Modern Growth Problems and Growth Techniques

Transcription

1 1 Modern Growth Problems and Growth Techniques Björgvin Hjörvarsson 1 and Rossitza Pentcheva 2 1 Department of Physics, Uppsala University, Box 530, Uppsala, Sweden bjorgvin.hjorvarsson@fysik.uu.se 2 Department of Earth and Environmental Sciences, Section Crystallography, University of Munich, Theresienstrasse 41, Munich, Germany pentcheva@lrz.uni-muenchen.de Abstract. The growth and characterization of magnetic materials has progressed substantially during the last decades. In this chapter we give a brief overview of this vastly growing field of research. We highlight some of the relevant growth techniques for different materials classes but we do not intend to be complete with respect to technical details or materials systems. We also outline some of the concepts and theories of the growth of modern magnetic materials, emphasizing the role of first principles calculations in providing microscopic understanding of the growth mechanisms. We discuss the growth of metallic and oxide single crystal films, multilayers and superlattices and the influence of thickness, strain, crystallinity, structure and morphology on the resulting magnetic properties. 1.1 Growth and Characterization The growth and characterization of magnetic materials has progressed substantially during the last decades. In this chapter we give a brief overview of this vastly growing field of research. We highlight some of the relevant growth techniques for different materials classes but we do not intend to be complete with respect to technical details or materials systems. We also outline some of the concepts and theories of the growth of modern magnetic materials, emphasizing the role of first principles calculations in providing microscopic understanding of the growth mechanisms. We discuss the growth of metallic and oxide single crystal films, multilayers and superlattices and the influence of thickness, strain, crystallinity, structure and morphology on the resulting magnetic properties Concepts: The Thermodynamic Versus Kinetic Picture The first concepts used to describe crystal growth were based on general thermodynamic considerations [1]. The resulting structures were assumed to B. Hjörvarsson and R. Pentcheva: Modern Growth Problems and Growth Techniques, STMP 227, 1 44 (2007) DOI / c Springer-Verlag Berlin Heidelberg 2007

2 2 B. Hjörvarsson and R. Pentcheva be in thermal equilibrium and therefore determined by the minimum of the free energy. Depending on the balance between the surface energies of the substrate material, γ S, the adsorbate layer, γ A, and the interface energy, γ I, three different growth modes are distinguished (shown in Fig. 1.1):. Δγ = γ S γ A γ I. (1.1) If Δγ > 0, wetting of the substrate by the deposited material is expected, resulting in layer-by-layer growth mode (Franck van der Merwe mode). In this case, a new layer starts to grow only after the first one is completed. In the opposite case, Δγ < 0, the formation of three dimensional islands is likely to occur (Vollmer-Weber growth mode). An intermediate situation can appear when Δγ > 0 and the growing film is strained. Initially the film grows in a layer-by-layer mode, up to a critical thickness where the growth becomes island like. This mode is often referred to as Stransky-Krastanov growth mode. The transition from a layer-by-layer growth to island like growth can be viewed as governed by the strain state and the elastic properties of the growing material. Although useful, this classification has its limitations, it does, for example, not include surface alloying. Furthermore in the homoeptiaxial case, i.e. when substrate and adlayer are of the same material, it predicts a layer-by-layer growth. This is not generally valid and different growth modes are observed for homoepitaxial systems. For example, Ag forms three dimensional islands on strained Ag(111) [2, 3]. Furthermore, ferromagnetic materials (e.g. Co, Fe) have a higher surface energy than the noble metals (e.g. Cu, Ag). Accordning to this simplified view, Co and Fe should not grow in a layer-by-layer mode on Cu and Au substrates. However, Co deposited on Cu(001) grows up to twenty monolayers (ML) in a layer by layer mode [4]. These are just a few examples illustrating the possibilities and limitations in the thermodynamic description of growth of materials. The main reason for the failures is that the growth process is by definition a non equilibrium process and the thermodynamic equilibrium condition is thereby not fulfilled. In this non-equilibrium kinetic process one or more steps can be rate limiting. By understanding the underlying processes, the growth procedures can be γ S γ A γ I Frank-van der Merwe Vollmer-Weber Stranski-Krastanov Fig The illustration on the left hand side describes the basis for the thermodynamic description of growth. The relative energy of the interface (I), the surface of the growing layer (A) and the substrate surface (S) is assumed to determine the resulting growth mode. Schematic illustration of the three basic types are also included in the figure

3 1 Modern Growth Problems and Growth Techniques 3 used to control the morphology and the crystallinity and thereby some of the emerging material properties. In an atomistic approach, the growth of a film can be described as a result of a number of microscopic processes such as adsorption, diffusion and desorption of adatoms (cf. Fig. 1.2). Diffusion of atoms can take place on flat regions of the substrate, along or across step edges or around island corners. Therefore, besides adatom diffusion also the adatom-adatom and adatomstep interaction determine island nucleation and growth. In the framework of transition state theory (TST) [5], surface diffusion is described by diffusion rates D which are determined by diffusion barriers, E d, and prefactors, D 0. D = D 0 e E d k B T. (1.2) Time scales relevant for sample growth are of the order of seconds and minutes and the length scales of kinetically controlled structures and islands are of the order of 100 Å and involve a large number of atoms (> 10 5 ). On the other hand the detailed quantum mechanical description of atomistic processes is currently restricted to relatively small system sizes, up to about 10 4 atoms. Typical time scales of e.g. ab initio molecular dynamics are of the order of picoseconds which limits its application to the determination of possible processes, probable paths, diffusion barriers and attempted jump rate (prefactors) of the adatoms. A phenomenological or statistical description of growth can for example be obtained by using nucleation theory [6] or kinetic Monte Carlo simulations. These methods are often based on empirical or semiempirical parameters and their predictive power is therefore limited. In nucleation theory, growth is described by rate equations, yielding the time evolution of the adatom and island density. When the desorption rate is negligible a simple relation between the saturation island density n x, deposition rate R, diffusion rate D and temperature is obtained [6]: n x (R/D) i/(i+2). (1.3) adsorption deposition diffusion via hopping substitutional adsorption Fig Atomistic picture of growth, including different processes like deposition, adsorption, diffusion of adatoms on the terrace, incorporation into existing islands, as well as incorporation via exchange in the substrate layer

4 4 B. Hjörvarsson and R. Pentcheva Typically, the critical island size corresponds to i = 1, which implies that two adatoms form a stable configuration. The linear dependence that follows between ln n x and 1/T is often used to extract the diffusion barrier from the island density at a constant deposition rate. It can also be used to determine the critical island size i from the deposition rate dependence of island density at a constant temperature. The graphical representation of ln n x (1/T )isoften referred to as Arrhenius plots. The rate equations express the time evolution of the average adatom and island density. Because it is a mean field approach, immediate and constant adatom density in the vicinity the growing islands is assumed, while in reality the islands have depletion zones with lower adatom densities. For systems where medium-range interactions are important (e.g. on strained surfaces), island densities predicted from nucleation theory can differ by as much as an order of magnitude when only short range interactions are considered. This was shown in a DFT-KMC-study [7] and emphasizes the need for including stress and strain in the theoretical considerations. For a review of the microscopic view on metal homoepitaxy, see [8]. Nucleation theory is restricted to adatoms forming islands on the surface and an exchange with the underlying material is not taken into account. However, exchange processes between adsorbate and substrate can significantly influence the resulting heteroepitaxial growth modes. An attempt to include exchange mechanism, within nucleation theory, was obtained by introducing systems with a critical island size of zero, i = 0 (see e.g. [9]). Some further aspects of metal heteroepitaxy will be discussed in Sect When describing the growth of thin films within the framework of statistical mechanics, the exact motion of an adatom is irrelevant. The motion is treated as a random process, while the probability as a functional of the energy of a particular configuration is exact. The kinetic Monte Carlo approach provides a statistical description of the evolution with time, enabling realistic description of a non-equilibrium processes. Combining DFT results on diffusion barriers and chemical interactions on the atomic scale with a statistical description of the time evolution in a kinetic Monte Carlo simulation (ab initio kmc) is therefore a promising route to bridge the gap between the time and length scales of first principle calculations and experimental observations [10, 11, 12] Growth Techniques After this brief overview on the theoretical description of growth, we now consider some of the experimental aspects of the involved processes. The most commonly used deposition techniques for growing magnetic materials are: molecular beam epitaxy (MBE); magnetron sputtering ion beam sputtering

5 1 Modern Growth Problems and Growth Techniques 5 pulsed laser deposition (PLD); metal organic chemical vapor depostition (MOCVD); These deposition techniques have many similarities, but do also differ substantially with respect to the underlying physical processes. In this section, the complementary aspects will be emphasized and general requirements for successful deposition of different materials will be addressed. The crystal coherency and the chemical purity are the most important parameters describing the quality of as grown samples. Surface oxidation and intermixing at the substrate interface are issues of concern for single films, while interface mixing and thickness variation become important when discussing multilayered structures. The chemical purity of the as grown structures depends strongly on the purity of the ambient. For this purpose, the vacuum conditions can be used as a qualifier. Under ultra high vacuum (UHV) conditions, corresponding to 10 9 mbar or lower, the impinging rate of the residual gas is below 10 3 /s. Thus, with a growth rate of 1 monolayer (ML) per second, the impurity level of the samples originating from the vacuum environment can be below 10 3 (atomic ratio), which is comparable to a representative purity of commercially available deposition material. The residual gas in a tight UHV system is typically governed by H 2.The partial pressure of water (p H2O) is typically one or two orders of magnitude below that of H 2.p O2 is often below the detection limit of most residual gas analyzers (10 13 mbar) and can be ignored in this context. Although p H2O is orders of magnitude below that of p H2, the influence of H 2 O can dominate the impurity levels. The sticking coefficient of H 2 O is close to unity while that of H 2 is typically 1 at 300 K. Furthermore, the sticking coefficient is strongly dependent on the chemical composition of the surface as well as the temperature and has therefore to be considered with extreme care. When using magnetron sputtering, the pressure during growth is typically in the 10 3 to 10 1 mbar range. Using UHV growth chambers and ultra pure gases, the partial pressure of impurities in the sputtering gas can be in the same range as in e.g. MBE systems. Impurity levels of bottled gases are at the ppm level at best which can result in adequate chemical purity, if the gas is not contaminated in the gas handling system. The composition of the low pressure ambient can influence the physical properties of the growing material. For example, the presence of water will inevitability cause H impurities, significantly altering e.g. optical and elastic properties of oxide films [13, 14]. Furthermore, when growing metallic layers, this will result in both H and O impurities in the films. The partial pressure of active gases in the sample environment is therefore a good measure of the possibility to grow samples with high chemical purity. The noble gases are not to be considered in this context, due to their inertness with respect to chemical reactions. The use of all metal connections and thoroughly outgassing the gas lines as well as the deposition system (baking) can be a simple route to improve the

