such Fd, as surface h, disorder i and defects. They also note that due to the large spatial extent of the 4d orbitals in the ruthenates, correlation effects are anticipated to be less important as stronger hybridisation provides more effective screening and a reduced emission spectroscopy Hubbard (XPS) U. However as u ) thesepeaks last claims in various can be disputed, ruthenates as is [10 12] explained including the next section. ol to investigate the chemical RuO 2 where their origin has been the subject of a long d molecules [1]. Usually the debate [10 14]. The peak s, located at about 2 ev lower ergy depending on2.4 the chemial shift ) is utilized for this peak u, is absent in an insulator Y 2 Ru 2 O 7 and grows as Core-level binding XPS energy satellites than thatin of Mott-Hubbard the rather broad unscreened- type systems ures arising from the TheCoulomb charge-transfer the mechanism, system becomes as used to more calculate metallic. the spectrum The charge-transfer resulting from X-ray photoelectronhole spectroscopy cre- mechanism (XPS) of 3d transition as in 3dmetal TMCs compounds, is not adequate is not adequate to explain to explain the satellite electrons and a core be a fruitful source structures of inforgly correlated systems is much such smaller as than tween thathe of main a Pd 3d and spectrum, satellitewhile peaksa is larger much separation smaller is than expected due to the observed these experimentally. satellite structures, The energy since separation the energy between separation the main be- and satellite peaks on-metal compounds trend(tmcs) in chemical shifting[17]. that of a Pd 3d spectrum in PdO [15], while we expect e high T c and related Contrary cuprates to the assumptions larger separation of Rondinelli according et. al.[19], toelectron the chemical correlations trend around [3]. Ru 4d have been tures of Cu 2p core-level shown toxps play a role Quite in many recently, ruthenates Okada such [13] as tried Ca 2-x tosrexplain x RuO 4 and the pyrochlores[17]. two-peak Kim et. al. in confirmimg show the chargelating gap in the at parent the fermi-surface. com- effect [16] in addition to the charge-transfer mechanism. [17] that thestructure peaks in ruthenates in Sr 2 RuO are 4 the [12] result by the of screening nonlocal of the screening core-hole by quasiparticles nally classified as The Mott resulting insubasic parameterspeaks, such ascan thebe seen andinthe figure little2.9 change for different of thematerials. energy separation The screened in ruthenate and unscreened peaks are spectrum He could cannot, nowhowever, be described fit the bystrong a combination intensityofvariation unscreened and screened nergy and the charge-transfer indicated with theseries letterswith s andreasonable u respectively. values of parameters. th of information has been structures of charge-transfer s, a similar understanding of been lacking for the Motttter, we propose a new mechatellites for the Mott-Hubbard e Ru 3d spectra of various d by this model. rrelation effect has been con- MCs because 4d orbitals are 2, for example, is traditionally l [6]. However, recent studies tes such as Ca 2 x Sr x RuO 4 [7] various interesting properties effect among Ru 4d electrons. tra also show some hint of the Intensity (arb. units) Ru 3d RuO 2 Bi 2 Ru 2 O 7 Sr 2 RuO 4 SrRuO 3 CaRuO 3 Y 2 Ru 2 O 7 u Relative Energy (ev) Figure 2.9: FIG. Screen 1. and Ru unscreened 3d XPS spectra peaksofinruo several 2 (after ruthenates. [10]), Bi(Image 2 Ru 2 O 7 from [17]) of Ru 4d electrons. As shown (after [11]), Sr 2 RuO 4 (after [12]), SrRuO 3, CaRuO 3, Y 2 Ru 2 O 7 bit doublet, of which To describe splitting the XPS (after spectrum [11]). sfrom andfirst u denote principles, screened Kimand et.al. unscreened start frompeaks, a single-band Hubbard f roughly two model components, which is mapped respectively, ontofollowing a single impurity [11]. All the Anderson spectra model are aligned as prescribed by the in Dynamical ) and unscreened Mean (denoted Field Theory unscreened-peak (DMFT)[17]. Apositions core-holeofisru then 3d 5=2 added. to the converged model parameters in order to simulate the XPS process. The spectrum can now be seen as the contribution of three valence-electron configurations in the XPS final state: the d =04=93(12)=126404(4)$ The American Physical Society 1, d 2 L (Lower Hubbard band), and the d 2 C (Coherent band) configuration, where d denotes the impurity level. The impurity level in the form of the core-hole is subsequently screened by holes in the coherent peak and lower Hubbard bands in the d 2 L, and the d 2 C configurations. See figure 2.10 for a schematic diagram. When the correlation energy decreases, it becomes easier for valence electrons to move around, allowing them to more effectively screen the core-hole. As can be seen in figure 2.11, for a strongly correlated system with W/U = 0.75, the contribution of screening is not significant. In the limit of W/U 0 the separation of the screened and unscreened peak is given by Q U. Where Q is present because the potential is lowered by the attractive coulomb force resulting from the core-hole left by the XPS process. For a weakly correlated system with W/U = 1.5, both the d 2 L s 9

15 -8 - near the unscreened peak (the separation is Q U in the limit W=U! 0). Since the configuration d 2 L represents a single-hole state in the LHB, the core-level peak of a Mott insulator has extra linewidth broadening re- Chapter 2: General theoretical aspects Intensity (arb. units) x = 2.0 x = 1.6 x = 1.0 x = 0.4 x = 0.0 FIG. 4. Schematic diagrams of three major valence-electron configurations contributing to the core-level spectra. Note that FIG. 5. Compa the impurity level at the core-hole site is lowered by the attractive Coulomb potential Q. (a) Y 2 x Bi x Ru 2 O model calculatio VOLUME 93, NUMBER 12 PHYSICAL REVIEW LETT Figure 2.10: Schematic diagram of the three major valence-electron configurations. The impurity level at the core-hole site is lowered by the attractive coulomb potential Q. (Image from [17]) and the d 2 C configuration contribute to the unscreened peak. The separation of the unscreened and screened peak is now given by Q U/2. The weight of the coherent peak, or bandwidth W, dictates the intensity of the screened peak. Intensity (arb. units) (a) W/U = 0.75 total d 1 d 2 C d 2 L (b) W/U = 1.5 total Energy/U each core-level peak. For the four-site model, we should consider three energy levels in the conduction band representing a lower Hubbard band (LHB), a coherent peak, and an upper Hubbard band, along with the impurity level at the core-hole site in the presence of a core hole. Then there are three major valence-electron configurations in the XPS final state, i.e., d 1 ( d 1 L 2 C 1, for N s 4), d 2 C, and d 2 L, where d denotes the impurity level and C (L) denote the coherent peak (LHB). 10 These valence-electron configurations are shown schematically in Fig. 4. It can be easily noticed from the bar diagram at the bottom of d 1 d 2 C FIG. Figure : Core-level Hubbard band spectra valence configurations. in the four-site (Image from model [17]) for (a) W=U 0:75 and (b) 1.5. The calculated spectra are broadened by a Lorentzian of 0.25 full width at half maximum. The bar diagram at the bottom indicates the weights of different valence-electron configurations in each XPS final state. The three valence-electron configurations d 1, d 2 C, and d 2 L are shown schematically in Fig. 4 int he Fig. 3 caption. d 2 L flecting the LH value and the d 2 C is now av core hole, wh peak, whose e of the unscree the weight of To test the a systematic Y 2 x Bi x Ru 2 O 7 bandwidth-con theoretical cal valence-band shows the met We took Ru Y 2 x Bi x Ru 2 O 7 using MgK s oxygen partia nealing proces contamination binding energy C 1s peak. Ru 3d XPS obtained are inelastic backg contribution b shapes.we can of the screened tion increases. the transport p

16 Chapter 3 Hypothesis As previously mentioned, thin films of SrRuO 3 know a metal-insulator transition (MIT) when going from a thickness of 3 ML to 4 ML. A candidate mechanism for this thickness-dependent MIT in SrRuO 3 is the transition from the material s anti-ferromagnetic to its ferromagnetic phase. The energetically most favourable phase is determined by ratio between the Hubbard band bandwidth W and the electron interaction energy U. While the electron correlation model described in paragraph 2.4 could be a perfectly valid description of the physics of the MIT in SrRuO 3, it does not yet give concrete physical insight as to why a MIT occurs specifically when the thickness of the film changes from 3 ML to 4 ML of SrRuO 3 on SrTiO 3. X-ray Photoelectron Spectroscopy (XPS) can be used to measure the electron correlation energy U and bandwidth W. The screened peak of the Ru 3d spectrum as discussed in chapter 2 is a measure of the U and W parameters. Visible in the Ru 3d spectrum as an emergent extra peak (or shoulder in literature) of both the j=3/2 and j=1/2 peaks, this peak increases in intensity with the thickness of the SrRuO 3 film. Bulk SrRuO 3, which can be assumed to be completely relaxed as compared to strained thin-film growth, also displays this peak. The electron correlation energy as measured with XPS, is a representation of the electronic structure of ruthenium. Surrounded in the perovskite with oxygen atoms, ruthenium hybridises its 4d orbitals with oxygen s 2p orbitals and forms oxygen octahedra. These octahedra can be seen to form a network, spanning the crystal. Figure 2.8 showed part of such a structure. The oxygen octahedrals also carry rotations and are rotating in and out of phase with specific angles and symmetries throughout the crystal. In this way, it becomes supportable that these rotations play an important role in the electronic structure of this perovskite. The classical view of perovskites is that the octahedral rotations in them are only influenced by the dimensions of their unit cell. In thin-films this means that strain, caused by a mismatch between the lattice parameters of the substrate and the epitaxially grown material would play the largest role in the properties of SrRuO 3. However, there are hints that octahedral rotations in the thin-film are also influenced by the rotations of octahedra in the substrate. In this way, thin-films would at least for the first few monolayers couple to the rotations present in the substrate. Hints for this coupling have been shown in DFT calculations[20]. It could also be possible that the topmost few layers of the substrate undergo octahedral rotations imposed by the thin-film, relaxing the rotations present in the strained thin film into the substrate. Something which has been shown on La x Sr 1-x MnO 3 grown on SrTiO 3 [21]. The question then becomes if it would be possible to influence the octahedral rotations in the thin-film independently of the strain, in order to investigate what role octahedral rotations play. To influence the octahedral rotations in SrRuO 3, an epitaxial layer could be grown between the substrate and the SrRuO 3 thin film. Materials that could be used for this intermediate layer would have differing octahedral rotations, or no rotations at all. By growing epitaxially on the same substrate, the strain is kept constant. Many perovskites are well characterised in terms of octahedral rotations. As such, materials like LaAlO 3 or DyScO 3 could serve as materials with differing octahedral rotations. 11

17 Chapter 3: Hypothesis The effects of depositing an intermediate layer of material without octahedral rotations could result in the decoupling of the octahedral rotations of the substrate from the SrRuO 3 film. This could allow the SrRuO 3 to revert to a more bulk-like state with regards to its octahedra, causing the emergence of bulk-like properties which could be measured with XPS and transport measurements. A compound that could serve as a material without octahedral rotations is SrCuO 2. SrCuO 2 has a tetragonal phase (also called the infinite-layer phase) and an orthorhombic phase. The tetragonal phase has a layered structure consisting of alternating Sr and Cu-O planes. In the orthorhombic phase SrCuO 2 has chain structures of edge-sharing Cu-O plaquettes which are also weakly coupled to each other. Both phases could thus be able to decouple the octahedral rotations. Figure 3.1(a) shows the way in which orthorhombic SrCuO 2 would prevent the coupling of octahedral rotations. The chains of Cu-O plaquettes formed in the SrCuO 2 are separated by Sr atoms and are not coupled to each other. Figure 3.1(b) shows the way infinite-layer SrCuO 2 would use its layers (coming out of the plane of the paper at the second interface) to prevent octahedral coupling. (a) (b) Figure 3.1: Schematic representation the effect of intermediate layers of (a) orthorhombic (chainlike) SrCuO 2 and (b), tetragonal (infinite-layer) SrCuO 2. Big(small) black dots represent the B(O)-site atoms, while grey dots represent the Sr atoms in SrCuO 2. From bottom to top the materials would be SrTiO 3, SrCuO 2, and SrRuO 3 with the interfaces indicated with dashed lines. The ideal case for testing these octahedral rotations would then be to start with the STO (001) substrate, followed by a number of monolayers of intermediate material, followed by the desired number of monolayers of SrRuO 3. A capping of this film with the same intermediate material should then be performed, to preserve symmetry and to keep surface states and reconstructions in the SrRuO 3 to a minimum. 3.1 Goals To investigate the decoupling of octahedral rotations, the aim is to grow intermediate layers of materials with differing and no octahedral rotations between the SrTiO 3 substrate and SrRuO 3 film. The ultimate goal would be to induce conductivity at lower thicknesses of SrRuO 3. For the growth of an intermediate material without octahedral rotations, the growth of infinitelayer SrCuO 2 is preferred. Epitaxial films of SrCuO 2 will be grown and analysed using a variety of techniques including XPS and X-Ray Diffraction (XRD) to investigate the crystal structure. Other materials, which do have oxygen octahedra, can also be grown as intermediate layers. The effects of this on the SrRuO 3 thin film will also be investigated. The materials used will be LaAlO 3, DyScO 3. For an overview of the relevant stack configurations, see figure 3.2. The influence of octahedral rotations in the films will be investigated using XPS and transport measurements. The screened peak in the Ru 3d spectrum is a candidate for predicting transport properties on the basis of XPS. The manner in which the position and weight of the screened peak can predict the ability of the film to insulate or conduct will be investigated. 12