6 6 B. Hjörvarsson and R. Pentcheva sample quality. This approach reduces the contamination of the process gas, however, the purity of the commercially available gases can be insufficient. The use of gas purifiers is therefore sometimes needed for obtaining the required purity levels. Getter materials and cold traps based on molecular sieves can be used for this purpose. The combination of all-metal bakeable gas lines and gas purifiers allows significant reduction of impurity levels. Although the sputtering gases are chemically inert, the inclosure of the sputtering gases can cause significant contributions to the chemical composition. Typically the sputtering gases (Ar, Kr, etc.) are found in voids created during growth [15]. This effect can be profound in the context of high growth rates of polycrystalline materials but is almost always negligible in single crystal materials. The substrate temperature determines the sticking coefficient, the surface mobility and the desorption rates, and is therefore one of the most important process parameters. To determine the actual growth temperature is experimentally challenging and the quoted numbers are typically rather inaccurate. This is of special importance with respect to the growth of heteroepitaxial materials. This is important especially in the case of heteroepitaxial growth. Interdiffusion and crystallization are competing processes which can result in a narrow temperature range available for the required growth processes. However, reproducible temperatures can easily be obtained in any growth system, which enables reproducibility of the growth while using the same setup. The concern is therefore the transfer of experimental procedures between growth systems, as temperature calibrations are often crude and are material and substrate dependent. Substrates Oxides and semicoductors are frequently used as substrates for growing magnetic films, multilayers as well as superlattices. The benefits from this choice are many. First of all, the availability and price. Single crystal MgO, Al 2 O 3, SrTiO 3, Si, Ge, GaAs, etc., with different orientations, are commercially available. The surface and bulk crystalline quality varies and the influence of exposure to ambient air differs substantially. The variation is not limited to different suppliers, large differences in substrate quality can be found from one and the same vendor. This is clearly seen in Fig. 1.3, which illustrates rocking curves around the MgO(002) peaks from two different substrates. One of the substrates shows a well defined peak (solid curve) consistent with a single crystal structure, while the second one exhibits number of peaks corresponding to crystallites which are only partially oriented. These substrates were obtained from the same vendor and represent two batches of the same material. The pre-treatment of substrates are both performed ex- and in-situ, depending on the purpose of the treatment. For example, ex-situ heat treatment of Al 2 O 3 (1500 Cfor1 5 hours, see Fig. 1.4) and SrTiO 3 are commonly

7 1 Modern Growth Problems and Growth Techniques Counts Counts , , ,5 Angle (deg.) Fig Rocking curve from a MgO(001) sample representing a good (solid) and a bad (dashed line) batch from the same supplier. The measurements were performed using Cu K α radiation around the (002) peak of MgO [17] used, increasing both the terrasse width and crystalline quality. Corresponding heat treatment of MgO destroys the surface completely, easily identified by an opaque appearance. In-situ treatment almost always involves extended annealing, removing water adsorbed at the surfaces. The required temperature and time depends on the adsorption energy of water on the surface, which can vary substantially between different materials and crystallographic orientations [16]. High temperature annealing of semiconductor substrates can result in inward diffusion of the near surface oxides. Pre-sputtering is required to avoid this, but post annealing is required for obtaining smooth surface after sputter cleaning. Pre-sputtering is normally not utilized when working with oxide substrates. Fig AFM pictures showing the development of a step pattern during the annealing in air of sapphire substrates for six hours at different temperatures [18]

8 8 B. Hjörvarsson and R. Pentcheva Molecular Beam Epitaxy and Sputtering Molecular Beam Epitaxy (MBE) and sputtering techniques are commonly utilized for growing magnetic thin films and superlattices. The main differences between these two techniques is the energy of the material flux, reaching the substrate and the different conditions of the evaporating material. In MBE growth, the material is heated to a temperature giving the desired vapor pressure, which can be either below or above the melting temperature of the material. For example, Mg has a high sublimation rate far below the melting temperature. Mg is therefore typically not melted during an evaporation process. This involves both limitations and possibilities, as in situ cleaning/outgassing of e.g. Mg becomes difficult. Secondly, the evaporation temperature of the material defines the kinetic energy of the atoms reaching the substrate and later the growing film. Thus, the energy of the impinging atoms are typically far below 0.2 ev in an evaporation process. In UHV based evaporation systems such as MBE, the mean free path is much larger than the system size. In magnetron based sputtering processes, the typical mean free path is of the order of centimeters. The flux from the evaporation source can therefore be regarded as highly directional, while the flux from magnetron sources is close to random at the substrate surface (covering 2π). This has profound implications on the possibilities of using masks to obtain patterned growth [19], for which e.g. MBE is much better suited as compared to magnetron sputtering. There are two main routes for evaporating materials, namely through direct heating as accomplished in effusion cells (through radiation or direct heating from heating elements) and by direct bombardment using high energy electron beam (e-beam). Effusion cells have much better stability with respect to the flux of the evaporated material and is therefore the method of choice when well defined layer thicknesses are required. The stability is critical when growing, for example multilayers and superlattices, where the layer thicknesses have to be extremely well defined. The instability in the material flux from an e-beam source originates from the dynamics of the melted region. Scanning the electron beam across the target material often increases the stability, but the resulting fluctuations in the flux are still much larger than obtained from effusion cells. The limitations in the use of effusion cells originate from the chemical and thermal stability of the crucibles which are in direct contact with the evaporating material. Typical crucible materials are Al 2 O 3, BeO, C (pyrolytic graphite), Ta and W. For the growth of materials consisting of more than one element from a single source, the use of evaporation techniques can be inferior to that of sputtering. As the flux of the evaporating material is determined by the vapor pressure of the elements, the growing film is likely to have significantly different chemical composition, as compared to the composition of the source material. In a sputtering process, the surface composition changes initially but

9 1 Modern Growth Problems and Growth Techniques 9 eventually compensates the different cross sections of the elements, resulting in a close matching of the film composition to that of the source (target material). Sputtering techniques can be viewed as complementary to MBE, not allowing any in situ treatment of the target material, but having wide flexibility with respect to the kinetic energy of the material flux. The kinetic energies relevant in the sputtering process must be divided into two classes, neutrals and charged particles. The energy of these can, in principle, be adjusted independently. The neutrals as well as the charged particles are directly affected by collisions with the residual gas, while a bias of the sample only affects the ratio of the kinetic energy of positive and negatively charged particles. Positive bias retards the positively charged sputtering gas and target material, while it accelerates the electrons and negatively charged atoms of the target material toward the growing film. This technique has been used to alter the morphology as well as the crystallinity of thin films and superlattices. The simplest and often used route to utilize the electric potential, is to keep the sample at a floating bias. An alternative mode of operation is reactive sputtering. By introducing reactive gases in the growth chamber, these will readily react with e.g. metals in an ionic or a neutral state. Reactive sputtering can be used to grow oxides, nitrides or any material where one of the components can be introduced in the gas phase feeding the sputtering process. The plasma chemistry can also be used for obtaining phases which can not be synthesized by regular chemistry under normal conditions. There are two main modes of sputtering, namely Radio Frequency sputtering (RF-sputtering) and Direct Current (DC-sputtering). Although the DC mode can be used for reactive sputtering, target poisoning poses a severe challenge, demanding high degree of control of the pressure and the electric potential at the target. RF sputtering is one possible route to remedy this, even allowing the use of an insulating target material. The corresponding MBE approach is denoted Oxygen Plasma Assisted Molecular Beam Epitaxy (OPAMBE), and is used for oxide growth as the name indicates. Although sputtering has a wide range of applications, there are some materials that cannot successfully be deposited using this technique. As an example, growth of high purity actinide and lanthanide films is typically restricted to ultra high vacuum evaporation. The chemical purity of the purchased materials is specified for all elements but hydrogen, which is typically high. The processing required for obtaining good quality films, involves therefore extended outgassing prior to deposition for reducing the hydrogen content. This processing is not compatible with standard sputtering techniques. The growth of the actinides and lanthanides is therefore typically restricted to UHV evaporation. Although great precaution is taken with respect to in-situ purification, substantial amounts of hydrogen is inevitably present in rare earth films, as will be discussed later.

10 10 B. Hjörvarsson and R. Pentcheva Other Techniques Both MOCVD and PLD are widely used for growing oxides and semiconductors. In the first technique the chemistry in the reaction chamber can be tailored, while some restrictions apply with respect to the most reactive elements such as Sc, Y, Lu, Lr and the rare earth elements. The use of in-situ tools is restricted due to the high pressures in the reaction chamber. In the magnetic community MOCVD is mainly used for the growth of doped ferromagnetic semiconductors. The use of ion beam techniques for growth of magnetic samples for research purposes, has increased substantially. The main advantage of this technique is the versatility, allowing large number of target materials in the one and same setup. Although the use has increased, it is still not widely used. Pulsed-laser deposition (PLD) has been utilized for the growth of single films and superlattices. This technique is often favored for the growth of complex oxides, allowing complete control of the ambient gas. Here, a laser beam is focussed on a target in an UHV chamber. Material ablated by the laser pulses is deposited on the substrate. This technique has two major advantages: First, the target has already the desired stoichiometry making the oxidation step superfluous and second, the amount of deposited material can be controlled/calibrated by the number of pulses, allowing high degree of precision of the thicknesses of the layers while growing thin films and superlattices [20, 21]. A finite oxygen pressure is often required to obtain stoichiometric oxides Characterization Techniques In situ Characterization Both in-situ and ex-situ characterization tools are used for obtaining information on the composition and structure of the as-grown materials. These can be viewed, in many respects, as complementary. The big advantage of the in-situ techniques is the absence of the limitations invoked by the exposure to ambient, while the advantage of ex-situ characterization lies in the versatility. Some aspects of the structural quality of the growing materials can be determined in-situ, however, most of the relevant work is done ex-situ, when operating a production device for thin films and heterostructures. The utilization of in-situ characterization is demanding and significantly increases the complexity level of the operation. Therefore many production systems are only equipped with a limited amount of in-situ possibilities. Examples of common in-situ equipment are typical surface characterization tools such as Auger spectroscopy, Reflective High Energy Electron Diffraction (RHEED), Low Energy Electron Diffraction (LEED) as well as Scanning Probes such as Atomic Force Microscopy (AFM) and Scanning Tunneling Microscopy (STM). The most widely used probe for in-situ measurements is the electron. For example, probing the energy of the ejected Auger electrons gives information

11 1 Modern Growth Problems and Growth Techniques 11 about the near surface composition, diffraction from the surface yields the surface structure and electron energy loss spectra (EELS) even yields information on the oxidation state. When determining the surface structure, there are two basic configurations, the electrons arrive to the surface almost parallel to the surface normal (LEED) or close to parallel to the surface plane (RHEED). Consequently, these techniques probe different directions and lengthscales, RHEED is highly sensitive with respect to steps and island growth while LEED yields information on the atomic distances and periodicity in the near surface region of the sample. An advantage of LEED with respect to other structural techniques such as e.g. XRD, is the higher sensitivity to the oxygen positions and the lower penetration depth making the method sensitive to the near surface structure. Thus, LEED is used to determine the crystal quality, in-plane lattice parameters, and the development of superstructures. By simulations of the LEED I V curves, information on the atomic positions and interlayer distances of the outermost layers can be obtained. RHEED is mainly used to monitor the thickness of evolving films and the growth quality. Oscillations in the RHEED intensity correspond to the formation of complete layers (closing of layers) and thus can be directly related to the thickness of the deposited film. The use of STM and AFM allows the visualization of the real space surface morphology, thereby serving as a complementary tool to reciprocal techniques such as LEED and RHEED. Both techniques have been used to follow the diffusion of adatoms on surfaces as well as monitoring the surface quality of thin films and heterostructures. Also diffusion parameters from island density measurements using Arrhenius plots are described in the literature. The combination of these techniques allows contact microscopy, enabling identification of chemical elements with near to atomic resolution. Thus, these scanning probe techniques are exceedingly valuable tools on all growth systems. There are also techniques which have to be classified as in-situ techniques, although these are not used as a routine equipment for monitoring growth in conventional growth systems. One of the most important ones is the force ion microscopy (FIM) which allows the study of individual diffusion processes and determination of diffusion barriers. For a comprehensive review see Tsong [22]. Ex situ Characterization Techniques As mentioned above, some aspects of the structural quality of the growing materials is preferably determined in-situ, however, when operating a production device for thin films and heterostructures most of the routine characterization is performed ex-situ. Two of the most important parameters describing the resulting films are the chemical composition and the crystalline structure. The composition of thin films is conveniently determined by Ion Beam Analysis techniques (IBA), such as Rutherford Back Scattering (RBS) and

12 12 B. Hjörvarsson and R. Pentcheva Nuclear Resonance Analysis (NRA). These techniques have been used to determine the composition of wide variety of materials such as metals [23, 24, 25], oxides [14], nitrides [26], carbides [27] and hydrides [28, 29, 30]. The signal in a RBS experiment scales as Z 2, thus, the sensitivity increases strongly with increasing atomic number. High energy scattering can be used for changing the scattering cross section, when a better sensitivity is required for low Z materials. The basic idea is to overcome the Coulomb barrier, entering a region where the cross section is highly varying with the energy of the impinging ions. The presence of resonances can even be used to obtain isotope selective depth profiling in materials. For example, the resonance between α and O 16 can be used for depth profiling of oxygen [31] with a resolution and sensitivity which is far better than most other techniques. When the depth resolution is less important than the detection limit, Elastic Recoil Detection Analysis (ERDA) can be utilized [32]. These techniques can all be operated in an absolute mode, counting the number of ions hitting the sample. This in combination with the known scattering cross section allows absolute determination of the concentration with an accuracy only limited by the determined stopping power of the ions. Typical accuracies are below few atomic percent, while the precision canbemuchbetter. Most of the structural information is obtained by x-ray scattering using conventional laboratory sources. The combination of reflectivity and diffraction allows the probing of all relevant length scales, from the overall thickness of the film to the the atomic distances. The x-ray scattering yields the crystal coherency of the material, the sharpness of the chemical modulations i n multilayers and superlattices as well as the thickness variation at all relevant lengthscales [33, 34, 35]. However, interdiffusion and thickness variation (see Fig. 1.5) in a superlattice can not be separated by solely performing specular scattering experiments. Only by combining off-specular and specular scattering the relative weight of the components can be separated. Simulations and fitting routines for off-specular scattering are currently not generally available. The combination of x-ray scattering and Transmission Electron Microscopy (TEM) can be useful. For example, when investigating combinations of oxides and metals, the difference in the scattering cross section is often large enough to obtain near atomic resolution of the interface composition. This combined with large contrast in the x-ray scattering allows detailed comparison giving unique insight in the actual local and global variation in the chemical composition in samples. However, TEM is highly destructive technique. The sample preparation involves slicing and thinning of a cross sectional part of the sample. Consequently, the investigated samples can not be used for any other measurements. Examples from investigations using most of these techniques will be given in the following sections. Investigations of the structural quality is not restricted to the use of x-ray scattering. For example, the use of ion beam channeling can be highly rewarding. When the ions channel through a single crystal film, the back scattering yield is small. If the film and the substrate have a coherent interface, the yield

13 1 Modern Growth Problems and Growth Techniques 13 Fig Illustration of interdiffusion (left) and roughness (right) at interfaces.if it is possible to insert a surface which separates all the atoms of the two types, there is no interdiffusion will remain small. On the other hand, if the interface is incoherent, there will be substantial back scattering from the region corresponding to the interface between the film and the substrate. Ion beam scattering can thereby give highly relevant information about the interface quality as illustrated in the investigations of the initial growth of Cr on Fe [37]. Not only is the epitaxial relation between the substrate and the film established, the thermal vibrations of the individual components can also be extracted. This has been demonstrated using high quality sputter deposited Fe/Cr(001) superlattices by Rüders et al. [23]. One important aspect of the use of nuclear scattering is the isotope selectivity. The cross section is unique for the isotope combination in the scattering process, allowing investigations of e.g. O in oxides by growing isotope layers of O 16 and O 18. This region in the scattering cross section corresponds to scattering above the Coulomb barrier, at which the two particles can be viewed as a compound nucleus in the moment of scattering. For a comprehensive introduction to different nuclear scattering techniques see for example [38] and [39]. The isotope selectivity is also prominent in neutron scattering experiments. Furthermore, the scattering cross section is strongly varying with the atomic number and the choice of isotopes, which makes neutron scattering extremely useful as a complement to x-ray scattering. Neutron reflectivity has been used to determine the composition variation in multilayers and superlattices as well as the magnetic profiles in thin films, multilayers and superlattices. Neutron reflectivity is one of the few methods that allow the full determination of the magnetic structure in materials [40, 41]. Full determination of the magnetic order can be obtained using polarization analysis. This includes the determination of the magnetic ordering in layered magnets as well as stripes and islands. Magneto Optical Kerr Effect (MOKE) is one of the most used ex- and insitu technique to determine changes in magnetization. For a review of the use of MOKE, see e.g. [42]. The MOKE is one of the most versatile approaches to

14 14 B. Hjörvarsson and R. Pentcheva probe the changes in magnetization with temperature, to measure anisotropies as well as the susceptibility. The absolute moment can not be determined by MOKE. Superconducting Quantum Interference Devices (SQUID) are therefore often used for calibrating MOKE results. The magnetic properties of the material can serve as a qualifier with respect to the structural quality. For example, the ordering temperature of extremely thin layers is strongly depending on the thickness. Consequently, if the thickness is changing substantially at lengthscales equal or larger than the magnetic interaction, this would result in a ill defined ordering temperature. In Fig. 1.6, the magnetization of 3.4 ML Fe on GaAs (001) is displayed. The results clearly support the presence of highly uniform layer thickness at all but extremely short length scales and thereby carries information about the uniformity of the grown film. Fig Remanent Kerr rotation (a) and magnetic susceptibility (b) versus temperature for 3.4 ML Fe on GaAs (001) (2 6). Almost no tailing is observed in the magnetization and the susceptibility is narrow and well defined. From [36]