18 Chapter 3: Hypothesis The last goal will be to investigate the effects of capping has on the structural and electronic properties of thin films. LaAlO 3 / SrTiO 3 / DyScO 3 / SrCuO 2 SrRuO3 LaAlO 3 / SrTiO 3 / DyScO 3 / SrCuO 2 x ML x ML x ML SrTiO3 Figure 3.2: Proposed stack configurations, with the SrTiO 3 visible in green, and the SrRuO 3 sandwiched between layers of different materials. 13

19 Chapter 4 Experimental methods 4.1 In or ex situ? In the analysis and growth of materials there is an important distinction between in situ, latin for in position and ex situ, out of position. In situ implies the analysis of samples without them having left the environment where they were fabricated. This is important because by taking samples out of vacuum and by exposing them to air, contaminations on the surface such as carbon or water can change the properties of the sample. This is particularly of importance in surface sensitive techniques such as XPS, UPS or AFM. In XPS for example, the Ru 3d spectrum is perturbed by the presence of a carbon peak after exposure to air, making analysis of these spectra more difficult. The system used in this work for growing thin films is a fully closed in situ UHV system, containing both tools for growing materials and analysis. Analysis and characterisation can be achieved through XPS and AFM, while pulsed laser deposition (PLD) and vapour deposition allow for growth of materials. All these systems are interconnected and allow for samples to be grown and analysed without being exposed to contaminations. 4.2 Heteroepitaxy The term epitaxy comes from the Greek roots epi, meaning above, and taxis, meaning in ordered manner. It can be translated to arrange upon [22]. The term hetero can be translated from greek as meaning different. In the context of materials science, the term heteroepitaxy refers to deposition of thin films on a substrate in such a way that they are structurally congruent but of different materials. The layers deposited in such a fashion are called epitaxial layers. With today s tools such as pulsed laser deposition (PLD) or molecular beam epitaxy (MBE), it is possible to control the growth of these epitaxial layers to within atomic distances. Thin films with dimensions of only one atomic layer thick, i.e. monolayer (ML), can be created. This allows for a rich new area of physics with the possibility to create new devices and effects in the atomic limit. 4.3 Pulsed laser deposition Pulsed laser deposition (PLD) is a technique for thin-film growth where high energy laser pulses are focused on a target of the material that is to be deposited as a thin film. The laser pulses are focused into a vacuum chamber where they vaporise and ionise the target material, resulting in a plume of plasma being ejected perpendicular to the surface. The plasma expands and arrives at the substrate where the deposited species will start to nucleate and form a thin film. The substrate is typically heated to increase adsorption and promote the growth of smooth crystalline films. The vacuum chamber can also be filled with background gases, such as oxygen, which can be used in the oxygenation of oxides. Figure 4.1 shows a typical PLD setup. 14

20 Chapter 4: Experimental methods Figure 4.1: A typical PLD setup, with the laser being focused before the window and striking the target material. The plasma plume can be seen impinges on the substrate. (Image from [23]) The species and stoichiometry of the plasma can be controlled by the fluence (J/cm 2 ). This fluence is dictated by the mask used in the beam path, as well as the final spot size at the target which is dictated by the strength of the lens in front of the window. Most materials have a specific fluence associated with them that has been empirically determined to lead to high quality films Growth modes During deposition, it is possible to have several growth modes. Because all substrates have a miscut, atomic steps will be present at the surface. In step-flow growth, the step plateaus are small enough that the species will diffuse to the step edges before they have a chance to nucleate. The growth can be seen as steps travelling perpendicular to their normal. In layer by layer growth, the atoms nucleate at several sites on the step plateaus creating islands. These islands grow until they coalesce and complete one monolayer. In the 3-dimensional (3D) growth mode, islands are formed on top of the islands formed in layer by layer growth. Thus not creating smooth film, but forming higher islands which do not coalesce Reflection high energy electron diffraction Reflection high energy electron diffraction (RHEED) is a tool used to track the structural properties of crystalline solids during growth. In RHEED, an electron beam is created by generating electrons in a filament and accelerating them to an energy of several tens of KeV. The RHEED gun used during deposition in this work accelerates electrons to 30 KeV. The electron beam then diffracts of the surface of the sample and strikes a phosphor screen. The diffraction pattern on this phosphor screen is recorded using a camera. RHEED is an excellent tool to monitor growth during deposition. The electrons do not interfere with growth and the setup is such that no part of the apparatus impedes the plasma. In the case of a perfect 2-dimensional structure, the RHEED diffraction pattern will consist of spots corresponding to the rods in reciprocal space. In this case, Kikuchi lines, which are caused by inelastic scattering in reciprocal space, can also be seen. These Kikuchi lines are often used by crystallographers as roads in orientation space because they connect diffraction spots. With the introduction of imperfections the diffraction spots can become streaks or elongated spots due to the broadening of the rods in reciprocal space. To align the diffraction spots, the sample can be rotated about the surface normal. 15

21 Chapter 4: Experimental methods Perhaps the most useful application of RHEED is in the determination of the amount of monolayers grown during deposition. In the case of layer by layer growth, the formation of nucleation sites and subsequent island growth will cause a roughening of the surface. This roughening can be seen in the intensity of the diffractions spots because of the scattering of electrons on the roughness. As the layer increases in roughness the intensity of the RHEED signal will decrease until the layer is half filled. The RHEED signal will recover to a maximum intensity as the remaining part of the layer is filled. In this way, one oscillation would equal the growth of one monolayer. The intensity oscillations visible in layer by layer growth are not visible during 3D growth. With increasing imperfections, such as height differences in 3D growth, the electron beam will pass through the 3D structures and cause secondary spots. These 3D spots are distinguishable from 2D spots because they will not move when the angle of the incident electrons changes. Step flow growth mode can be recognised by the sharp recovery of RHEED intensity after each laser pulse[24]. 4.4 Materials properties of LaAlO 3, DyScO 3, SrRuO 3, and SrTiO 3 The following table 4.1 recounts the materials properties of the perovskites used in this work. The material properties of SrCuO 2 will be discussed in chapter 5. Material Lattice constants (l.c.) (pseudo-)cubic l.c. rotation system a b c a glazer Sr 2+ Ru 4+ O a a c + (α = β = 7.56, γ = ) [25]. 3 Dy 3+ Sc 3+ O a b + a [26] 3 La 3+ Al 3+ O a a a [27] 3 Sr 2+ Ti 4+ O no rotations 3 Table 4.1: Valence, lattice parameters and octahedral rotation systems of used perovskite materials in bulk. 4.5 Atomic Force Microscopy Atomic force microscopy (AFM) is a high-resolution imaging technique based on forces experienced by a mechanical probe. The mechanical probe often consists of a cantilever with a sharp tip. A laser is aligned to the end of the cantilever and its reflection is measured to determine the deflection of the tip. The movement of the tip is controlled by piezoelectric actuators. This method allows high-resolution images of the atomic surface down to distances of fractions of a nanometer. Several modes of operation for AFM are available: contact mode, non-contact mode, tapping mode. All AFM measurements in this work were made in tapping mode. Tapping mode works by oscillating the cantilever close to its eigenfrequency and measuring the change in vibration frequency due to van der Waals forces as the tip scans the surface. 4.6 X-ray Photoelectron Spectroscopy Measurements and Setup X-ray Photoelectron Spectroscopy or XPS is a surface-sensitive measurement technique which can probe the electronic and chemical environments of elements in materials. X-ray photons with specific energy are generated in a source and injected into the material. These photons interact with the specific elements and knock out (photo)electrons close to the nucleus, leaving behind a so-called core-hole. Measuring the energy of the photo-electrons now allows the probing of the 16

22 #!"# $%&"'&$()*# +#,-.# /$)'0%12)0)%# '1"/&/0/# 13#,4%+5#/16%')*#)()'0%1"#)")%75#+"+(58)%#+"9# )()'0%1"#9)0)'01%:#;+'<#$+%0#=&((#>)#>%&)3(5#&"0%196')9#<)%):# Chapter 4: Experimental methods # binding energy of these electron to their specific elements, revealing the electronic and chemical environments. The equation for the electron binding energy then reads:!"/#01(,2#&.)(3+&# E binding = E photon (E kinetic + Φ) (4.1) # where E photon is the energy of the x-ray photon, E kinetic the kinetic energy of the photoelectron, and?<)#21/0#&2$1%0+"0#'1"/&9)%+0&1"#&"#'<11/&"7#,4%+5#/16%')#31%#,-.#&/#)")%75#%)/1(60&1":#@7# Φ the work function of the material. A #BCDEF*G#)HI#+"9#J(#A The basic apparatus consists #BCKLG*G#)HI+%)#0<)#0=1#3%)M6)"0(5# of an X-ray source, a monochromator, 6/)9# sample, /16%')/:# andn1%# electron )33&'&)"0# detector. First, x-rays are generated by accelerating electrons up to 15 kev from a filament into an anode $%196'0&1"#13#,4%+5/#>5#)()'0%1"#>12>+%92)"0*#+"#+'')()%+0&"7#O1(0+7)#13#CE#AH#&/#+$$(&)9#&"# material (figure 4.2). The filament is usually made of either magnesium (Mg) or Aluminium +"9# The J(# x-rays /16%')/:# by/16%')/# bremsstrahlung +%)# 9)/&7")9# from this =&0<# anode 96+(#/16%')/*# will be of type 1")# +"19)# Kα; resulting 3+')# >)&"7# from the transitions 2p 3/2 1s and from 2p 1/2 1s. The characteristic energies (widths) of the principal '1+0)9#=&0<#@7#+"9#0<)#10<)%#=&0<#J(*#+/#/<1="#&"#N&7:CD:## lines are (0.70) ev for Mg Kα and (0.85) ev for Al Kα. [28] ############################################ # #######################4-5)(+#/6#7),819%.:+#01(,2#');+# Figure 4.2: Schematic of an x-ray source. P)0=))"#0<)#+"19)#+"9#0<)#+"+(5/&/#'<+2>)%*#+#0<&"#J(#1%#P)#=&"91=#&/#&"0)%$1/)9#01#$%10)'0# The spread of energy besides the principles lines of these sources is rather broad due to being generated from bremsstrahlung. Therefore, it is useful to monochromate Kα radiation by firstorder diffraction using!"#$%&'(")*+),-.&/) an α-quartz crystal..&/)0*1*"#.*%&'*. All XPS experiments carried out in this work use a 0<)#/+2$():# monochromated x-ray source.?<)#(&")#=&90<#13#,4%+5/#3%12#@7#+"9#j(#/16%')/#&/#+%16"9#q:rsq:l#)h:#j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igure 4.3: Schematic of a monochromator refracting x-rays from a quartz-crystal. The electron detector consists of an input lens, a concentric hemispherical analyser (CHA) (figure 4.4), and multiple channel electron multipliers also called channeltrons. The input lens decelerates the electrons from the sample and focuses them onto a slit at the entrance of the hemisphere. The CHA has an internal electric field which can be varied. Electrons in the field are 17

23 Chapter 4: Experimental methods deflected proportional to their kinetic energy. By varying the electric field, electrons with specific kinetic energies can be selected to reach the exit aperture and the detector. &'%()*+%,-.# /0%# )1# )2%*$# 3*44%$%5)# 61$7#405()*15&"#)2%$%#*&#8#(15)8()#'1)%5)*8,# -V 2 #*5#9*:.;.#<1#)2%#%5%$:-#14#)2%#'21)1%,%()$15&#3%)%()%3#=-#)2%#%,%()$15#%5%$:-#8 -V.# A2%# 61$7# 405()*15# 14# )2%# 858,->%$# 1 *&# (15&)85)# 853# )2%# *5(*3%5)# '21)15# 28&# Entrance width W s r 1 r 0 r2 Exit width W F -.#<1#)2%#=*53*5:#%5%$:-#(85#=%#%8&*,-#3%30(%3#8&B# % # Figure 4.4: Hemispherical analyser with incoming electron at S with potentials at the inner and outer shell at V 2 and V 1.! D "! C By connecting the analyser and the sample, the fermi levels are levelled and it is possible to draw the following diagram (figure 4.5) regarding the energies. There is a contact potential between sample and analyser but by levelling the fermi energy the following relations hold: " &' ############################################################################################## &' ############################################################################################################# % # E b = E ph (E k + Φ s ) (4.2) E b = E ph (E k + Φ a) (4.3) where E b is the binding energy, E ph the energy of the photon, Φ s,a the work function of the material and analyser, E k the kinetic energy of the electron after exiting the sample, and and E the kinetic k energy of the electron as seen by the analyser. ############# #!"#$%&'(')*&%#+',-.&/&'01'&/",,"0*'2*3'3&4&-4"0*'01'5.040&6&-4%0*,7' Figure 4.5: Energy relations between emission and detection of photoelectrons. The binding energy can now be described by knowing the work function of the analyser and by the measured energy in the analyser. Something that would not be possible without levelling the fermi energy of the sample and the detector. In practice the work-function of the detector is &4&%/"*24"0*'01'4.&'0;"324"0*',424&,' known. In the case of insulating samples some electrons will not reach the surface of the material or cannot overcome the work-function. Because the sample is insulating, these electrons cannot flow to ground and this leaves them in or on the material, causing a build-up of charge. The effect on the spectrum will be a shift upward in binding energy since the kinetic energy measured in the detector will be lower. *I'1$)85)# 8'',*(8)*15# 14# JK<# *&# )1# 3*&)*5:0*&2# )2%# 1L*38)*15# &)8)%&# 14# )2%# % 18 3*5:#)1#)2%*$#3*44%$%5)#=*53*5:#%5%$:*%&.#A2%#(285:%#14#)2%#'21)1%I*&&*15#,*5%& %5)#(2%I*(8,#%5+*$15I%5)&#*&#(8,,%3#M(2%I*(8,#&2*4)N.#H5%#I13%,#)1#3%&($*=%#)2%#' #