15 1.2 Growth of Metals 1 Modern Growth Problems and Growth Techniques 15 The first step towards fabricating a multilayer or superlattice stack is the growth of a thin film. In this section we discuss the growth of transition metals on a noble metal substrate and of rare earths. The growth of transition metals on other transition metals as well as oxides is covered in Sect Transition Metals on a Noble Metal Substrate Growth of thin magnetic films on a nonmagnetic substrate provides the possibility to design materials that do not exist in the bulk phase: e.g. Co (hcp in bulk) grows fcc on a Cu substrate. Similarly, Fe (bcc in bulk) was inferred to adapt a fcc structure on Cu(001) up to a thickness of 10 ML, while the structure is transformed into a body centered tetragonal phase at thicknesses above 10 ML [43, 44, 45, 46]. Three (0 4 ML, 4 10 ML and > 10 ML) thickness regions of Fe on Cu(001) are identified exhibiting widely different magnetic properties, closely connected with the structure of the films. Recently the local atomic structure of the first iron layers has been revisited and there is still a substantial controversy whether the structure is bct as determined using STM, LEED and DFT-calculations [47, 48, 49] or fcc [50]. The obtained structure is a result of an intricate balance between lattice mismatch, strain and bond strength between adsorbate and substrate, versus adsorbate-adsorbate. We will discuss some specific aspects of heteroepitaxial versus homoepitaxial growth using Co on Cu(001) as an example. Co and Cu are immiscible in the bulk and have only a small lattice mismatch (2%). Still, intermixing influences the obtained interface quality substantially. In the initial stages of metal homoepitaxy, nucleation theory [6] predicts the logarithm of the island density to decreases linearly with increasing temperature (cf. 1.3). Instead of a linear dependence, a complex N-shaped non-arrhenius behavior of the island density (illustrated in Fig. 1.7) is obtained for Co/Cu(001) from an ab initio thermodynamics kinetic Monte Carlo study, using diffusion barriers from DFT [51, 52]. At low temperatures, the heteroepitaxial case resembles homoepitaxial growth with adatoms diffusing on the surface, forming nearly square islands (see STM image at 295 K in Fig. 1.7). At approximately 340 K the activation of atomic exchange leads to a minimum in ln n x and a subsequent increase of island density due to pinning at substitutionally adsorbed Co adatoms [53]. At higher temperatures, the exchange mechanism and diffusion of the substrate material on the surface results in a bimodal island size distribution with large Cu islands decorated with Co as well as a large number of small predominantly Co islands (see STM image at 415 K in Fig. 1.7). Island densities obtained from He-scattering experiments [51] confirm the predicted N-shape of ln n x (1/T ) (cf. Fig. 1.7). This unusual behavior was also observed in STM-measurements [53, 54, 55]. Experimental results for

16 16 B. Hjörvarsson and R. Pentcheva T=415K T=293 K K 300K 200K island density n x [cm -2 ] x x10 13 n tot n Co n Cu /T[K -1 ] Fig N-shaped non-arrhenius behavior of island density in the initial growth of Co on Cu(001) obtained from DFT-kMC simulations (solid diamonds) and ion scattering experiments (solid circles) for F = ML/s [51]. Empty diamonds: theoretical results for F = 0.1 ML/s. Open circles: island densities derived from STM-images (shown above) for F = ML/s [53]. Inset: linear plot of island density between 340 and 410 K for Co (n Co x ) (solid triangles) and Cu (n Cu x )(open triangles) islands. Experimental error bars comprise statistical and, for high and low temperatures, possible systematic errors Fe/Cu(001) [9], Fe/Au(001) [56], Ni/Cu(001) [57], and Co on Ag(001) [58] imply, that the scenario described above could be relevant for a broader class of materials. Epitaxial growth of Co films up to 20 ML has been reported in the literature. However, bilayer growth was observed for the first two layers [4, 54], and the second layer starts to grow before the first one is completed. At elevated temperatures a Cu capping layer is formed [4, 59]. To explain these experimental observations, DFT calculations were performed for different configurations such as monolayers, bilayers and sandwich structures [60]. The corresponding formation energy for 1 ML of Co is shown in Fig The tendency towards

17 1 Modern Growth Problems and Growth Techniques 17 Fig Formation energy of different ferromagnetically ordered configurations for a total cobalt coverage of 1 ML as a function of the cobalt island thickness N. The structures consist of clean Cu(001) and a compact island with N Co layers ( ) or N Co-layers capped by copper. The area covered by the cobalt islands is 1 of N the whole surface. Especially for the copper terminated systems the separation in higher than bilayer cobalt islands is unlikely because of a negligible energy gain the growth of bilayer islands results from the affinity of Co to maximize the number of Co-Co bonds versus Co-substrate bonds. This is also expressed in a strong relaxation (Δd/d Cu = 13.4%) of the interlayer spacing between the two Co-layers. A Co double layer capped by 1 ML of Cu was found to be the thermodynamically most stable configuration. Similar results were found for Co films grown on Cu(111) [61]. A recent molecular dynamics simulation based on Tight Binding Second Moment Approximation (TBSMA) [62] identifies an upward diffusion mechanism at island edges as the origin of bilayer growth [63]. The interaction between ultra thin transition metal films and noble metal substrates is relatively weak. Consequently, the surface strain of such films may be much larger that the one suggested from the bulk phases. For example in the extreme case of a free standing Co monolayer (missing interaction with the substrate) DFT-calculations predict an equilibrium lattice constant 12.2% smaller than the one for Cu, while the lattice constant of bulk fcc Co is only 2% smaller than Cu [60]. Related to this is the so called mesoscopic strain: based on ab initio calculations Stepanyuk et al. [64] found that the shape of the growing islands and the underlying substrate is strongly deformed by the inhomogeneous stress field around Co islands Rare Earth The rare earth (RE) elements (actinides and lanthanides) are extremely reactive and therefore difficult to purify and to maintain impurity free. These can be purchased in a reasonably pure form with respect to all elements but

18 18 B. Hjörvarsson and R. Pentcheva hydrogen. Typically, RE materials are refined in situ, by extensive annealing and outgassing procedures to minimize the hydrogen content. This severely limits the possible deposition techniques and MBE appears to be the technique of choice for successful deposition. Here, we regard Y as representative for RE materials due to the similarities with respect to the outermost electron states, yielding similar chemical properties. The reactivity of the RE materials influences the sample design and most researchers use diffusion barriers to hinder e.g. oxygen transport from the substrate to the film, forming RE oxide. For example, the growth of RE materials on Al 2 O 3 often involves a Nb layer serving simultaneously as a diffusion barrier and a seeding layer, as described by Kwo et al. [65]. This aspect will be discussed further when addressing the growth of RE superlattices. A comprehensive review of the growth of Nb (110) on Al 2 O 3 is found in [67]. The choice of capping layer is important for hindering deterioration of the material. The capping has to wet the RE film and form a stable continuous layer hindering reactions with the ambient atmosphere. Even here, Nb has been used successfully, as Nb forms a self passivating oxide layer at ambient conditions [68]. A good counterexample is the use of gold as a capping layer. At first glance, gold appears to be the ideal material choice for capping, it is inert and it is possible to form what appears to be continuous films. However, Au is highly unsuitable as a capping layer as the RE films deteriorate rapidly[66]. The root of the deterioration is the adsorption of H 2 O on the Au surface. Water diffuses readily on grain boundaries, reaching the underlying film where (a) (b) Yield (Counts/2μC) [H]/[Y] (Atomic Ratio) Sapphire Niobium Mixed! -Y and fcc YH2 YOxHz Gold Energy (MeV) Fig Illustration of the deterioration of a representative film by H 2O. (a) The hydrogen content was determined by nuclear resonance analysis using the N-15 method, and the oxygen content by Rutherford back scattering spectrometry. The hydrogen diffuses deep into the film, while the oxygen forms oxides in the near surface region. Notice the absence of hydrogen in the Nb layers. An illustration of the deduced structure is shown in (b) [66]

19 1 Modern Growth Problems and Growth Techniques Nb YHx Nb Al2O Normalized Yield (counts/μc) H/Y (atomic ratio) Energy (MeV) Fig Illustration of the hydrogen content of a Y film. The Nb capping is hindering the reaction of H 2O with the Y, however, substantial hydrogen concentration is still found in the film it dissociates. This results in the formation of oxides, hydroxides and hydrides. This is illustrated in Fig. 1.9 [66]. Although Nb appears to be an excellent capping material, substantial amount of hydrogen is still found in Nb capped MBE grown RE materials. Typical procedures for the MBE growth of RE materials involves extended (days) outgassing of the target materials, followed by evaporation onto a substrate at elevated temperatures. In-situ capping and subsequent dry oxidation of the Nb capping still results in significant hydrogen content of the grown films. Representative results on the determined hydrogen content are illustrated in Fig [69]. The hydrogen content of Y and RE films are typically in the same range, as expected from the similar chemical properties, as confirmed by measurements on Y, Gd and Ho. The hydrogen content of RE materials is frequently ignored, which seriously influences the reliability of the deduced film properties. 1.3 Growth of Magnetic Oxides and Magnetic Semiconductors While detailed atomistic models for the homoepitaxial growth of metals have been put forward and attempts to incorporate some aspects of heteroepitaxial growth have been made (see discussion in Sect ), a kinetic description of oxide growth is largely lacking. One of the reasons is the complexity of the oxide structures, the different nature of bonding and the variety of chemical species that are involved, leading to a multitude of potentially relevant diffusion processes. Besides the temperature and the deposition rate, the partial pressure of oxygen is an important parameter in the epitaxial growth of oxides.

20 20 B. Hjörvarsson and R. Pentcheva Post-growth treatment (e.g. annealing) in vacuum may lead to reduction of the oxide. Vice versa, post-growth annealing in oxygen atmosphere can help to reduce oxygen vacancies. Concerning the characterization of the grown film, most of the surface science techniques require ultra high vacuum (UHV) conditions. The film properties and structure may be altered by the ambient, which imposes a problem. Furthermore, the insulating nature of oxides hampers the application of imaging techniques, such as scanning tunneling microscopy which requires a reasonably conducting sample. This limitation can be circumvented by using thin oxide films grown on a metal support, providing sufficient conductivity. The surface stoichiometry and structure has important consequences for the reactivity but also for the magnetic and electronic properties of the material. Oxides and their surfaces are typically classified according to electrostatic considerations. The most commonly used are Tasker s scheme [70] and the autocompensation rule [71]. Originating from semiconductor physics and a covalent picture of bonding, the autocompensation rule states that on a stable surface all anion- (cation-) derived dangling bonds have to be filled (empty). Tasker s classification, which emphasizes the ionic nature of bonding, is shown in Fig Here, oxide surfaces are divided in three groups, according to the charge of the layers Q and the dipole moment μ perpendicular to the surface. Systems of type one have neutral layers and no dipole moment perpendicular to the surface (Q =0,μ = 0). Systems of type two and three consist of charged layers. In type two systems the repeat unit has no dipole moment perpendicular to the surface, while in type three it has a nonvanishing dipole moment perpendicular to the surface. It should be noted that depending on the termination, surfaces of the same orientation can be polar or non-polar. For type three surfaces both the scheme of Tasker and the autocompensation rule postulate a diverging surface energy ( polar catasptrophy ) that can only be compensated by strong changes of the surface stoichiometry either by reconstructions or by faceting. Although usefull, these concepts have their limitations, because both models rely on the bonding and valence state in the bulk which may substantially be altered at the surface. Type 1: Q=0, µ=0 Type 2: Q=0, µ=0 Type 3: Q=0, µ=0 Fig Classification of polar oxide surfaces after the scheme of Tasker [70]. See text for details

21 1 Modern Growth Problems and Growth Techniques 21 DFT calculations have shown that lattice relaxations, where electronic charge redistribution often leads to metallization of the surface, can be an effective mechanisms to reduce and even compensate surface polarity (see review of Noguera [72]). Other mechanisms emerging from the correlated nature of transition metal oxides will be discussed in Sect The development of ab initio thermodynamics [73, 74, 75] contributed substantially to identify cases where simple electrostatic arguments fail. In ab initio thermodynamics density functional theory is combined with concepts from thermodynamics to describe the surface stability at ambient pressures and temperatures. The main idea is that the most favorable surface configuration minimizes the surfaces energy. The latter depends on the Gibbs free energy of the system G slab M xo y(001) as well as on the chemical potentials of the constituents. γ(t,p)= 1 2A [ ] G slab M N xo y(001) Mμ M (T,p) N O μ O (T,p). (1.4) When the entropic contributions are small or cancel out, one can susbtitute G slab M xo y(001) with the total energy from DFT-calculations. The chemical potentials μ M (T,p)andμ O (T,p) are not independent of each other. The condition that the surface is not only in equilibrium with the gas reservoir (e.g. oxygen pressure in the atmosphere) but also with the bulk oxide M x O y results in only one independent variable, the oxygen chemical potential which can be translated into partial pressures at a particular (growth) temperature. For further details, see e.g. [75]. A further aspect that has to be noted is that the treatment of transition metal oxides represents a challenge for DFT- methods due to the correlation effects in the d states, localized oxygen orbitals and magnetism. For such systems all electron methods provide the highest accuracy and for the treatment of on-site Coulomb repulsion methods that go beyond the local density approximation (LDA) or the generalized gradient approximation (GGA) like e.g. the LDA+U method [77] are gaining importance. Such methods have mainly been applied to bulk systems and only recently to surfaces and interfaces Binary Oxides In this section we limit the discussion to few examples of the growth of oxides with ferro- or ferrimagnetic coupling. Prominent examples are the halfmetallic ferromagnets, Fe 3 O 4 and CrO 2. Fe 3 O 4 Magnetite is the oldest known magnetic material and its importance ranges from geology to magnetic recording. The predicted half metallic behavior [78] paired with a high Curie temperature makes it a prospective material for spintronics applications. This has generated substantial research activities on magnetite.

22 22 B. Hjörvarsson and R. Pentcheva Magnetite is a ferrimagnet, that crystallizes in the inverse spinel structure where oxygen ions form a slightly distorted fcc lattice. Trivalent iron ions occupy one fourth of the tetrahedral sites (Fe A ), while 50% of the octahedral sites are occupied by mixed valence Fe B -ions. While a p(1 1)-structure has been reported on the (111)-surface, Fe 3 O 4 (001) shows a ( 2 2)R45 - reconstruction. The origin of the latter was subject of a controversial debate in the literature. Various models for a compensated (001)-surface were proposed, where the surface reconstruction was understood as an ordering of surface defects [79, 80, 81, 82, 83, 84]. The surface phase diagram, obtained within the framework of ab initio thermodynamics (shown in Fig. 1.12a) for all possible surface models, revealed that a modified bulk termination yields the lowest energy over the entire range of accessible oxygen pressures. In this so called modified B-layer (top and side view displayed in Fig. 1.12b) the surface reconstruction is a result of a wavelike Jahn-Teller-distortion and not an ordering of surfaces vacancies as in previous models. This termination does not fulfill the electrostatic models and was therefore ignored in the structural analysis so far. Experimental evidence for this structure is obtained from XRD [85], LEED and STM measurements [76]. The wave like pattern in the surface layer is clearly visible in the STM image of the Fe 3 O 4 (001)-surface and the STM simulation in Tersoff-Hamann model using the charge density from the DFT-calculation shown in Fig. 1.12c). Both the DFT and spinpolarized photoemission measurements [76] show that the surface stabilization is accompanied by strong changes in the electronic properties: e.g. a half-metal to metal transition takes place from bulk to the surface. Besides natural samples synthetic single crystals as well as epitaxial films are used in experiment. Koltun et al. [86] grew synthetic Fe 3 O 4 crystals using the floating zone technique from pre-sintered magnetite bars prepared from iron oxalate. After crystallization the samples were annealed at 1473 K for 20hinapartialoxygenpressureof mbar. a) b) c) Fig (a)surface phase diagram of the Fe 3O 4(001)-surface; (b) side and topviewofthemodified B-termination, oxygen, tetrahedral and octahedral iron are marked by white, grey and dark grey circles, respectively; (c) STM image of the Fe 3O 4(001)-surface [76] together with an STM simulation of the modified B- termination both showing the wave like structure in the 110-direction

23 1 Modern Growth Problems and Growth Techniques 23 Epitaxial films were grown on a variety of substrates. Due to the nearly perfect lattice match (0.31%), MgO is an excellent candidate for the growth of epitaxial magnetite films. Fe 3 O 4 -films in the (001)-orientation are typically grown on MgO(001). Other susbtrates used are SrTiO 3 (100) [87](lattice mismatch -7.1%), MgAl 2 O 4 (100) [88] (lattice mismatch 3.8%), and GaAs(100) [89]. The latter substrate is particularly interesting for the incorporation of Fe 3 O 4 in spintronic devices. Two main techniques are used for the synthesis of Fe 3 O 4 : (i) oxidation of Fe films with oxidizing agents as O 2,NO 2 [84] or oxygen plasma; (ii) oxidation during deposition of Fe in an oxygen rich environment. On MgO(100), Fe 3 O 4 -films were grown using O 2 assisted MBE. With a substrate temperature of 525 K and a growth rate of 22.5 Å /min good quality films were obtained [90]. Oxygen plasma assisted MBE was used by Kim et al. to grow Fe 3 O 4 (001) on MgO(001) [92]. The best quality was obtained under oxygen poor conditions (p O2 = mbar) and an iron deposition rate of 0.6 Å/s with an electron-cyclotron resonance (ECR) plasma source runnung at 200 W. Voogt et al. [84] used NO 2 as an oxidizing agent and a similar substrate temperature of 525 K. The growth was monitored by RHEED and the time to form a ML of Fe 3 O 4 (001) was estimated to be 46 s. The threshold of Mg interdiffusion at K represents an upper limit for the growth temperature. For spintronics application it is desirable to combine the half-metallic oxides with semiconductor devices. Lu et al. [91] grew Fe 3 O 4 (001) on GaAs(100): initially a bcc epitaxial Fe-layer was grown on GaAs(100), suppressing the formation of secondry phases (e.g. FeAs) to avoid the development of a magnetically dead interface layer. Subsequently the Fe-film was oxidized at p O2 = mbar. An aspect that needs further investigation is whether the half-metallic behavior is preserved at the interface to MgO(001) or GaAs(001). For the growth of Fe 3 O 4 (111), different substrates have been used, e.g. MgO (111) [94, 95] and Al 2 O 3 [96, 97], as well as metallic Pt(111). Weiss et al. [98] repeatedly deposited and oxidized layers of iron on Pt(111). The LEED analysis of well ordered films suggested a termination with 1/4 monolayer of Fe over a distorted hexagonal oxygen layer. Dedkov et al. oxidized a Fe(110) film grown on W(110) [99]. Depending on the oxygen pressure and the post growth annealing procedures, lattice parameters corresponding to the formation of FeO and Fe 3 O 4 (111) were obtained. A FeO(111)-surface was formed after 100 L oxygen exposure and post-annealing at 525 K. The FeO-film was transformed into a Fe 3 O 4 (111) film after subsequent exposure to 200 L oxygen and post-annealing at 525 K. Alternatively, Fe 3 O 4 (111) was obtained after an extended exposure to 900 L oxygen. Fonin et al. [93] grew Fe 3 O 4 (111) by oxidizing a Fe(110)-film grown on Mo(110)/Al 2 O 3 (11 20) at 700 Cand p O2 = mbar. A TEM cross section of the film demonstrating the four different regions of the sample with sharp interfaces is shown in Fig. 1.13a). A lower spin-polarization of 60% was measured for this sample, as compared to the nearly fully spin-polarized (80%) one grown on W(110) [99]. This result

24 24 B. Hjörvarsson and R. Pentcheva is consistent with DFT calculations [100] of bulk magnetite showing that uniaxial strain reduces the degree of spin-polarization. A 100 ÅthickFe 3 O 4 (111)-film was grown on Al 2 O 3 (0001) [97] by codeposition of Fe from an effusion cell and atomic O using a plasma source, at a substrate temperature of 450 K and postannealing at 900 K. Such interfaces are interesting as magnetic tunneling junctions (MTJ). For Fe 3 O 4 (111) films grownonan(1 1)-OH terminated MgO(111)-surface [94] and Al 2 O 3 (0001) [96] a phase separation in Fe and FeO nanoinclusions were observed at the interface. Unlike MgO(100), the MgO(111)-surface is polar. Substrate polarity was identified as the driving mechanism towards phase separation at the interface. In regions between the Fe crystalites atomically abrupt interfaces were observed. A HRTEM image is shown in Fig DFT-GGA calculations [95] find that these are stabilized through electronic screening and metallization at the interface in contrast to the stoichiometry change expected from classical electrostatic models. CrO 2 Since the prediction of half metallic behavior of CrO 2 [101], this oxide has attracted attention as a potential material for spintronic devices. Unfortunately, rutile CrO 2 is unstable at room temperature and transforms irreversibly into a) b) MgO Fig TEM cross section micrograph of a) the Fe 3O 4(111)/Fe(110/Mo(110) /Al 2O 3(11 20) system [93] with sharp interfaces and b) Fe 3O 4(111) grown on MgO(111) showing the formation of Fe(110) nanocrystals at the interface [94]