24 Chapter 4: Experimental methods Surface sensitivity While XPS can probe the electronic and chemical environments of solids, it can only do so up to a small depth from where electrons can escape. A naive approach to this is the Beer-Lambert law: I z = I 0 e z/λsin(θ) (4.4) where the measured intensity I z is described by the intensity from the surface atoms I 0 ; attenuated by a term dependent on depth z, take-off angle θ (with respect to the surface plane), and the inelastic mean free path (IMFP) λ. In reality not only inelastic collisions determine the measured intensity, but also processes involving elastic scattering. This discovery lead to significantly greater values for λ, now better understood as an attenuation length. Also, it is known that electron attenuation in solids is not exponential and are dependent on experimental geometry. [28] Assuming exponential attenuation the sampling depth, where 95% of the signal is detected can be defined as three times the escape depth. Therefore: d s = 3λ AL sin(θ) (4.5) where d s is the sampling depth, λ AL the attenuation length, and θ the take-off angle. As the energy of the x-ray photon increases, so does the kinetic energy of the electrons. Because the electrons are higher in energy they will have a higher probability of escaping the solid, increasing the attenuation length and sampling depth. This also explains why ultraviolet photoelectron spectroscopy (UPS) is an extremely surface sensitive technique. Its photons have an energy of 21.2 ev. By decreasing the take-off angle, less of the bulk is seen and the technique becomes more surface sensitive. Sampling depths at kinetic energies for up to 1200 ev can be as high as 1.7 nm, 6.9 nm, and 9.7 nm for 10, 45, and 90 respectively. [28] The X-ray source used in all experiments in this work, produces monochromated photons with an energy of ev. Fitting Before fitting the line shapes of XPS peaks the background signal must be determined. The background signal in XPS is caused by the elastic scattering of electrons in the material. Because these electrons lose energy, they appear to originate from higher binding energies. Background subtraction was performed using a manual multipoint Shirley background approximation. Shirley backgrounds are a computationally inexpensive iterative procedure defined by: Em ax F n (E) = J(E) k n F n 1 (E )de (4.6) where F(E) is the background signal at E. J(E) is the measured spectrum, and k n an arbitrary constant. The underlaying assumption for the Shirley algorithm is that the background caused by inelastic scattering processes at an energy, E, is proportional to the number of electron with kinetic energies higher than E[28]. More thorough treatments of the background can be found in the Tougaard universal cross-section method[29]. The Heisenberg principle dictates that the interaction time will cause an uncertainty in the energy. Due to this so-called life-time broadening, XPS peaks will not appear as single lines, but instead be broadened by a factor of Γτ = evs[28]. Where Γ is the uncertainty in the energy, and τ the interaction time. So for a 1 fs interaction time, the line broadening will be about 0.1 ev. The resulting distribution of energy will be gaussian. The line of an XPS peak can be approximated by the convolution of a lorentzian and gaussian peak, known as a Voigt function. The fitting parameters for a peak are then area, lorentzian peak E 19

25 Chapter 4: Experimental methods width, gaussian peak width, and position. Constraints are also added to account for effects such as the intensity ratio between peaks of different spin-orbit coupling. An educated guess is then made from the spectrum by eye or theoretical predictions about multiplet contributions, to determine how many peaks and at what approximate position they should be located. This is important for the initial parameters. A local minimum, as with many fitting algorithms, can be reached quickly with the wrong initial parameters. With these fitting parameters a non-linear regression fit is then performed on the XPS line shape after shirley background subtraction. Non-linear fits mitigate some of the problems with standard least-squares fitting in that the likelihood of ending up in a local fitting minimum is lower Simulations The method used to simulate spectra in this work is based on charge transfer multiplet theory. This theory overcomes some of the limitations of single electron excitation models in the descriptions of rare earth (RE) systems and transition metal (TM) compounds. The CTM4XAS simulation program is used for all calculations. CTM4XAS uses a modified version of the atomic structure code originally developed by Cowan[30], and can generate spectra for X-ray Absorption Spectroscopy (XAS), X-ray Magnetic Circular Dichroism (XMCD), and XPS. The steps in the simulation program are as follows. First, the specific one-electron radial are calculated using a Hartree-Fock method. Then, for each possible value of the total angular momentum J, eigenvalues (energy levels) and eigenvectors, and other characteristics such as radiation spectra, and radiative transition probabilities are calculated. Then, the irreducible representations of the configurations are used to branch the configurations into the correct crystal field symmetry. If charge transfer is calculated, dipole transitions for XPS transition with and without ligand holes are calculated. These are then mixed essentially using a monopole transition in the initial and final states. An overview of such a calculation for the Ru 4+ 3d xps process is seen in figure 4.6. Appendix A contains a configuration file that can be used for such calculations. dipole 3d 10 4d 4 p 0 3d 9 4d 4 p 1 mixing t mixing t` 3d 10 4d 5 p 0 L dipole 3d 9 4d 5 p 1 L Figure 4.6: The configurations and their relations for the calculations of charge transfer in Ru 3d. Crystal Field In the case of crystal structures, ligands are ions that bind to a central atom to form a coordination complex. The five d-orbitals of the transition metal are initially degenerate. By introducing a ligand atom, the orbitals that are closest to the ligand atom will increase their energy through coulomb repulsion. The orbitals farther away will experience a decrease in energy. The difference in energy levels of the orbitals is called the crystal field splitting 0. The electronic structure of these complexes depends on the symmetries with which the ligand atoms bond. The crystal field caused by these ligands leads to different crystal field splittings of different orbitals, as seen in figure 4.7. The three lower energy-orbitals in octahedral symmetry are referred to as t 2g and the two higher energy-orbitals as e g. The calculation of the crystal field in the simulations is performed by an application of group theory. By determining the irreducible representations of the symmetry groups, branching relations can be found such as the ones in table 4.2. From the branching table it can be seen that for a d- orbital, the SO 3 representation splits into E+T 2 for an octahedral (Oh) symmetry. Other branching tables for going to lower symmetry, such as from Oh symmetry to square planar (D4h) symmetry, can also be determined. The branching relations are then applied to the atomic transition matrices elements. 20

26 Chapter 4: Experimental methods (a) (b) (c) (d) (e) (f) Figure 4.7: Three different ligand configurations (octahedral, square planar, and tetrahedral) and their resulting crystal field splitting. (Images from [31]) SO 3 O h (Mulliken) S 0 A 1 P 1 T 1 D 2 E + T 2 F 3 A 2 + T 1 + T 2 G 4 A 1 + E + T 1 + T2 Table 4.2: Branching rules for symmetry elements from SO 3 to O h symmetry. 21

27 Chapter 4: Experimental methods Figure 4.8: Cu 2p spectra of several Cu compounds. (Image from [32]) Charge transfer Charge transfer is the mechanism by which an electron is donated from the ligand to the transition metal ion. This donation leaves a hole in the ligand, which is denoted by L. For a Ru 4+ oxidation state, the electron configuration would be 3d 10 4d 4. When a ligand donates an electron the 4d shell will have 5 electrons, giving a 3d 10 4d 5 L configuration. For calculating the spectrum, the simulation now has to calculate two dipole transitions; one with charge transfer, and one without. For the XPS process, the dipole transition involves the excitation of an electron from a core-level to vacuum. In the simulation, this process is represented by the transfer of one d-shell electron to the continuum. Due to selection rules, an electron ejected from a d-shell will have character p. This is indicated in figure 4.6 as p 0 before the dipole transition, and as p 1 after. After the initial and final states of both configurations are calculated, they are mixed using a monopole transition calculation. Staying with the Ru 4+ oxidation state, the simulation parameter of charge transfer energy determines the mixing of the 3d 10 4d 5 L and 3d 10 4d 4 state. Higher energies will see lower contributions of charge-transfer effects in the spectra. Another simulation parameter is the Hubbard repulsion energy U dd and core-hole potential U pd. These last parameters are essentially in competition and only the difference U dd U pd would matter. The other important parameters are the hopping or hybridisation parameters which for octahedral symmetry imply that V(e g ) = T(b 1 ) = T(a 1 ) and V(t 2g ) = V(b 2 ) = V(e)[33]. 4.7 Transport measurements To measure the transport properties, such as resistivity, a Quantum Design physical property measurement system (PPMS) was used. To measure the electrical properties of flat samples of arbitrary shape if the contacts are sufficiently small and located in the circumference of the sample a measurement setup was devised by van der Pauw[34]. The sample is placed on a measurement puck that fits inside the PPMS. Electrical contacts are wire bonded from the sample to the measurement puck using silver bonds. For a square sample with electrical contacts in each corner, labeled 1 through 4, current is supplied first through one edge of the sample (corner 1 to 2) and voltage is measured on the opposite edge (corner 3 to 4). Then, current is supplied through the edge perpendicular to the first edge (corner 1 to 3) and voltage is measured on the opposite side (corner 2 to 4). Figure 4.9 shows this setup. 22

28 Chapter 4: Experimental methods Figure 4.9: A van der Pauw setup with electrical contacts in each corner. This way, the resistances are and R 12,34 = V 34 I 12 (4.7) R 13,24 = V 43 I 12 (4.8) The sheet resistance is given by the following relation e π R 12,34 Rs + e π R 13,24 Rs = 1 (4.9) where R s is the sheet resistance. This equation is not solvable exactly, and is usually solved numerically using iterative approaches such as the Newton-Raphson algorithm. To calculate the resistivity ρ of a thin film of known thickness d the usual resistivity relation ρ = R A l where A is the cross-sectional area and l the length of the piece of material reduces to: ρ = R s d (4.10) 4.8 XRD X-ray diffraction (XRD) was used to obtain information on the crystal structures of samples. XRD is a non-destructive technique and allows for the quick characterisation of crystal structure. A limitation of this technique in common laboratory settings is the limited intensity of the source. For films too thin, the diffracted intensity will not be high enough to get clear diffraction peaks. To increase the signal it is possible to go to synchrotron facilities, which get high-intensity x-ray beams resulting from bremsstrahlung in the. XRD relies on the constructive interference when reflecting off of planes of atoms in the crystal structures under a certain angle. Using Bragg s law (4.11) the angle of the principal diffraction peaks can be calculated. nλ = 2dsin(θ) (4.11) Where λ is the wavelength of the x-ray source, d the spacing between the planes in the atomic lattice, and θ the angle between the incident ray and the scattering plane. Measurements in this work were made using a Philip X pert thin film diffractometer with a Cu Kα1 source with a photon wavelength of Å. A Bruker D8 was also used. The Bruker D8 has a monochromator so alternate spectral lines, such as Cu Kα2, are insignificant in diffraction experiments. It also has the ability to make reciprocal space maps. 4.9 SrTiO 3 substrate treatment All films grown in this work are deposited on SrTiO 3 substrates. SrTiO 3 has a cubic perovskite unit cell and with a lattice parameter of Å. The lattice match of SrTiO 3 with other oxides is one of the main reasons this material is widely used. The lattice match of SrTiO 3 with SrRuO 3 is 23