25 1 Modern Growth Problems and Growth Techniques 25 Cr 2 O 3, which is an insulating antiferromagnet. In order to stabilize the rutile structure, epitaxial films were grown on TiO 2 (001) and Al 2 O 3 (0001) using chemical vapor deposition CVD [102]. On Al 2 O 3 (0001) a 400 ÅCr 2 O 3 -layer is formed prior to the growth of CrO 2 (001). A 1000 ÅthickCrO 2 (001) film of good crystal quality was obtained on TiO 2 with no Cr 2 O 3 formation. The magnetic ordering temperature of this film was 385 K. Heterostructures based on the insulating TiO 2 as a barrier are interesting for magnetic tunnel junctions. On the other hand, metallic RuO 2 -spacers are interesting as GMR-elements. CrO 2 (a =4.421 Å, c =2.916 Å) and RuO 2 (a =4.499 Å, c =3.107 Å) have also a good lattice match. Upon deposition of RuO 2 on CrO 2 /TiO 2 the Cr 2 O 3 termination of the surface is transformed to CrO 2 but despite the restored conductivity a relatively low magnetoresistence of the sample indicates susbtantial chemical and magnetic disorder associated with this transformation [103]. Besides TiO 2 and RuO 2, also SnO 2 was used as a substrate, however the measured magnetoresistance was still relatively low. w-nio Analogous to the growth of ferromagnetic materials on a non-magnetic substrate, an attractive possibility is opened by the growth of oxides on semiconductors where the oxide adopts the structure of the substrate, which does not exist in the bulk. Recently, Wu et al. [104] predicted, using LDA+U calculations, that NiO in the wurzite structure (w-nio) should be halfmetallic and on substrates like ZnO or GaN the ferromagnetic coupling should be more stable than the antiferromagnetic Ferromagnetic Semiconductors The design of semiconductors with magnetic and/or spin-related properties and high Curie temperature for spintronic devices is a demanding and by far not completely resolved issue. In this subsection we will briefly summarize the current knowledge on one of the most intensively studied III-V semiconductors, Mn doped GaAs (T C 170 K), and then discuss several systems where the prediction of room temperature ferromagnetism (e.g. Co:TiO 2, p- type Mn:ZnO, Mn:GaN) has envigorated a lot of research in the last years. A major issue in the fabrication of doped semiconductors is the incorporation of dopants in the lattice and whether a homogeneous distribution can be achieved. Mn:GaAs Unlike II-VI semiconductors where there is no solubility limit for 3d dopands, the solubility limit is very low for III-V systems (e.g. 0.1% for Mn). Therefore a secondary phase like MnAs is formed under typical growth conditions. To avoid this, MBE growth is performed at low temperatures (LT-MBE) of K. The Mn ions primarily occupy cation sites (Mn Ga )wheretheyactas

26 26 B. Hjörvarsson and R. Pentcheva acceptors introducing both magnetic moments and holes. The current understanding of the underlying mechanism is that of hole-mediated ferromagnetism in p-type materials. Since the growth proceeds at highly non-equilibrium conditions where kinetic effects dominate, there are two main types of defects that form and influence the magnetic properties and conductivity: the As antisites (As Ga ) and the Mn-interstitials (Mn I ). Both act as donors, i.e. reduce the number of holes. Additionally, Mn I couple antiferromagnetically to Mn Ga and thus suppress ferromagnetism. Yu et al. [105] provided evidence that the reduction of T C is directly related to formation of Mn I and the latter depends on the doping of the barrier layer on which Mn:GaAs is deposited. Therefore the reduction of both types ofdefectsisthemainpathtowards obtaining a higher T C. This is done through a control of the As flux during growth and a post-deposition low temperature ( K) annealing step. As shown by polarized neutron reflectometry on as grown and annealed samples, annealing improves substantially the homogeneity of Mn and thus the magnetic properties [106]. For further reading the reader is referred to several reviews e.g. [107, 108, 109]. Doped Thin Film Oxides Co:TiO 2 : The interplay of growth conditions (temperature, deposition rate), but also the partial pressure of oxygen plays a decisive role on the quality of the samples. For example PLD growth of Co:TiO 2 (anatase) from a mixed metal oxide target may result in Co -nanoinclusions within the TiO 2 matrix if the O 2 -pressure is too low or if the Co:Ti ratio is too high [110, 111, 112]. A continuous epitaxial film with no signs for Co enrichment was obtained for depostion rates of 0.1 Å/s and T = 650 C. On the other hand OPA-MBE material is found to be FM at RT for x Co 5 7% (1.1 μ B /Co) [113]. Similar morphologies but a lower saturation magnetization of 0.6 μ B /Co is obtained for Ar-sputtering of Ti and Co metal targets at T 650 Cusingwaterasan oxidant [114]. The incorporation of hydrogen could be of importance in this context, as discussed in Sect above. Another issue, similar to the ones arising at oxide interfaces, is the mechanism of charge compensation: e.g. in Co:TiO 2 Co 2+ substitutes for Ti 4+. DFT-GGA calculations predict Co 2+ -segregation together with the formation of an oxygen vacancy [115]. Annealing in vacuum is not likely to lead to Co-oxidation. For example, XANES measurements revealed metallic Co at T>750 C [116]. Post-growth annealing results in enrichment of Co at surfaces, grain boundaries and interfaces. Also the origin of room temperature ferromagnetism is not well understood, e.g. in Co:TiO 2 crystallographically perfect samples were obtained for T 550 CandR = Å/s which turned out to be nearly nonmagnetic [117]. This leads to the assumption that defect formation at surfaces and interfaces plays a significant role in triggering magnetism.

27 1 Modern Growth Problems and Growth Techniques 27 Cr:TiO 2 : PLD growth of Cr: TiO 2 (rutile) on Al 2 O 3 (012) resulted in conducting films (reflecting a finite density of O-vacancies) and a saturation moment of 2.9 μ B /Co for x =0.07. OPAMBE films grown on TiO 2 (110) were on the other hand insulating and nonmagnetic. XPS showed that Cr 3+ substitutes for Ti 4+. Post growth annealing in vacuum reduced the film and made it n-type, showing weak room temperature ferromagnetism with 1 μ B /Co. Cr:TiO 2 (anatase): LaAlO 3 or SrTiO 3 were chosen as substrates for OPAMBE growth. The in-plane lattice mismatch between anatase and the perovskite structure along (001)-direction differs by an order of magnitude for LaAlO 3 and SrTiO 3 ( 0.26 vs. 3.1%), respectively. For R 0.1 Å/s and T 650 C the as grown films exhibited room temperature ferromagnetism which was enhanced after annealing in vacuum. The films exhibit a high Curie temperature (T c =690 K). An improved crystalline quality is obtained for R Å/s and T 550 C [117] but again as for Co:TiO 2 magnetic properties deteriorate. In summary, room temperature ferromagnetism in Co or Cr:TiO 2 appears to be driven by defects and is not an intrinsic property of the material. Co:ZnO: The appearance of room temperature ferromagnetism in Co:ZnO depends critically on electron doping. Epitaxial films of Co:ZnO grown on Al 2 O 3 (012) by MOCVD [118] were found paramagnetic and insulating. However, the interdiffusion of atomic Zn in these samples (Zn occupies interstitial sites) results in a weakly ferromagnetic semiconducting sample [119]. Ti:Fe 2 O 3 : A nontraditional candidate for room temperature ferromagnetism is Ti doped Fe 2 O 3. The host is a (canted) antiferromagnetic insulator at room temperature, but the incorporation of Ti in the lattice is expected to lead to ferrimagnetism. Chambers and collaborators [116, 120] grew a 710 Å α-fe 2 O 3 -film on Al 2 O 3 using a 130 Å thick buffer layer of Cr 2 O 3 to reduce the in-plance lattice mismatch (5.8%). The OPAMBE growth was performed at T sub = 550 C with growth rate of 0.25 Å/s at an oxygen pressure of mbar. XRD measurements suggested a high degree of crystallinity. The magnetic signal depends sensitively on whether Ti is incorporated on one spin sublattice or randomly distributed on both spin sublattices. The measured saturation magnetization is however much lower (approximately 0.5 μ B /Ti) than the expected 5 μ B /Ti indicating that only about 12% of Ti contributes to a FM ordered phase and the the majority of Ti is randomly distributed in the lattice. 1.4 Multilayers and Superlattices General Considerations A multilayer is a general term describing one dimensional variation in composition. A multilayer can be single crystalline, polycrystalline, amorphous or a combination of (poly-)crystalline and amorphous. When a multilayer is single crystalline and has many repetitions, it is denoted superlattice, see Fig

28 28 B. Hjörvarsson and R. Pentcheva L c Capping :! c L SL L A Seed layer Substrate L B Fig Illustration of a typical superlattice structure. A seed layer is grown on a substrate, followed by the growth of the superlattice. The structure is thereafter typically covered by a capping layer, hindering the deterioration of the superlattice A superlattice can therefore be thought of as a single crystal multilayer, with well defined atomic distances as well as chemical repeat distance (Λ). Λ defines the unit cell in the growth direction, where Λ=L a +L b where L a and L b are the thicknesses of layer a and b, respectively. The variation in the chemical composition is a route to create a modulation in the electronic states, forming new material classes with unique properties. However, formation of a superlattice is only the first step. Altering the chemical composition in 3 dimensions would allow much larger degree of freedom, forming unique electronic states defined by the extension and the compositional variation in the material. This can be viewed as the ultimate task of materials processing of today. The growth of superlattices bears large similarities to the growth of single films. However, there are also significant differences both with respect to the growth procedures as well as the analysis of the samples. The inherent lattice parameters of the constituents have a special role, which is often used to judge the possibility for the growth of high quality superattices. The basic ideas resemble in many ways the criteria for the growth of single films on a substrate, as will be apparent below Metallic Superlattices We will use the combination of Fe, Mo and V as examples for the possibilities and limitations of the growth of metallic superlattices. All these elements are bcc with a bulk lattice parameter of 2.86, 3.16 and 3.02 Å, respectively and can therefore, in principle, form congruent single crystals. The difference in the lattice parameter of Fe and V is 5%, Mo and V is 4% and finally Fe and