29 Chapter 4: Experimental methods such that SrRuO 3 grows with a (compressive) strain of only -0.5%. For the growth of SrRuO 3 on SrTiO 3 it is important that the surface is clean and has unit cell steps. This can be accomplished by solvents such as acetone or isopropyl-alcohol and annealing. This does not however take care of something called mixed-termination. Because SrTiO 3 is built up out of alternating layers of SrO and TiO 2 it is possible, after cleaving of the sample, that the topmost layer of the substrate is a mixture of these terminations. With annealing, the SrO terminations tend to move to the step edges and form a morphology that is recognisable by sharp edges. These patterns can easily be observed through AFM techniques. This mixed termination is a problem for the growth of SrRuO 3, since the diffusivity of this material is different for different terminations, leading often to island growth. To eliminate mixed termination, the weak acid HF can be used[35]. By exposing the substrate to water, the SrO termination is hydrogenated to form Sr(OH) 2, which can be etched away using HF. The resulting surface is virtually free from SrO termination and after annealing the step edges are long and well defined. First, the as-received substrates are set in a teflon holder for easy handling and dipping in various chemicals in the treatment steps. The substrates are then cleaned by emerging them in solvents, specifically ethanol and acetone, and placed in an ultrasonic bath for approximately 15 minutes per solvent. To hydrogenate the SrO, the substrates are placed in water in an ultrasonic bath for 30 minutes. The substrates are then taken out of the water and immediately placed inside a buffered HF solution (NH 4 F (87.5%), HF (12.5%)). The substrates in the solution are then placed in the ultrasonic bath for 30 seconds after which they are immediately taken out and dipped in three beakers of water to dissolve the acid and finally placed in ethanol. The substrates are then blown dry with nitrogen and placed in an oven for annealing at 950 C for 90 minutes with a constant flow of oxygen at 150 ml/min. Some substrates have been known to diffuse underlying SrO to the surface during annealing. To resolve this, substrates can be re-etched using the same procedure without the anneal step and by placing the substrates in HF for only 10 seconds. An example of the effect of this treatment is seen in figure (a) (b) Figure 4.10: AFM image 4.10(a) of the result of HF treatment and annealing showing steps and 4.10(b) a line profile across the steps showing only unit cell steps. 24

30 Chapter 5 SrCuO 2 This chapter will describe the properties and structures of SrCuO 2 when grown as a thin film on SrTiO 3 (100). SrCuO 2, a parent compound of high-temperature superconductors, has been grown in thin films at low pressure before, using a variety of techniques including laser molecular beam epitaxy, pulsed laser deposition[36 40], magnetron sputtering and metal-organic chemical vapour deposition[41]. Compared to the films used later on in this report, the SrCuO 2 films described here are relatively thick (30-40 ML). This is because the characterisation techniques used are not able to give structural information on films much thinner. 5.1 Structure There are three possible structures of SrCuO 2 that need to be taken into account: orthorhombic SrCuO 2, tetragonal (or infinite-layer) SrCuO 2, and orthorhombic Sr 2 CuO 3. While SrCuO 2 is expected to be formed in most cases, unforeseen stoichiometry considerations could lead to the formation of Sr 2 CuO 3. Figure 5.1 shows the atomic configurations in their respective unit cells, accompanied by a multiple unit cell view which shows in which way the Cu-O plaquettes are formed. Orthorhombic SrCuO 2, which is the stable bulk phase, contains zigzagging plaquettes of edgesharing Cu-O chains parallel to the c-axis, stacked along the b-direction (lattice parameters a= Å, b= Å, c= Å). The electrons in the compound are likely to move along the Cu-O chains, and the coupling between neighbouring chains is weak[42]. The plaquettes are configured as shown in figure 5.2b. Tetragonal, or infinite-layer, SrCuO 2 consists of planes of corner-shared Cu-O plaquettes perpendicular to the c-axis and alternating with layers of Sr (figure 5.2d). Its lattice parameters are a=b=3.926 Å, and c=3.432 Å. Orthorhombic Sr 2 CuO 3 (lattice parameters a=12.69 Å, b= Å, c=3.494 Å) contains linear chains of corner-sharing Cu-O plaquettes shown in figure 5.2a. In summary, these three phases form structures with three distinguishing features: a zigzagging edge-shared chain, a corner-shared infinite-layer plane, and corner-shared linear chain. By growing thin films epitaxially on SrTiO 3 it is possible to control orientation and phase. It is expected from strain principles that due to the mismatch of lattice parameters of tetragonal SrCuO 2 with SrTiO 3 (-0.5%), that the growth of the tetragonal phase is most preferred. Orthorhombic SrCuO 2, with a strain of 4% in the a/c plane will likely grow with its b-axis perpendicular to the surface. Because of growth on the charge neutral layers of SrTiO 3, competition between the polar nature of tetragonal SrCuO 2 (Sr 2+ /CuO 2 2 ) and orthorhombic SrCuO 2 (SrO 0 /CuO 0 ) could play an important role in the structure that is grown for the first few monolayers. It is expected that at first, the polar nature of the tetragonal phase will lose to the orthorhombic phase. An atomic reconstruction is the most energetically favourable way to resolve the polarity. At higher thicknesses polarity is 25

31 Chapter 5: SrCuO2 (a) (b) (c) (d) (e) (f) Figure 5.1: The unit cells of (a) orthorhombic SrCuO2, (c) tetragonal SrCuO2 and (e) Sr2 CuO3 respectively. The corresponding Cu-O plaquettes can be seen in (b),(d), and (f), showing the different structures formed by the different phases. Sr, Cu and O-atoms are denoted by green, blue, and red orbs respectively. Material Lattice constants (Å) a b c SrCuO2 (Orthorhombic) SrCuO2 (Tetragonal) Sr2 CuO3 (Orthorhombic) Table 5.1: Lattice parameters of Srx CuOx 26

32 Chapter 5: SrCuO 2 57 BRIEF REPORTS 139 FIG. 1. Sketch of the Cu-O networks in the different cuprates examined. Cu atoms (), O atoms (). a The linear chain in Sr 2 CuO 3, b the zigzag chain in SrCuO 2, c the chain of edgeshared plaquettes in Li 2 CuO 2, d the Cu A O 2 plane in Sr 2 CuO 2 Cl 2 and the Cu 3 O 4 plane in Ba 2 Cu 3 O 4 Cl 2 containing an extra Cu B site (), and e the crenellated chains of Ba 3 Cu 2 O 4 Cl 2. Figure 5.2: Different plaquettes and crystal structures present in (a) Sr 2 CuO 3, (b) orthorhombic SrCuO 2, and (d) tetragonal SrCuO 2. (Image from [43]) O-Bi-O bridging units 14 and no coherent Cu-O network is realized. Ba 3 Cu 2 O 4 Cl 2 contains chains of edge-shared CuO 4 plaquettes, which are so arranged as to give a crenellated Cu-O network as depicted in Fig Sr 2 CuO 2 Cl 2 is considered to be the paradigm 2D spin- 1 2 Heisenberg antiferromagnet. 16 Ba 2 Cu 3 O 4 Cl 2 has a similar structure to that of Sr 2 CuO 2 Cl 2, 17 although its Cu-O network is composed of a regular CuO 2 plane denoted by Cu A in Fig. 1 with an additional Cu site Cu B ) which is connected to the regular Cu A O 2 plane via a 90 Cu A -O-Cu B configuration. Consequently, the Cu atoms in Ba 2 Cu 3 O 4 Cl 2 show two antiferromagnetic phase transitions at widely different temperatures and two different ZRS dispersion functions connected with these different Cu subsystems. 3 The experiments were carried out using a Perkin-Elmer photoemission system equipped with a monochromatic Al K source giving a resolution of about 0.4 ev. Since the samples are insulating, corrections for charging effects were undertaken, resulting in an estimated accuracy of the given absolute energy values of 0.3 ev. The measurements were performed at room temperature and the crystals were cleaved in situ under ultrahigh-vacuum conditions. The O 1s spectra of all samples show negligible emission at binding energies greater than 531 ev, which indicates clean samples. We found that the shape of the Cu 2p spectra depends critically on the cleanliness of the oxygen signal. In Fig. 2 we show the Cu 2p 3/2 spectra of the cuprates. An integral background has been subtracted and the spectra are normalized to the leading peak. The area of the Ba 2 Cu 3 O 4 Cl 2 spectrum is multiplied by a factor of 1.5 compared to that of Sr 2 CuO 2 Cl 2, assuming it to be proportional to the number of Cu atoms in the Cu-O plane for reasons to be shown later. The spectral features for all crystals studied are summarized in Table I. In the following, we will describe the salient features of the Cu 2p main lines and then discuss the role of the electronic states near the chemical potential in the screening processes responsible for the observed structures. In contrast to all previous XPS studies known to us, the Cu 2p main line spectra show either a rich fine structure or are fairly narrow symmetric lines. The position of the lowest-binding-energy feature denoted by A is the same within our experimental accuracy for all the systems studied. The spectrum of screened significantly and will result in the material being better able to form smooth films since atomic reconstruction does not distort the structure[44]. DFT calculations have shown that the transition from orthorhombic to the tetragonal infinite-layer phase should occur, starting from a film thickness of around 5 ML[44]. This means that below this 5 ML threshold the orthorhombic phase of SrCuO 2 is expected, provided that the stoichiometry of the ablated target is correct. This idea is further supported by TEM measurements made on another sample (STO/10 ML LaAlO 3 /1 ML SrCuO 2 /2 ML SrTiO 3 ), also grown at this university. In these measurements, the d-spacing was measured to be 3.9 Å. This would be around 1/4 of the unit cell of orthorhombic SrCuO Growth FIG. 2. Cu 2p 3/2 photoemission spectra of the single-crystalline cuprates. The 1D chain systems a SrCuO 2, b Sr 2 CuO 3, c Li 2 CuO 2, d Ba 3 Cu 2 O 4 Cl 2, e 0D Bi 2 CuO 4, and planar f Sr 2 CuO 2 Cl 2 and g Ba 2 Cu 3 O 4 Cl 2. Also shown is the difference spectrum gf. Bi 2 CuO 4 is very similar to a previously published one. 18 A recent low-resolution XPS study also gave a comparable spectrum for polycrystalline Sr 2 CuO 3, 11 although, in the present case using high-resolution XPS, a shoulder denoted C accompanying the leading main line feature A is clearly resolved. The spectrum for the zigzag chain SrCuO 2 closely resembles that of Sr 2 CuO 3, whereby the intensity of C is larger in the former. In addition, referring to Table I, the satellite to main line intensity ratio I s /I m is larger for SrCuO 2 than for Sr 2 CuO 3, indicating a larger for the zigzag chain in the language of single-site models. In contrast, for Bi 2 CuO 4, Li 2 CuO 2, and Ba 3 Cu 2 O 4 Cl 2 a comparatively narrow and symmetric main line B is observed at lower binding energy than feature C in the other 1D cuprates. At the position of feature A either no or a small spectral intensity is observed. The spectrum of Sr 2 CuO 2 Cl 2 has been discussed in more detail elsewhere. 12 The broad main line is composed of two features denoted A and B, and a prominent shoulder C appears at higher binding energies. Compared to Sr 2 CuO 2 Cl 2, the feature B in Ba 2 Cu 3 O 4 Cl 2 is much more pronounced. For both 2D systems, I s /I m is nearly the same whereas C is positioned at different binding energies. Significantly, the position of feature B is very close to that of the main lines of Bi 2 CuO 4, Li 2 CuO 2, and Ba 3 Cu 2 O 4 Cl 2. To understand these features we will focus first on the 0D and 1D systems and relate the main line features to the differently coupled CuO 4 units in Bi 2 CuO 4, Li 2 CuO 2, Ba 3 Cu 2 O 4 Cl 2, SrCuO 2, and Sr 2 CuO 3, assuming the simplified geometry of the Cu-O networks displayed in Fig. 1. In the case of Bi 2 CuO 4, the interpretation of the main line spectrum is straightforward: Since no coherent planar or linear Single terminated SrTiO 3 (100) substrates were prepared using the buffered HF treatment described in chapter 4, and subsequently measured with AFM to gauge substrate smoothness, single/double termination and step size. The growth temperature was aimed to be 600 C and measured with a heat-spy. This temperature is the result of literature study which showed growth of infinitelayer from around 500 C[38] to 650 C[41]. A sintered SrCuO 2.5 target was pre-ablated and subsequently used for deposition with a laser energy of 55 mj per pulse resulting in a fluence of 2.11 J/cm 2 under a partial oxygen pressure of mbar. The laser repetition rate was 1 Hz. During deposition sample roughness was monitored using in situ RHEED to monitor the number of monolayers grown. After deposition the samples were cooled down at 100 mbar oxygen by switching off the heater and allowing it to return to room temperature. As the RHEED intensity oscillations of the main spot show in figure 5.3, a maximum is achieved after about 8 laser pulses, indicating that the surface roughness reaches a minimum and one monolayer is grown. After about two monolayers the growth mode changes and clear oscillations are no longer visible. The combination of spikes at 1 Hz in the signal, and a constant intensity suggest pulse-recovery. This could indicate steady-state or step-flow like growth. The growth-rate was taken to be 8 pulses/ml, and 40 ML of SrCuO 2 were grown on STO (001). Figure 5.4 shows the RHEED patterns after the growth of each different material. The spacing between the spots in-plane perpendicular to the beam give a measure for the spacing of the reciprocal rods, and thus for the in-place lattice parameters. The spacing of the spots was found to be 34, 33.9, and 76 pixels for the SrTiO 3 substrate, the SrCuO 2 film, and after the SrRuO 3 film (rotated 26.5 azimuthally), respectively. The same pixel spacing for the SrTiO 3 substrate and SrCuO 2 film indicates that in at least one direction, the SrCuO 2 is fully strained. The ratio of the SrRuO 3 (rotated 26.5 ) and the SrCuO 2 film spacing was found to be 0.446, which is consistent with the ratio of rotation sin(26.5 ) = The SrCuO 2 film is thus fully in-plane strained after deposition, at the deposition temperature. AFM images before growth show a substrate that is single terminated, with only single unit cell and smooth long running steps, also indicating single termination(figure 5.5). After deposition the steps are difficult to distinguish. The film was measured with AFM to have a maximum height difference of 1 nm and an rms roughness of 0.14 nm. 27

33 Chapter 5: SrCuO 2 Intensity (a.u.) growth of one monolayer manual RHEED intensity change laser turned off Time (s) Figure 5.3: RHEED intensity over time for the growth of 40 ML of SrCuO 2. RHEED intensity over a larger timescale. Insets show the (a) (b) (c) Figure 5.4: RHEED images (a) of the SrTiO 3 substrate before deposition, (b) after the growth of 40 ML of SrCuO 2, and (c) after the subsequent growth of 3 ML of SrRuO 3 rotated 26.5 around the azimuthal angle. 28

34 Chapter 5: SrCuO 2 (a) (b) (c) Figure 5.5: AFM image showing (a) the SrTiO 3 substrate before deposition and (b) the 40 ML thick film after deposition with its (c) line profile of the surface. 29

35 Chapter 5: SrCuO XRD Previous diffraction patterns measured on SrCuO 2 grown by molecular vapour deposition(mvcd) by Chang[41] can be seen in figure 5.6. The orthorhombic phase clearly shows the (0 6 0) and (0 10 0) peaks with the latter being lower in intensity, whereas the tetragonal phase shows the (001) and (002) diffractions peaks with the (002) being higher in intensity. Figure 5.6: FIG. XRD1. spectrum X-ray /2 for tetragonal, scan for films orthorhombic depositedsrcuo at various 2, andsubstrate orthorhombic temperatures. a Substrate temperature 800 C, orthorhombic SrCuO 2 with a grown with different C Sr 2 CuO 3 [41] FIG. 3. The x-ray / To determine b-axis the orientation. phase of SrCuO b Substrate temperature 700 C, orthorhombic 2, X-ray diffraction (XRD) was used. The out-of-plane lattice of the (Sr 1x Ca x )CuO parameters Sr 2 of CuO different 3 with phases a-axis orientation. in the grownc material Substrate were temperature measured and 650 analysed. C, tetragonal SrCuO on a 40 2 A θ/2θ scan creasing Ca concentra was performed ML with thick c-axis sample orientation. of SrCuO 2 on a SrTiO 3 substrate, which also had a capping creasing Ca concentra of 3 ML of SrRuO 3. c-axis lattice constant Figure 5.7 shows that the sample contains the SrTiO 3 (001), (002), and (003) substrate peaks. Secondary spectral peaks from a x-ray source are indicated in figure 5.7. These correspond to the Kβ spectral lines of the substrate peaks of SrTiO 3. Furthermore, there are peaks corresponding to the aluminium plate used to support the sample during the measurement. mize fluoride formation, the partial pressure of water is increased by increasing the water bubbler temperature to 23 C. The temperature range for obtaining the infinite-layer compound is known to be very narrow. 16 This results primarily from the competition between the phase stability of the tetragonal phase and the amorphous phase. Both are more stable at lower temperatures. In the present case, there is also a competition between fluoride phase and oxide phase formation arising from the use of fluorinated metal organic precursors. X-ray scans were performed to determine in-plane orientation. A diffractometer equipped with a four circle goniometer was used to measure the off-axis reflection of the 101 family of planes. The film was rotated in plane during the diffraction measurements along its fourfold symmetry axis, which is parallel to the30plane normal of the substrate The remaining out-of-plane diffraction peaks correspond to the (001) and (002) peaks for the tetragonal phase of SrCuO 2 with a d-spacing of d=3.44 ± 0.01 Å and are located at 25.8 and The spectrum also contains the (0 6 0), (0 8 0), (0 10 0), and (0 12 0) peaks of orthorhombic SrCuO 2 at 33.0, 44.77, 56.5, and respectively, corresponding to a 1 st order d-spacing of d=16.28 ± 0.05 which is only 0.3% mismatch from the theoretical c lattice parameter of orthorhombic SrCuO 2. This could indicate that the orthorhombic phase in the material is no longer fully strained, which could mean that the top part of the material consists of unstrained orthorhombic SrCuO 2, with the strained tetragonal phase below. The grown film could thus be seen as a layered composition of tetragonal SrCuO 2, starting at the surface extending a number of monolayers. With increasing thickness the strain is not able to support the unstable tetragonal phase, possibly resulting in unstrained orthorhombic SrCuO 2. surface. The geom Fig. 2. For epitax spaced 90 apart width at half-max 0.3. Diffraction dimensional epita The depositio was next examin (Sr 1x Ca x )CuO 2 Ca compositions cium substitution higher 2 values From the diffracti data reported by determined. The c sition of the depos be substituted up

36 Intensity (arb. units) Tetr. SrCuO 2 (001) Orth. SrCuO 2 (060) Aluminium plate K-β Orth. SrCuO 2 (080) Tetr. SrCuO 2 (002) Orth. SrCuO2 (0 10 0) K-β Aluminium plate Chapter 5: SrCuO 2 2θ h k l * * Table 5.2: Diffraction peak parameters for tetragonal SrCuO 2 (prepared at 60kbar at 1050 C for 30 minutes, from the ICDD crystallographic database). Entries indicated with * were calculated from the table s (002) entry with d= STO (001) STO (002) STO (003) Orth. SrCuO2 (0 12 0) Θ Figure 5.7: XRD spectrum of 40 ML of SrRuO 3 on SrTiO 3, capped with 3 ML of SrRuO Photoelectron spectroscopy The photoelectron spectra of Cu-compounds are relatively well understood, and give a good idea of the band structures and chemical environment of Cu. The 2p doublet of Cu carries the most information about the oxidation state of copper[45]. An example of a spectrum is shown in figure 5.8. The difference between the valence states (copper metal no valency, Cu 1+ 2 O) and Cu 2+ O is immediately apparent. The 2p 3/2 and 2p 1/2 peaks in Cu metal can be linked to the 2p 5 3d 10 final state. For CuO, with a 2+ oxidation state of copper, the oxidation changes the electron configuration of Cu from 3d 10 4s 1 to 3d 9 and two extra so-called satellite peaks appear. The satellite peaks are due to this 3d 9 ground state of copper. Hybridisation of the bands between copper and oxygen causes oxygen to donate an electron to copper through the charge transfer mechanism, leaving a hole in in the ligand. The total configuration is written as 3d 10 L, where the hole is indicated by L. So in contrast to the Cu metal, the 3d 10 peak is then no longer caused by the ground state, but by charge transfer. The 3d 10 L is a lot broader for CuO than for a Cu metal. This is due to electron delocalisation and the screening of the valence hole[46]. While XPS allows easy distinction between valence states, the difference in spectral shape for one valence state among different compounds is very insightful. It is instructive to observe some previously measured spectra of several cuprates in literature. Figure 5.9 shows these spectra for a number of cuprates, including the orthorhombic phases of SrCuO 2 and Sr 2 CuO 3. The differences are subtle but mainly reside in the shoulder of the 3d 10 L. Figure 5.9 also shows (f)sr 2 CuO 2 Cl 2 and (g)ba 2 Cu 3 O 4 Cl 2, which are compounds that also contain Cu-O planes, as in infinite-layer SrCuO 2. The Cu 2p spectrum of the infinite-layer structure Sr 0.9 La 0.1 CuO 2 and La 1.85 Sr 0.15 CuO 4 has also 31

37 Chapter 5: SrCuO2 Chapter 5. Tetragonal CuO Figure 5.5: Photoemission spectra of Cu, Cu2 O, and CuO (tenorite). For both Cu as well Figure Photoemission spectra of Cu, CuO (tenorite). as Cu25.8: O only 2 peaks are visible, which are Cu related to the 2p35/2 3d 10 and (Image the 2p15/2from 3d 10 [45]) final 2 O, and states, which have almost the same binding energy in both materials. The CuO spectrum is more complex and has 4 peaks due to the possibility of oxygen donating an electron to the copper 3d band. The final states for the CuO peaks are labeled in the figure. The largest peak is called the main peak and the smaller peak related to the ground state is called the satellite peak in the text. The peaks of the CuO are shifted to higher binding energy by more than 1 ev. They are also much broader than the peaks of the other two materials due to delocalized screening and multiplet splitting, which is explained in the text. BRIEF REPORTS 139 a lot broader for CuO than it is for Cu or Cu2 O, which has been explained by Van Veenendaal and Sawatsky. 34,35 By performing calculations they show that there is a competition between screening electrons coming from local ligand atoms and electrons coming from ligand atoms of neighboring Cu metal ions to screen the charge on the Cu. This causes an extra contribution to the main peak, which shows up as a shoulder at higher binding energy in the case of CuO. The multiplet splitting in the 2p5 3d9 satellite peak is more subtle and is the subject of many papers. 36, Magnetic properties The electrons in a Mott insulator cannot move, but their spins can still fluctuate. The spins usually arrange themselves in an antiparallel fashion, due to virtual charge the Cu-O networks in the different cuprates 38 fluctuations, which leads to long-range antiferromagnetic ordering. The temperature ( ), O atoms ( ). a The linear chain in at which this ordering takes place is the Ne el Temperature (TN ). zag chain in SrCuO2, c the chain of edgethe2cl Ne el temperatures for transition metal monoxide compounds with a rock salt Li2CuO2, d the Cu A O 2 plane in Sr2CuO 2 in Ba2Cu3O4Cl2 containing an extra Cu B siteare graphically represented in figure 5.6 on the facing page. The y-axis structure ellated chains of Ba3Cu2O4Cl2. is scaled logarithmically and the Ne el temperature increases exponentially with the transition element number. The Ne el temperature for CuO does not follow this trend, nits14 and no coherent Cu-O but network is remember, CuO has a monoclinic structure, which could be the cause of the reduced Cl2 contains chains of edge-shared CuO 4 transition temperature. If CuO could be forced into a rock salt structure its Ne el are so arranged as to give a crenellated epicted in Fig onsidered to be the paradigm842d spin- 21 FIG. 2. Cu 2 p 3/2 photoemission spectra of the single-crystalline romagnet.16 Ba2Cu3O4Cl2 has a similar 5.9: The XPS1D spectra several cuprates. chain of systems a cuprates. SrCuO2, b (Image Sr2CuOfrom Sr2CuO2Cl2,17 although its Cu-O network Figure 3, c [43]) Li2CuO2, d Ba3Cu2O4Cl2, e 0D Bi2CuO4, and planar f regular CuO2 plane denoted by Cu A in Sr2CuO2Cl2 and g Ba2Cu3O4Cl2. Also shown is the difference itional Cu site Cu B ) which is connected spectrum g f. O 2 plane via a 90 Cu A -O-Cu B configuly, the Cu atoms in Ba2Cu3O4Cl2 show Bi2CuO4 is very similar to a previously published one.18 A etic phase transitions at widely different recent low-resolution XPS study also gave a comparable two different ZRS dispersion functions spectrum for polycrystalline Sr2CuO3,11 although, in the se different Cu subsystems.3 32 present case using high-resolution XPS, a shoulder denoted s were carried out using a Perkin-Elmer C accompanying the leading main line feature A is clearly em equipped with a monochromatic Al resolved. The spectrum for the zigzag chain SrCuO2 closely a resolution of about 0.4 ev. Since the resembles that of Sr CuO, whereby the intensity of C is ing, corrections for charging effects were