29 1 Modern Growth Problems and Growth Techniques 29 Mo it is close to 10%. The initial growth of V on Fe has been investigated using RHEED [121] following the in-plane lattice parameter changes of the V, in a wide temperature range. A 200 ML Fe (001) layer was initially grown on a MgO(001) substrate, on which the V was deposited. Between 300 and 800 K, RHEED oscillations were observed up to 9 ML, consistent with a well defined layer by layer growth of V. However, the in-plane lattice parameter at the V surface was observed to relax from that of Fe when the number of V layers exceeded 7 ML. At 800 K, no RHEED oscillations were observed, which indicates the presence of strong island formation and an upper boundary for a well defined growth. Thus, it is possible to grow V on Fe up to 7 ML, although the difference of the lattice parameter is as large as 5%. When the number of repeats of the superlattice period is large, the average in-plane lattice parameter reflects the thickness ratio of the constituents. Thus, when growing e.g. Fe/V superlattices, the average in-plane lattice parameter can be estimated directly from the thickness ratio, as the elastic constants of Fe and V are rather similar. For example, when the layers have equal thicknesses (L Fe =L V ), the average in-plane lattice parameter will be close to the average lattice parameter of the constituents. This reduces the in-plane strain from 5 to below 3%, which should be reflected in an increased critical thickness for the growth of coherent layers. This was observed in RHEED investigations of Fe(3)/V(x) (001) superlattices in which the critical layer thickness was determined to be around 16 ML [122]. The growth of Fe/V(001) superlattices has been thoroughly investigated by a number of authors (see for example [123, 124]). The temperature dependence of the growth on MgO (001) was established by Isberg et al. [124] where the quality of the SL were shown to depend strongly on the growth temperature. The best crystalline quality was obtained with a substrate temperature in the temperature range K. The thickness variation of the layers was also at its best in the same range and was determined to be 1 Å. This variation in layer thickness represents the lower limit for a non phase locked growth, resulting in incomplete formation of the individual Fe and V layers. Thickness variation corresponding to one monolayer of both the Fe and V layers is thus inevitable. Representative X-ray results are shown in Figs and The optimal growth temperature of Mo/V superlattices is close to 1000 K, which is substantially higher than for Fe/V superlattices. This correlates with the substantially higher melting temperature of Mo compared to Fe. The growth of this material combination was pioneered by the group of Fisher [125, 126, 127], demonstrating both a (001) and (110) growth on MgO(001) and Al 2 O 3 (11 20) respectively. The influence of the critical thickness of the layers in Mo/V superlattices on the superconducting properties was discovered by Karkut et al. [125]. A critical thickness of 16 ML was inferred for both the (001) and (110) grown Mo/V superlattices, when the ratio of the layer thicknesses was close to unity. The difference of the lattice parameters of Mo and Fe with respect to V is similar, but with different sign. Thus, a comparable critical thickness for epitaxial growth is observed for compressive and tensile

30 30 B. Hjörvarsson and R. Pentcheva Fig Representative reflectivity data from a Fe/V(001) superlattice with Λ = 25.1 Å. From [122] biaxial strain in V, as expected from symmetry reasons. A compressive strain in one layer is balanced by a tensile strain in the second. Thus, the sign of the strain appears to be irrelevant, and a hint of the relation between the lattice mismatch and the critical thickness of the layers emerges. Above the critical thickness, the formation of dislocations and other defects results in buckling and increases therefore the variation in the layer thicknesses [126, 127, 128]. This leads to relaxation of the in plane lattice parameter, where the variation is accomplished by the presence of defects. The thickness variation is governed by the V layers, because V atoms have larger surface mobility at the actual growth temperature [128]. Fig Representative diffraction data from a Fe/V(001) superlattice with Λ = 25.1 Å The inset highlights the crystalline quality of the superlattice structure. The presence of Laue oscillations implies a interference between the scattering from the first and the last monolayer of the Fe/V(001) stack. From [122]

31 1 Modern Growth Problems and Growth Techniques 31 Let us now consider the growth of Fe on single crystal Mo(110), with a lattice mismatch of close to 10%. The growth results in a complicated defect generation already in the first monolayer. The first layer grows pseudomorphically [129] followed by a pronounced relaxation in the third layer. Thus, the difference in the lattice parameters of Fe and Mo appear to be beyond the limit for epitaxial growth. The growth and characterization of Fe/Mo multilayers is described in the literature [130, 131], but no reports are found on succesful growth of Fe/Mo superlattices. The choice of growth temperature is a compromise between two constraints, surface mobility and interdiffusion. Surface mobility of adatoms is increased with increased temperature, but so is the interdiffusion. This limitation is clearly seen in many of the transition metal superlattices and constitutes therefore a substantial challenge for optimizing the growth. One route to circumvent this limitation is the use of surfactants. The basic idea is to decrease the activation energy of the surface diffusion and thereby increasing the surface mobility. As seen in (1.2) the weight of the change in activation energy (E d ) has the same influence on the diffusion rate (D) as the change in temperature. One possible surfactant is hydrogen. The influence of hydrogen on the diffusion of Pt adatoms on Pt(111) was investigated by STM [133]. A clear increase of the diffusion rate was demonstrated. However, the presence of hydrogen has also been found to inhibit surface mobility [134]. The main benefit of the use of hydrogen in this context, is the compatibility with the vacuum processes. Hydrogen is simply removed by evacuation from the deposition chamber. The use of hydrogen as a surfactant was recently demonstrated by Remhof et al. [132], where a substantial increase of the quality of Fe/V(001) superlattices was obtained. Representative x-ray reflectivity results are illustrated in Fig Co/Cu Superlattices of Co/Cu have been widely studied during the last decades. Successful growth of an epitaxial Co/Cu on a single crystal Cu(001) under UHV conditions was demonstrated by Cebollada et al. [135]. The Cu substrates were cut from a single crystal bar and oriented within 0.3 of the [001] direction using Laue diffraction. The substrates were cleaned in-situ by cycles of Ar+ sputtering and annealing to 1000 K. The composition of the surface was investigated using Auger electron spectroscopy and the crystalline quality was investigated by thermal-energy atom scattering (TEAS). The resulting surface consisted of, on average, 300 Å wide flat terraces separated by monoatomic steps. The authors used extremely low deposition rate, 0.01 Å/s as compared to few Å/s, to reduce the amount of imperfections. The thickness of the evolving layers was monitored and determined by counting the number of oscillations in the TEAS signal. The thickness obtained this way was confirmed by calibrated quartz balance. The samples were all covered by

32 32 B. Hjörvarsson and R. Pentcheva Fig Small angle x-ray reflectivity scans recorded at E = kev. The upper curve displays the reflectivity of the sample sputter deposited at p H2 = mbar at T = 320 C. The lower curve shows the reflectivity of a reference sample grown without the presence of hydrogen. The scans are shifted along the intensity axis for clarity. The inset shows the diffuse scattering recorded at the first superlattice peak [132] 1000 Å Cu, to hinder oxidation of the underlying superlattice. The quality of the resulting structure was established by neutron reflectivity and diffraction. Concerning the growth of Co/Cu superlattices, we tie up to discussion of the initial growth of Co on Cu(001) in Sect The lattice parameter of fcc Co at room temperature is Å, while the lattice parameter of Cu is Å under the same conditions. The lattice mismatch is therefore slightly below 2%. When growing thin Co layers on Cu(001) single crystal, the Co adapts the in-plane lattice parameter of Cu. This results in a tetragonal distortion of the fcc lattice, where the out-of-plane lattice parameter is contracted and the in plane lattice parameter is expanded, as compared to the bulk value [57]. The restoring force at the interface between the Co and the Cu substrate is fixed but the strain energy associated to the tetragonal distortion increases linearly with the thickness of the Co film. Thus, at some critical thickness, it will be energetically more favorable to form defects, relieving the strain and forming a non distorted fcc lattice. In the LEED study of Navas et al. [136] Co was found to grow in the strained state up to at least 10 monolayers. The inherent electronic structure of the Co layers is therefore significantly different from bulk Co. The relation between the structural and magnetic properties of single Co films and Co/Cu superlattices is discussed by de Miquel et al. [137]. The subsequent growth of Cu on the thin Co films retained the bulk lattice parameter of Cu. Thus, while growing Co/Cu(001) superlattices on a Cu(001)

33 1 Modern Growth Problems and Growth Techniques 33 substrate, only the Co layer should exhibit a tetragonal distortion, as long as the superlattice is grown in a coherent mode. Fe/Cr The work on Fe/Cr superlattices was pioneered using evaporation techniques. For example, Etienne and collaborators used effusion cells for the growth of both Fe and Cr on GaAs under UHV conditions [138]. The surface structure of the substrate was improved by growing a Å GaAs(001) buffer at 1000 K. The surface reconstructed from 2 4to2 6 after a brief Ga-exposure at 725 K. The metallic superlattices were grown on such buffer layers, starting with Fe. The growth temperatures of the superlattice were in the range K, the drift reflecting the thermal balance with the cooling and heating devices in the MBE system. The presence of layering was confirmed by a combination of Auger spectroscopy and sputtering and the surface crystallinity was investigated using RHEED. The chemical purity was investigated by Auger spectroscopy and the samples were found to be free of both oxygen and carbon. The authors did not present any x-ray data. The resulting giant magnetoresistance (GMR) of these structures was in the range of 100%, taking the high field resistance as a reference. Epitaxial growth of Fe/Cr(001) by sputtering on MgO(001) was later demonstrated by Fullerton et al. [139]. In this case, the authors used a Cr buffer layer grown at 873 K which leads to an improved surface flatness and wetting of the initial layers of the superlattice. The bcc Cr is rotated with respect to the MgO substrate in the same way as discussed for Fe and V on MgO. The Fe/Cr superlattice was grown at 450 K and the resulting structure was investigated by x-ray reflectivity and diffraction. Typical rocking curves of the (002) diffraction peak yielded a FWHM of 0.7 and a FWHM of 0.2 in 2θ, usingcuk α radiation. The authors did not dwell on the choice of growth temperatures. This approach resulted in a 150% GMR, which is substantially higher than obtained by Etienne and collaborators. The structural quality is of large importance for the resulting physical properties, such as GMR. This is clearly seen when comparing the results from single crystal structures with polycrystalline samples. Parkin and York [140] grew polycrystalline Fe/Cr multilayers by sputtering using Si substrates. The structure was a combination of (110) and (001) textured crystallites, of which the (001) increased in weight with increasing temperature in the range 300 K to 475 K. These samples yielded a maximum 40% GMR effect and the resistivity of the optimized samples was 4.4 μω cm at 4.2 K, as compared with 14 μω cm for single crystal superlattices [139]. The growth of Fe/Cr(001) on SrTiO 3 using evaporation, was discussed by Ono and Shinjo [141]. The crystalline quality was substantially worse as compared with the results of Fullerton et al. [139], although great care was taken concerning the surface flatness of the substrates. Chemical etching was used to obtain better flatness and in-situ RHEED was used to investigate