38 Chapter 5: SrCuO 2 been measured before as can be seen in figure 5.10[47]. Notable is the suppression of the 3d 10 L shoulder and shift of the 3d 9 satellite structure R P Vasquez et al Figure 5. Cu 2p 3/2 core levels measured from chemically-etched surfaces of (a) La 1.85Sr 0.15CuO 4 epitaxial film, (b) La 2CuO 4 epitaxial film, (c) Sr 0.9La 0.1CuO 2 polycrystalline pellet, and (d) Nd 1.85Ce 0.15CuO 4 epitaxial film. Figure 5.10: Cu 2p3/2 core levels measured from chemically-etched surfaces of (a) La 1.85 Sr 0.15 CuO 4 epitaxial film, (b) La 2 CuO 4 epitaxial film, (c) Sr 0.9 La 0.1 CuO 2 polycrystalline pellet, and (d) Nd 1.85 Ce 0.15 CuO 4 epitaxial film. (Image from [47]) Measurements The Cu 2p 3/2 core level spectra are presented in figure 5. The multiplet at the higher binding energy, referred to as a satellite peak in the literature, corresponds to states of predominant 2p 5 3d 9 L character, while the main peak at the lower binding energy corresponds to states of predominant 2p 5 3d 10 L character resulting from ligand-to-metal (O 2p Cu 3d) charge transfer [28 30]. Different final states can have sufficient energy separation to yield resolvable features in the main peak, which is most evident in the spectrum measured from Sr 0.9 La 0.1 CuO 2 in figure 5(c). Such features have previously been resolved in Cu 2p spectra measured from SrCuO 2 [10], in which the Cu O layers consist of zig zag chains, and from Sr 2 CuO 3 [10, 31], in which the Cu O layers consist of linear chains. Calculations for Sr 2 CuO 3 based on a threeband Hubbard model have also been shown to be consistent with the measured Cu 2p spectrum [31], with the main peak features corresponding to final states with different screening of the photoexcited core hole. The Cu 2p spectra are commonly analysed within a simple configuration interaction model utilizing a two-band Hamiltonian [3, 29, 30, 32, 33]. Within this model, the O 2p Cu 3d charge transfer energy,theo2p Cu3dhybridizationstrengthT, and the on-site Coulomb interaction between Cu 2p and Cu 3d holes U, are related to the experimentally-determined energy separation between the poorly-screened satellite and well-screened main peak (E s E m ) and to the ratio of the intensities of the satellite and main peak (I s /I m ). The experimental values of I s /I m and E s E m obtained for La 1.85 Sr 0.15 CuO 4 are consistent with the earlier results for hole-doped cuprates [3]. The electron-doped materials both exhibit higher values of E s E m and lower values of I s /I m than values measured from hole-doped cuprates. One might have expected that the Cu 2p spectra of Nd 1.85 Ce 0.15 CuO 4 δ and Sr 0.9 La 0.1 CuO 2 would be most similar. Both materials are electron doped, both have similar Cu O bond lengths [1, 34] (1.97 Å) which are significantly greater than those in La 2 x Sr x CuO 4 [35] (1.89 Å) and the To investigate the SrCuO 2 -SrRuO 3 interface, two types of samples were fabricated. One type where approximately 40 ML of SrCuO 2 were deposited on STO (001), capped with 3 to 4 ML of SrRuO 3. The second type did not have this capping. This allows for the inspection of any surface reconstruction of SrCuO 2 that might occur in this process, while also allowing to peer at the chemical environments of the material. Because we are mostly interested in the Cu-O plaquettes in the material, and because oxygen spectral shapes are often easily influenced by other extraneous factors such as water-vapour and other small disturbances, the focus is mainly on the Cu 2p (j=3/2 and j=1/2) core-hole peaks. XPS measurements were performed on several samples (figure 5.11) with the detector positioned at an 80 angle to the surface plane. The X-ray photons were produced by an Al Kα source and subsequently monochromated with a quartz monochromator. The spectrum shows several features, starting with a main peak at 933 ev, corresponding to 3d 10 L. It has a shoulder feature at 935 ev that is reduced in intensity as the sample is capped with SrRuO 3. The smaller peaks from 940 ev to 945 ev are called the satellite peaks in literature and correspond to the 2p 5 3d 9 final state. Due to capping the satellite structure changes clearly. The satellite structure appears to shift slightly to lower binding energy with the peak highest in energy at 945 ev reducing drastically in intensity. The feature at 953 ev is the 3d 10 L (j=1/2) peak and mirrors the effects and features we observe in the main 3d 10 L (j=3/2) peak. There is also a 3d 9 (j=1/2) peak at around 962 ev but is not shown in the figure. Because XPS only penetrates roughly 20 Å(approximately 8 ML) of SrCuO 2 into the film, there is a strong indication that surface reconstruction plays a role in the uncapped versus the capped film since XPS spectra (shown in figure 5.11) differ markedly in the shoulder of the 3d 10 L peak and in the highest binding energy satellite peak. This reconstruction might be abated by the capping of the SrCuO 2 film with SrRuO 3. There is also an effect visible from different oxygen pressures at cooldown. A higher oxygen pressure at cooldown resulted in an increase of the height of the 3d 10 L shoulder peak Simulations Simulations were performed to investigate the influence of charge-transfer parameters and how this would match experiments. 33

39 Chapter 5: SrCuO SCO 30 ML uncapped, 100 mbar cooldown SCO 30 ML uncapped, mbar cooldown SCO 22 ML, capped 4 ML SRO SCO 40 ML, capped 3 ML SRO 2p 5 3d 10 L (j=3/2) Intensity (a.u.) p 5 3d 10 L (j=1/2) 2p 5 3d 9 (j=3/2) Binding energy (ev) Figure 5.11: Cu 2p XPS measurement of SrRuO 3 with an intermediate layer of SrCuO 2 on SrTiO 3. The two top measurements are samples capped with SrRuO 3, the bottom two are left without a capping layer. I m /I s = 0.33, 0.25, 0.39, 0.40 from top to bottom. Multiplet calculations with charge-transfer were performed in square planar (D4h) and octahedral (Oh) symmetry. For D4h symmetry the irreducible representations are Γ = b 1g, a 1g, b 2g, and e g. The hybridisation terms V(Γ) influence the spectral shape the most for the 2p 3/2 and 2p 1/2 components, with most of the emphasis on the 2p 5 3d 9 (j=3/2) shape. For the relation V(Γ) in D4h symmetry we assume from geometric considerations that V(a 1g ) = V(b 1g )/ 3, V(b 2g ) = V(e g )/ 2, and V(b 2g ) = V(b 1g )/2[32]. For the higher octahedral (Oh) symmetry the hybridisation strengths V(e g ) and V(T 2g ) are the varied parameters. The other parameters, such as the charge transfer = 2 ev, the electron-hole correlation U dd = 6eV and the core-hole potential U dc = 8 ev. The resulting spectral lines were broadened with a 0.4 ev wide lorentzian-gaussian profile. By varying V(b 1g ) from 1.5 ev to 3 ev and calculating the resulting remainder of the hybridisation terms a spectral shape is obtained that resembles the measured Cu 2p spectra. The full hybridisation parameters can be found in table 5.3. Hybridisation Parameter Configuration 1 Configuration 2 Configuration 3 V(b 1 ) 1.5 ev 2.25 ev 3 ev V(a 1 ) 0.86 ev 1.3 ev 1.7 ev V(b 2 ) 0.75 ev 1.13 ev 1.5 ev V(e g ) 1.1 ev 1.6 ev 2.1 ev Table 5.3: Cu 2p configuration parameters for D4h symmetry. Fixed parameters were = 2 ev, U pp = 6 ev, and U pd = 8 ev Figures 5.12 and 5.13 show the simulations done for Cu 2+ for Oh and D4h symmetries, respectively. The D4h symmetry shows an asymmetric satellite peak for all hybridisation strengths. This asymmetry is only present in Oh symmetries for low hybridisation energies. Figures 5.14 and 5.15 show the simulations done for Cu 3+ for Oh and D4h symmetries, respectively. For Oh symmetry, which branches only to e g and T 2g, the only parameters that are varied are the hybridisation in these branches. The shoulder peak at 935 ev seen in experiments could not be reproduced in our simulations without bringing copper to a 3+ valence state. The separation between the main peak and the lowest lying satellite peak is about 8 ev in experiments. This matches the separation seen in the simulations for both valencies and a hybridisation strength of about 3 ev. 34

40 Chapter 5: SrCuO V(b1g) = 3 ev V(b1g) = 2.25 ev V(b1g) = 1.5 ev Intensity (a.u.) Binding energy (ev) 20 Figure 5.12: Cu 2p XPS spectrum calculation for Cu 2+. V(b1g) is varied from 1.5 tot 3 ev and follows the relation for D4h symmetry with respect to the other hybridisation terms. Intensity (a.u.) V(Eg) = 1 ev, V(T2g) = 1 ev V(Eg) = 2 ev, V(T2g) = 2 ev V(Eg) = 3 ev, V(T2g) = 3 ev V(Eg) = 4 ev, V(T2g) = 4 ev Binding energy (ev) 20 Figure 5.13: Cu 2p XPS spectrum calculation for Cu 2+. V(b1g) is varied from 1.5 tot 3 ev and follows the relation for Oh symmetry with respect to the other hybridisation terms. 35

41 Chapter 5: SrCuO V(b1g) = 3 ev V(b1g) = 2.25 ev Intensity (a.u.) Binding energy (ev) 20 Figure 5.14: Cu 2p XPS spectrum calculation for Cu 3+. V(b1g) is varied from 1.5 tot 3 ev and follows the relation for D4h symmetry with respect to the other hybridisation terms V(Eg) = 3 ev, V(T2g) = 2 ev V(Eg) = 1.5 ev, V(T2g) = 1 ev Intensity (a.u.) Binding energy (ev) 20 Figure 5.15: Cu 2p XPS spectrum calculation for Cu 3+. V(Eg) is varied from 3 ev tot 1.5 ev and follows the relations for Oh symmetry. 36

42 Chapter 5: SrCuO 2 To compare simulations with experiments, one of the features that can be compared is the ratio in area between the satellite and main peak. This is denoted by I m /I s. The shift of the satellite structure (also present in the j=1/2 peak at 962 ev) to lower energies in figure 5.11, accompanied by an increase in the I m /I s ratio could signify, as shown in simulations, that the hybridisation strength becomes weaker. The simulations couldn t accurately predict the ratio between the main peak and satellite peak intensity I m /I s, which, in experiments was around 0.3 to 0.4 and in simulations around 0.17 (at the hybridisation strength that produced the right separation between satellite and main peak). 5.5 Discussion To take the literature into account, one observation can be made when viewing previously measured spectra. The edge-sharing Cu-O plaquettes tend to have a higher intensity of the 3d 10 L shoulder peak than the corner-sharing plaquettes specifically, Sr 2 CuO 3 compared to SrCuO 2 (figure 5.9). The intensity of the 3d 10 L shoulder peak decreases when SrRuO 3 is used to cap SrCuO 2. In the context of edge and corner-sharing, this would imply that SrCuO 2 films capped with SrRuO 3 display more corner-sharing type behaviour. The shoulder of the 3d 10 L peak in figure 5.11 is associated with a localised character of the valence hole redistribution[46] as shown in figure 5.16(b). So as the shoulder decreases with the capping of SrCuO 2 with SrRuO 3, the percentage of Ru sites that has a localised character decreases as well. With this reasoning, a chain-like structure is less common since there would be more degrees of freedom with a lower shoulder peak. A more delocalised character of the material is present, implying a more infinite-layer type structure. When capped with SrRuO 3, the 3d 10 L main peak shifts to a lower separation distance with respect to the 3d 9 quintet. This also happens in Ba 2 Cu 3 O 4 Cl 2 when compared to SrCuO 2. Ba 2 Cu 3 O 4 Cl 2 shares some aspects of the Cu-O plaquette planes of infinite-layer SrCuO 2 as can be seen in figure 5.2(d). 80 EUROPHYSICS LETTERS Intensity I (!) (arb. units) (a) Bi 2 CuO 4 (b) Sr 2 CuO 3 (a) (b) (c) Fig Binding energy! (ev) Fig. 2 Figure 5.16: Plaquettes with valence hole delocalisation. The core-hole is indicated by a cross. Fig. 1. Comparison of experimental data (dots), taken from ref. [1], and theoretical results for (a) Bi 2CuO Large 4,and(b)Sr (medium, 2CuO 3. small) The line dots spectra symbolise have beenaconvoluted large (medium, with a Gaussian small) function density of of the valence hole orig- ev (dashed width 1.8inally located lines), atand the 0.2central ev (solidplaquette. lines). For details Thesee final the text. states (a), (b), and (c) correspond to the lines at Fig. 2. approximately Valence hole delocalization 933 ev, in935 finalev, states and of the photoemission ev in figure process 5.9. for(image an infinite from CuO 3 [46]) chain. The core-hole Cu site is denoted by a cross. Large (medium, small) dots symbolize a large (medium, small) density of the valence hole originally located at the central plaquette. The final states (a), (b), Another and (c) correspond possible to the confirmation lines at approximately of the transition 933 ev, 935 ev, to infinite-layer and ev in fig. SrCuO 1(b). 2 at the SrCuO 2 -SrRuO 3 interface is the spectrum shown in figure Comparing these spectra found in literature to the experimental data in this work, some of the same trends can be discovered. The decrease of the upper estimate 3d 10 L shoulder, of The as well exchange as the termdecrease in eq. (1) inleads the to highest an asymmetric bindingsplitting energyof peak of the 3d 9 satellite can the satellite structure into a quintet (approximately at binding energy 941 ev) and a triplet be seen. (944 ev) final state. The narrow main line (934 ev) is associated with a local final state in which the Although valence hole themoves shoulder from the ofcu thesite main onto3d the 10 surrounding L peak cano most sites. Due likely to the be small associated with the previously Cu occupation mentioned in thislocalised final state the valence exchange hole splitting redistribution, in the main line there is small. have also been reports of mixed valence As shown in fig. 1(b) the result for an infinite CuO 3 chain with =2.7eV is in good contributions in (La agreement with the experimental S r)cuo with regards to the ligand-hole peak[48]. Indeed, XPS simulations as spectrum of Sr 2 CuO 3. The calculated ratio I s /I m =0.37 is equal topreviously the experimental discussed value [1]. produce The satellite this same emission feature. has a two-peak But thestructure ionisation similar energy required for the Cu 3+ to that of a single plaquette although the degeneracy of the quintet and triplet splitting is now lifted due to additional delocalization. The exchange splitting of the main line, on the other hand, is small. While the main line consists of several features 37 it is dominated by two contributions. There is a low-energy peak at 933 ev and a shoulder at higher energy (935 ev). The delocalization properties of the corresponding processes are shown in fig. 2. The low-energy peak (933 ev) is mainly due to a delocalization of the valence hole to its neighbouring plaquettes, fig. 2(a), which may be interpreted as a Zhang-Rice singlet formation.

43 Chapter 5: SrCuO 2 state is around 1500 kj mol 1 higher than that for Cu 2+ [49]. Therefore this valence state is not expected. The effect visible in the height of the shoulder 3d 10 L peak with different oxygen pressures during cooldown is likely caused by a decrease in oxygen vacancies in the Cu-O plaquettes. The XRD pattern of the SrCuO 2 capped with SrRuO 3 shows both tetragonal and orthorhombic phases. The out-of-plane lattice parameter for the orthorhombic phase indicates that it is not fully strained, while the out-of-plane lattice parameter for the tetragonal phase indicates that it is fully strained (with SrTiO 3 in-plane lattice parameters). While the RHEED pattern taken immediately after growth indicates that the SrCuO 2 film is fully strained, the film might relax when cooling down to room temperature. The picture therefore emerges that at first the tetragonal phase of SrCuO 2 is grown, but that the top part of the film relaxes to an orthorhombic phase when cooled down. Stabilisation of the film could be occurring by capping the film with SrRuO 3, causing less of the SrCuO 2 to relax to an orthorhombic phase. This is supported by XRD patterns of uncapped films (not shown) that show less intense peaks for the tetragonal phase. This is also consistent with XPS spectra, which show that with capping more of the infinite-layer type spectrum is seen. 5.6 Conclusions Charge-transfer simulations could not reproduce the behaviour of the 3d 10 L shoulder, most likely due to the fact that valence redistribution over multiple plaquettes is not a capability since all calculations are performed in the ligand model. XRD and XPS indicate that the measured films show both tetragonal and orthorhombic phases of SrCuO 2. The SrCuO 2 film likely consists first of around 5 ML of orthorhombic SrCuO 2 at the bottom of the film, followed by tetragonal SrCuO 2, followed by the orthorhombic phase. Strain might no longer be able to accommodate the tetragonal phase close to the surface, causing the formation of the orthorhombic phase. This could be confirmed using reciprocal space mapping. A possible factor in this could be the growth temperature. Consequently capping the SrCuO 2 with SrRuO 3 seems to, by inference from XPS and XRD, lead to a stabilisation of the SrCuO 2 leading to less formation of the orthorhombic phase. There was no evidence of the existence of Sr 2 CuO 3 in the films. The types of films grown will likely contain infinite-layer planes and would be well suited in this regard to test the decoupling of oxygen octahedra across an interface. The 5 ML thickness would be the minimum of thickness expected to grow the tetragonal phase. For very thin films (5-20 ML) the formation of the orthorhombic phase might be suppressed, with strain being able to support the desired tetragonal phase for the entire film. Even if the orthorhombic phase of SrCuO 2 were present in thinner films, it would still be suitable for decoupling octahedra because of the weak coupling between chains. 38

44 Chapter 6 The metal-insulator transition in SrRuO 3 Chapter 5 described the growth and structural characteristics of SrCuO 2. With that analysis complete, SrCuO 2 can be used in the investigation of the MIT in SrRuO 3 (SRO). For the analysis of the thickness dependence of the MIT in SrRuO 3 there are now four materials available to use as intermediate and capping layers: LaAlO 3 (LAO), SrTiO 3 (STO), DyScO 3 (DSO), and SrCuO 2 (SCO). By keeping the same substrate over all samples, strain is constant, and changes in properties due to octahedral rotations should become prevalent. This chapter will analyse the XPS spectra in connection to the transport properties and different heteroepitaxial structures with different intermediate materials. 6.1 Heteroepitaxial structures To investigate any effects of the intermediate materials and their octahedral rotations on the MIT in SrRuO 3, several types of heteroepitaxial structure were grown. The first type of structure is formed by growing an intermediate layer before growing the SrRuO 3 film. In the case of the use x monolayers of DyScO 3 as intermediate material and the use of SrTiO 3 as a substrate, this can be written as STO/x ML DSO/4 ML SRO. The second type of structure is a sandwich type structure, where the previous structure is capped with the same intermediate material. This would be notated as STO/x ML DSO/4 ML SRO/y ML DSO. Both heteroepitaxial structures are shown in figure 6.1. LaAlO 3 / SrTiO 3 / DyScO 3 / SrCuO 2 x ML SrRuO3 x ML SrRuO3 x ML LaAlO 3 / SrTiO 3 / DyScO 3 / SrCuO 2 x ML LaAlO 3 / SrTiO 3 / DyScO 3 / SrCuO 2 x ML SrTiO3 SrTiO3 (a) (b) Figure 6.1: Heteroepitaxial structures of the (a) uncapped type and (b) sandwich type. 39

45 Chapter 6: The metal-insulator transition in SrRuO Growth Single terminated SrTiO 3 (100) substrates were prepared using the buffered HF treatment described in chapter 4. For the growth of the perovskite intermediate/capping layer materials, different parameters were used. Growth temperatures were measured using an infrared heat spy and set to be 650 C. SrCuO 2 was grown at nearly the same conditions described in chapter 5, with only a difference in growth temperature: now around 650 C. in situ RHEED was used to measure the number of monolayers grown in all cases. After deposition, samples were allowed to cool down to room temperature at 100 mbar oxygen pressure without temperature control SrRuO 3 SrRuO 3 was grown at a fluence of 2.1 J/cm 2, with a partial pressure of mbar consisting of 50% argon and 50% oxygen. The partial oxygen pressure was previously optimised for the magnetic properties of SrRuO 3 by another member of this group. SrRuO 3 in the (100) direction is made up out of layers of SrO and RuO 2. When depositing SrRuO 3 ontio 2 terminated SrTiO 3, the first monolayer of SrRuO 3 deposited will continue the SrTiO 3 termination by stacking SrO first, and RuO 2 second. However, the volatile RuO 2 will evaporate, leaving the more stable SrO as the topmost, or termination, layer. This is known as a termination switch[50]. This can be seen in RHEED analysis as the first oscillation will last longer than the following monolayers, essentially because 2 monolayers worth of material has to be deposited in order to form one monolayer of SrRuO 3. Figure 6.2 shows the oscillations in the RHEED signal during growth of SrRuO 3 on 2 ML thick LaAlO 3 on top of an SrTiO 3 (001) substrate. Each oscillation in this case appears to take about 60 seconds. This is atypical for the growth of SrRuO 3, which for these settings on SrTiO 3 takes about 60 seconds for the first monolayer, and 30 seconds for each subsequent monolayer. The reason for this is unclear Intensity (arb.u.) growth of first monolayer laser turned off Time (s) Figure 6.2: RHEED intensity over time for the growth of 5 ML of SrRuO 3 on 2 ML of LaAlO 3 on SrTiO 3 at 650 C. The inset shows the RHEED image after deposition of the SrRuO LaAlO 3 LaAlO 3 was grown at fluence of 1.3 J/cm 2 under a partial oxygen pressure of 100% oxygen at 0.1 mbar. Figure 6.3 shows the RHEED intensity over time for the growth of LaAlO 3. The inset shows a RHEED pattern after deposition of the film, the defined spots with slight streaking indicate that growth was 2D but the surface is slightly roughened. 40

46 Chapter 6: The metal-insulator transition in SrRuO 3 Intensity (a.u.) growth of one monolayer laser turned off Time (s) Figure 6.3: RHEED intensity over time for the growth of LaAlO 3 on SrTiO 3 at 650 C. The inset shows the RHEED image after deposition of LaAlO DyScO 3 A fluence of 2.5 J/cm 2 at a partial pressure of 100% oxygen at mbar was used in the growth of DyScO Intensity (arb.u.) growth of first monolayer laser turned off Time (s) Figure 6.4: RHEED intensity over time for the growth of DyScO 3 on SrTiO 3 at 650 C. The inset shows the RHEED image after deposition of DyScO Fitting the Ru 3d photoelectron spectrum As can been seen in figure 6.5, the fitting of the Ru 3d spectrum consists of three peaks for each spin-orbit coupling. The constraints for the peaks are such that the peaks for j=5/2 and j=3/2 have an area ratio of 5/ /2 2+1 = 2/3. The Sr 3p peaks, which are close to the Ru 3d j=5/2 peak, are also fitted and subsequently subtracted. Then the spectrum is normalised to the total area of the Ru 3d peaks. An example of the fit, without a C 1s peak, can be seen in figure 6.6. The most ideal way to quantify the screened peak is to use simulations on Ru 3d in order to determine the U and W parameters. However, the unfilled d-band increases the complexity of calculating Ru 3d XPS spectra considerably compared to e.g. 2p spectra. The program used in this report could not calculate the Ru 3d spectrum due to this complexity. The input files for a Ru 3d calculation without charge transfer parameters is given in Appendix A. 41

47 Chapter 6: The metal-insulator transition in SrRuO 3 3d 3/2 C 1s 3d 5/2 (RuO 4? Shakeup?) Unknown Unscreened Screened Figure 6.5: Peaks used to fit the j=3/2 and j=1/2 Ru 3d XPS spectrum Intensity (arb. u.) Binding energy (ev) Figure 6.6: Ru 3d spectrum and the fitted peaks underneath. 6.4 Thickness series SrCuO 2 As discussed in chapter 5, SrCuO 2 is postulated to have a transition from orthorhombic to a tetragonal infinite-layer structure around 5 ML. To test if this orthorhombic to tetragonal transition occurs, and if this has an effect on the MIT of a film of SrRuO 3, several samples with varying SrCuO 2 thickness were fabricated. The heterostructure can be notated as STO/x ML SCO/4 ML SRO. The SrRuO 3 film was grown 3 or 4 ML thick in order to keep the film at or around the previously known critical thickness on a bare SrTiO 3 substrate. The XPS spectra for these samples can be seen in figure 6.7. As can be seen, the screened peak changes very little with varying SrCuO 2 thickness. The only parameter which seems to influence the screened peak, increasing its intensity, is the lowering of the oxygen pressure during growth of the SrRuO 3 layer. This is possibly due to an increase in the oxygen vacancies in the SrRuO 3. 42

48 Chapter 6: The metal-insulator transition in SrRuO ML SCO (6e 2 SRO pressure) 22 ML SCO (3e 1 SRO pressure) 9 ML SCO (3e 1 SRO pressure) 9 ML SCO (6e 2 SRO pressure) 4 ML SCO (3e 1 SRO pressure) Intensity (arb. u.) Binding energy (ev) Figure 6.7: Ru 3d XPS spectra for several uncapped STO/x ML SCO/4 ML SRO samples. The thickness of the intermediate SrCuO 2 film is varied, as well as the oxygen pressure during SrRuO 3 growth. 6.5 Surface effects In materials, surface effects due to e.g. dangling bonds can play a role in the properties of thin films. The role of surface effects in the electronic and photoelectronic spectrum of SrRuO 3 can be investigated by capping the film with a few layers of material. The effects of capping on SrRuO 3 are investigated by looking at the differences between Ru 3d XPS spectra for uncapped and capped samples, for all intermediate materials used. The thickness of the SrRuO 3 layer is kept approximately constant to rule out effects caused by varying the thickness such as an increasing screened peak SrTiO 3 Figure 6.8 shows the effects of capping SrRuO 3 films with SrTiO 3 in the Ru 3d XPS spectrum. Films of SrRuO 3 of varying thickness were grown on an SrTiO 3 substrate and subsequently left uncapped, or capped with another layer of SrTiO 3. The effects for even a thin film seems to be an increase of the screened peak at 281 ev SrCuO 2 For the effects of capping with SrCuO 2, the Ru 3d XPS spectrum is shown in figure 6.9. The effect of capping in this case seems to be a further reduction of the width of the Ru 3d j=1/2 and j=3/2 peaks with the unscreened peaks centred at ev and ev, respectively. The capped sample appears higher in intensity compared to the other samples. This is because the normalisation procedure uses the total area of the fitted peaks, which has become smaller because with the reduction of the screened peak DyScO 3 The effects of capping with DyScO 3 are seen in the XPS Ru 3d spectrum in figure For comparison, a sample with thick SrRuO 3 without capping grown on DyScO 3 is also included. Capping with DyScO 3 appears to cause a splitting in the j=3/2 peak and, to a less visible extent, 43

49 Chapter 6: The metal-insulator transition in SrRuO STO/2ML STO 4ML SRO 2ML STO STO/16ML SRO STO/35ML SRO 4000 Intensity (arb.u.) Binding energy (ev) Figure 6.8: Ru 3d XPS spectra of thin films of SrRuO 3 of varying thickness, capped with SrTiO 3 or left without a capping layer STO/15ML SCO/4ML SRO STO/22ML SCO/4ML SRO STO/3ML SCO/3ML SRO STO/10ML SCO/3ML SRO STO/2ML SCO/3-4 ML SRO/2ML SCO Intensity (arb.u.) Binding energy (ev) Figure 6.9: Ru 3d XPS spectra of thin films of SrRuO 3 of approximately equal thickness, capped with SrCuO 2 or left without a capping layer. 44

50 Chapter 6: The metal-insulator transition in SrRuO 3 also in the j=1/2 peak. The DyScO 3 substrate causes tensile (0.744%) strain in the SrRuO 3, as opposed to compressive (-0.441%) strain on SrTiO DSO/thick SRO STO/2ML DSO/4ML SRO/2ML DSO STO/3ML DSO/3 ML SRO 4000 Intensity (arb.u.) Binding energy (ev) Figure 6.10: Ru 3d XPS spectra of thin films of SrRuO 3 of varying thickness, capped with DyScO 3 or left without a capping layer LaAlO 3 The effect of using LaAlO 3 as intermediate material in a sandwich structure can be seen in figure A sample without a LaAlO 3 capping layer was not fabricated. Therefore, an simple uncapped SrRuO 3 film on a bare SrTiO 3 substrate is used as a reference. The spectral shape of the sample capped with LaAlO 3 does not appear different from the uncapped one. The screened peak of the uncapped SrRuO 3 does appear to be more pronounced STO/2ML LAO/7ML SRO/2ML LAO STO/6 ML SRO 4000 Intensity (arb.u.) Binding energy (ev) Figure 6.11: Ru 3d XPS spectra of thin films of SrRuO 3 of approximately equal thickness, capped with LaAlO 3 or left without a capping layer. 45

51 Chapter 6: The metal-insulator transition in SrRuO Sandwiching structures of SrRuO 3 It is expected that the Ru 3d screened peak reveals the transition from a metal to an insulator, depending on the parameters W and U. To determine the change in Hubbard parameters U and W, the final fit values are evaluated. Firstly, the real value of the bandwidth W cannot be given exactly, since an ab-initio model for this value is absent. To still get an idea of the value of W, the ratio r s,u = (A s /FWHM s )/(A u /FWHM u ) of the weight of the fitted screened peak (A s /FWHM s ) to the weight of the unscreened peak (A u /FWHM u ) is determined. The separation in binding energy of the screened and unscreened peak δ s,u is a measure electron correlation U. The separation in binding energy as previously discussed in chapter 2 is given as Q U/2. Since the attractive coulomb potential Q of the core-hole is not known, it is necessary to judge the energy U as a relation that is inverse with the separation of the screened and unscreened peak. As the separation decreases, the electron correlation energy U increases SrRuO 3 and LaAlO 3 The sandwich structure using LaAlO 3 as an intermediate material was investigated with varying thickness of the SrRuO 3 film. The Ru 3d XPS spectra are visible in figure 6.12, with the SrRuO 3 layer 4 ML and 7 ML thick. The screened peak visible as a shoulder at about 281 ev shows a slightly different separation distance δ s,u for both structures: 2.3 ev for 4 ML and 2.2 ev for 7 ML. The weight of the screened peak was determined from the fitted peaks to be 0.23 and 0.21 for 4 ML and 7 ML, respectively. By eye, the screened peak looks to be more pronounced in the 7 ML thick sample Sandwiching (Ru 3d) STO/2ML LAO/4ML SRO/2ML LAO STO/2ML LAO/7ML SRO/2ML LAO 4000 Intensity (a.u.) Binding energy (ev) Figure 6.12: Ru 3d XPS spectra for 4 and 7 ML of SrRuO 3 sandwiched with 2 ML of LaAlO The screened peak in all sandwich intermediate materials Sandwich structures with intermediate materials SrTiO 3, LaAlO 3, SrCuO 2, and DyScO 3 were grown to observe the effects of different octahedral rotations on the behaviour of the screened Ru 3d peak and transport properties. The XPS spectra of the sandwich structures shown in figure 6.13 were fitted and normalised using the fitting procedure described earlier. Table 6.1 shows the calculated parameters r s,u and δ s,u derived from the fitted peaks, for different materials in the sandwiching heterostructure. For DyScO 3 and SrCuO 2 with 4 ML of SrRuO 3, the screened peak could not be fitted. 46

52 Chapter 6: The metal-insulator transition in SrRuO STO/2ML LAO/4ML SRO/2ML LAO STO/4ML SRO/2ML STO STO/2ML DSO/4ML SRO/2ML DSO STO/2ML SCO/3 4ML SRO/2ML SCO STO/7ML SCO/15ML SRO/2ML SCO Intensity (arb. u.) Binding energy (ev) 280 Figure 6.13: Ru 3d XPS spectra of sandwiched structures of SrRuO 3. The materials LaAlO 3, DyScO 3, SrTiO 3, and SrCuO 2 are used as intermediate materials. Sandwich material δ s,u r s,u LaAlO SrTiO DyScO 3 Screened peak absent in fit - SrCuO 2 (with 4 ML SrRuO 3 ) Screened peak absent in fit - SrCuO 2 (with 15 ML SrRuO 3 ) Table 6.1: Fitted Ru 3d parameters for sandwiched heterostructures. As figure 6.13 shows, the sandwiched structure with SrCuO 2 and a 15 ML SrRuO 3 thick film still shows a very much repressed screened peak. This in contrast to e.g. the screened peak seen in a thick SrRuO 3 sample in a DyScO 3 sandwich structure seen in figure While the area ratio of the screened peaked is high, the peak separation is low compared to the sample with DyScO 3 and LaAlO 3. Something which also can readily be seen by eye. The SrCuO 2 sandwich structure with 4 ML of SrRuO 3 shows an even more repressed screened peak. Structures with LaAlO 3 show a prominent screened peak at a separation that is comparable to the SrTiO 3 structure, which are known to be conducting. DyScO 3 structures show a splitting and no screened peak could be fitted Transport measurements With the general trends of the W and U parameters now known from the previous section, the relationship between the ratio W/U and the transport properties of the films can be determined. The sandwich structures in the previous section were probed for their transport properties using the PPMS measuring system in a van der Pauw setup. The calculated sheet resistances are shown in figure In some cases it was not possible to measure resistances in both bridges of the van der Pauw setup due to e.g. too high resistance in the film. This made it impossible to calculate the sheet resistance. In these cases the resistivity was not calculated, but a comparison of all resistances can be seen in figure 6.14(a). The upward slope for negative temperature for the samples with the sandwich structures with LaAlO 3 and SrCuO 2, indicate that they are not conducting. The sandwich structure with SrTiO 3 is the only structure that shows conducting behaviour for films with a thickness of about 4 ML. Referencing the XPS spectra in figure 6.13 and the ratio in table 6.1, this could be understood by the differences in δ s,u and r s,u. For LaAlO 3, the ratio of W/U would be too small due to the 47

53 Chapter 6: The metal-insulator transition in SrRuO 3 Resistance (Ω) 10 7 STO/4 SRO/2 STO STO/2 SCO/4 SRO/2 SCO STO/40 SCO/3 SRO STO/2 DSO/4 SRO/2 DSO STO/6 SCO/6 SRO/2 SCO 10 6 STO/2 SCO/4 SRO/2 SCO STO/2 LAO/4 SRO/2 LAO STO/10 SCO/3 SRO STO/6 SCO/15 SRO/2 SCO Temperature (K) 10 7 STO/4 SRO/2 STO STO/40 SCO/3 SRO STO/2 DSO/4 SRO/2 DSO STO/6 SCO/6 SRO/2 SCO 10 6 STO/2 LAO/4 SRO/2 LAO Resistivity (µ Ω cm) Temperature (K) Figure 6.14: Transport properties of SrRuO 3 films in sandwich structures using several intermediate materials with (a) the (sheet) resistance as measured in a van der Pauw setup. Dashed lines indicate that only one bridge in the van der Pauw setup could be measured. A measurement indicated with a cross could not reliably be measured, and is given as the resistance at room temperature. (b) shows all films where resistivity could be calculated. 48

54 Chapter 6: The metal-insulator transition in SrRuO 3 weight W of the peak, represented by r s,u, is too small when observing the conductive sample of SrTiO 3 with a r s,u that is slightly higher. For SrCuO 2 (both with 4 ML and 15 ML SrCuO 2 ), the ratio r s,u would be high enough. However the peak separation δ s,u in this case indicates that U is high, leading to a lower W/U, overall. 6.7 Insulating behaviour of SrCuO 2 /SrRuO 3 /SrCuO 2 sandwich structures The conduction and spectral measurements of the sandwich structures using SrCuO 2 indicate that SrRuO 3, even when grown as thick as 15 ML, is still not conducting. This insulating behaviour at higher thicknesses is unexpected. For higher thicknesses, the current assumptions about octahedral rotations suggest that for higher thicknesses SrRuO 3 films should return to a conductive state due to octahedral relaxations. This anomalous behaviour could be caused by doping of Cu atoms from the SrCuO 2 films by means of intermixing. To investigate this idea, angle dependent XPS spectra were measured and subsequent peak area ratios of the constituents of the film (Cu, Ru, Sr) were calculated for several angles. Figure 6.15 shows the depth profile of a STO/6 ML SCO/15 ML SRO film. This film lacks capping so any intermixing of Cu atoms will be only from the first intermediate layer of SrCuO 2. The AFM image of the film can be seen in figure The film has small features that might be islands a maximum of 10 nm high. The rest of the film has areas which shows holes of a maximum of 2 nm. The film, not including the islands, has a root-mean-square roughness of 0.6 nm. The Ru 3d XPS spectrum shows the same behaviour for the screened peak as for the capped structure Ru 3d/Sr 3p Cu 2p L/Sr 3p Cu 2p/Sr 3p Cu 2p L/Ru 3d 1.2 Intensity ratio Detector Angle (θ) Figure 6.15: Angle dependent XPS ratios of a STO/6 ML SCO/15 ML SRO thin film. The depth profile shows a Cu 2p signal at 10. At this angle, the sampling depth would be around 1.5 nm (about 5 ML). The ratio between Cu and Ru shows a slight increase going from 10 to 30 and shows no difference going to 80. This could indicate that Cu is present in the top layers Reciprocal space-mapping To determine the crystal structure and lattice parameters of SrRuO 3 as grown in the SrCuO 2 /SrRuO 3 /SrCuO 2 sandwich structures, reciprocal space maps were made of a sample with an STO/6 ML SCO/15 ML SRO/2 ML SCO heterostructure. Previously measured reciprocal space maps found in literature of SrRuO 3 on SrTiO 3 and SrRuO 3 on DyScO 3 can be seen in figure Figures 6.19 and 6.18 show the reciprocal space map on the (204) and (114) reflections on the STO/6 ML SCO/15 ML SRO/2 ML SCO heterostructure. The SrTiO 3 substrate peak is clearly visible in the top half of the the scans, while the SrRuO 3 film peak streaks diagonally to the bottom of the images at a slight angle. The angle for the (114) and (204) reflection was measured to be approximately 49

55 Chapter 6: The metal-insulator transition in SrRuO 3 Figure 6.16: (a)afm image of a STO/6 ML SCO/15 ML SRO thin film and (b) the corresponding height profile Vailionis, Siemons, and Koster Appl. Phys. FIG. 3. Color online Schematic representat a b c and b tetragonal a=b c structures rhombic SRO, RuO 6 octahedra are rotated counter 001 directions and clockwise about 100 directi ture RuO 6 octahedra are rotated around 001 dir consistent with the sign of the strain indi FIG. 2. Color online Reciprocal lattice maps of SrRuO 3 layers grown on cell, in addition to the variations of a, b a SrTiO and b DyScO single crystal substrates. SrRuO 3 on Figure 6.17: Reciprocal lattice maps of SrRuO3 layers grown on SrTiO eters, utilizes this additional degree of fr SrTiO 3 clearly shows an orthorhombic unit cell, while SrRuO 3 layer 3 (001) and DyScO on 3 (110) single crystal substrates. DyScOSrRuO 3 is tetragonal. 3 on SrTiO Here we 3 used clearly Q =4 shows sin /, anwhere orthorhombic is the Bragg unit date cell, thewhile mismatch the between the substrate SrRuO 3 layer on DyScO angle 3 is andtetragonal. = Å. Here Indices weof used the layer s Q out o reflections f plane are = shown 4πsin(θ)/λ, negative where mismatch θ the becomes relatively bold. Bragg angle and λ = Å. Indices of the layer s reflections are shown in bold. and (Image CRO onfrom STO the ARO layer stab [51]). instead of an orthorhombic structure. Th in high- and medium-resolution modes at the Stanford Nanocharacterization Laboratory, Stanford University. films is induced by epitaxial strain. It bulk materials takes place at higher t The XRD results demonstrate that epitaxially grown that orthorhombic-to-tetragonal O-T tr ARuO 3 films exhibit 110 out-of-plane orientation with CRO occurs at different mismatch val 100 and 110 in-plane orientations. Reciprocal lattice from Table I, CRO becomes tetragonal maps from symmetrical and asymmetrical Bragg reflections match values of about 1.78%, while fo allowed us to determine the50unit cell size and shape of the occurs at much lower mismatch of ab ARuO 3 layers. As an example, Fig. 2 shows reciprocal lattice large dissimilarity can be explained by maps of the SrRuO 3 260, 444, 620 and 44-4 reflections as well as the SrTiO and DyScO 3 260, 444, materials. For bulk SRO orthorhombic hombicity factors ratio of a and b latti much smaller than that of CRO which