34 34 B. Hjörvarsson and R. Pentcheva the substrate as well as the surface quality of the resulting film. A strong influence of the substrate quality on the GMR was obtained, the chemically etched substrate yielded twice as large values for GMR. The x-ray contrast between Fe and Cr is poor due to the similarity in the electron density of these elements. This was a major obstacle for obtaining higher order Fourier components in, for example, x-ray reflectivity measurements. Bai et al. overcame this difficulty using resonant scattering techniques [142], allowing detailed simulations of the composition profile as well as the roughness of the samples. The experiments were performed on sputtered superlattices, using MgO(001) as a substrate. Rare Earth Superlattices The main part of the magnetic moment in Rare Earth materials (RE, in short for actinides and lanthanides) stems from the localized f-electrons. In contrast to itinerant magnets, due to this localization one can often safely ignore the effect of hybridization on the magnetic moment. Kwo et al. [65] demonstrated the growth of high quality Gd/Y superlattices and related the changes in moment and ordering temperature to the repeat distance in the samples. The growth of the superlattices was done in an UHV MBE system, with a base pressure in the mbar range. A buffer layer of Nb (011) was grown on an Al 2 O 3 (11 20) at 1170 K, hindering the transport of oxygen from the substrate to the RE film. A seeding layer of Y(0001) was subsequently grown at 970 K, resulting in a flat surface, enabling the growth of high quality Gd/Y(0001) superlattices at 470 K. Lower growth temperature is chosen for the superlattice to limit the interdiffusion of the constituents. An alloy interface region of 2 monolayers was established by combined x-ray and magnetic analysis. The interface region was inferred to be a GdY alloy, without magnetic ordering, even at low temperatures. The magnetic susceptibility is large, which confirms the proposed model of the compositional modulation. The growth and characterization of other RE superlattices was established by a number of groups most of which followed the ideas of Kwo et al. with respect to diffusion barrier for oxygen and a seeding layer. McMorrow and collaborators investigated the chemical structure of Ho/Lu and Ho/Y superlattices [143] using high resolution x-ray scattering to address the nature of the interface imperfections. By detailed investigations of the shape of the superlattice Bragg peaks, the presence of conformal roughness of the interfaces was established. The basic idea behind these investigations was to establish not only the type of roughness, but to separate roughness and intermixing. This is of primary interest as the physical properties are highly dependent on the type of interface imperfections. Intermixing can for example result in an absence of ferromagnetic ordering as discussed above [65], while both conformal and uncorrelated roughness will give rise to local anisotropy fields. A thorough description of the growth and structural characterization of rare earth superlattices is given by Majkrzak et al. [144].

35 1 Modern Growth Problems and Growth Techniques Metal-Oxide Superlattices and Magnetic Tunnel Junctions Magnetic tunnel junctions containing an oxide barrier sandwiched between two ferromagnetic layers are a prototypic system to achieve high tunnel magnetoresistance (TMR) values. In this subsection we discuss several issues concerning the quality of the interface that have impact on the measured TMR value. For structurally perfect interfaces a TMR value of several 1000% was predicted theoretically for Fe/MgO/Fe(001) [145]. Experimentally, TMR values up to 188% at room temperature (RT) were achieved [146] for a junction where Fe was grown using MBE. The MgO barrier was epitaxially grown using electron-beam evaporation of a stoichiometric source material. To avoid the oxdation of the bottom Fe layer the first MgO layer was deposited at p O2 = mbar. The top Fe electrode was deposited at 473 K substrate temperature. A measured asymmetric current-voltage characteristic was attributed to an assymmetry of the interface structure. This is rationalized by thermodynamic arguments (cf. (1.1): the lower surface energy of MgO (1.1 J/m 2 ) versus Fe (2.9 J/m 2 )) suggests layer-by-layer growth of MgO on Fe but not vice versa. Tusche et al. [147] found a substantial effect of the oxygen atmosphere on the quality of the interface. Starting with 2 MLs of MgO deposited on Fe(001) at a rate of R=0.125 ML/min by electron bombardment of a polycrystalline MgO rod under UHV conditions they deposited 8 ML of Fe using MBE at R=0.25 ML/min. In the sample where the Fe film was deposited at UHV conditions, 30% of the interface layer was found to be FeO with a subsequent disordered Fe layer. In a second sample where the initial 0.5 ML Fe were deposited in oxygen pressure of 10 7 mbar, surface x-ray diffraction (SXRD) measurements and quantitative analysis revealed that a coherent Fe layer is formed attributed to the formation of a nearly complete FeO layer between the MgO spacer and the Fe film. In a thermodynamic picture, the role of the FeO layer is understood as to reduce the interface energy (cf. (1.1)). DFT calculations found strong dependence of the transport properties on the structure of the interface [148]. Only a symmetric Fe/FeO/MgO/FeO/Fe junction was predicted to give rise to a giant TMR [147]. Parkin and coworkers measured 220% TMR at RT in a FeCo(001)/ MgO(001)/(Fe 70 Co 30 ) 80 B 20 sample where the bottom electrode was a polycrystalline bcc FeCo-layer with a (001)-texture [149]. Djayaprawira et al. [150] found a significantly higher RT TMR using amorphous CoFeB electrodes as compared to polycrystalline CoFe (values of TMR were 230% vs. 62%). In this experiment the metal reference and the free layer were deposited by dc magnetron sputtering, while rf sputtering was used for the MgO film. After annealing at 593 K a partial crystallization at the interface was observed in HRTEM. Microstructural analysis of the MgO(001) spacer shows a good crystal quality with a fibre texture. The role of the recrystallization after annealing and the distribution of B is still not well understood. First principles calculations find a preference for B to reside at the interface [151], which is expected to suppress the TMR.

36 36 B. Hjörvarsson and R. Pentcheva The authors conclude that inhibiting B segregation at the interface during processing is likely to enhance TMR. X-ray photoemission studies [152] on CoFeB/MgO bilayers find evidence for CoFeB oxidation during MgO deposition, while annealing in vacuum leads to B interdiffusion into MgO and MgB x O y formation. To avoid this a Mg-buffer layer is introduced between CoFeB and MgO. A Co/MgO/Co MTJ where Co is stabilized in the bcc structure was predicted from DFT to have an even higher MR than Fe/MgO/Fe [153]. Indeed, MBE grown Co-based MTJ showed a MR of 410% [154]. To retain the metastable bcc structure the thickness of the Co layers on both sides of the insulating MgO barrier were limited to 4 ML and grown at RT. The authors reported that Co does not wet the MgO(001) spacer but grows in a 3D manner. The subsequent annealing step at 525 K for 30 min was used to improve the crystalinity of the sample. Oxide Superlattices Transition metal oxide superlattices open new possibilities to make artificial materials with magnetic and electronic properties that differ from the bulk components. Analogous to oxide surfaces (cf. Sect ) the question of polarity and disruption of charge neutrality arises also at oxide/oxide interfaces. For example perovskites possess a natural charge modulation in the [001]- direction, e.g. in LaTiO 3 positively charged (LaO) + alternate with negatively charged (TiO 2 ), while in SrTiO 3 both the SrO and TiO 2 -layers are neutral. Thus the interface between these two insulators represents a simple realization of a polar discontinuity. Using PLD from a single crystal STO target and a polycrystalline La 2 Ti 2 O 7, Ohtomo et al. [155] fabricated superlattices of LaTiO 3 and SrTiO 3 with an atomically controlled number of layers of each material which therefore are refered to as digital. The films were grown at 970 K at an oxygen pressure of 10 5 mbar to stabilize both valence states of Ti and subsequently annealed at 670 K to fill residual oxygen vacancies. An anular dark field TEM image is shown in Fig RHEED oscillations were used to monitor growth. Ohtomo et al. found that although the parent compounds are a Mott and a band insulator respectively, the heterostructure is conducting with electron energy loss spectra suggesting mixed Ti-valence in the interface region. Based on Hubbard models, Okamoto and Millis [156] proposed an electronic reconstructuion of this interface. Another system showing unexpected behavior are heterostructures of the two simple band insulators LaAlO 3 and SrTiO 3. Here, both the A and B sublattice cations in the perovskite structure change across the interface giving rise to two different types of interfaces: an n-type between a LaO and a TiO 2 - layer that was found conducting with a high electron mobility and a p-type betweenasroandanalo 2 -layer that showed insulating behavior despite the charge mismatch [157]. Using PLD, the n-type LAO/STO IF was grown

37 1 Modern Growth Problems and Growth Techniques 37 a) b) Ti 3+ Ti 4+ Fig a) an anular dark field STEM image of bright LTO layers in a STO host. The boosted up view above shows a 1 5 LTO/STO superlattice [155]; b) 45 checkerboard charge density distribution of the occupied 3d states in the charge and orbitally ordered TiO 2 layer at the LTO/STO-IF. The positions of O-, Ti 3+ and Ti 4+ -ions are marked by white, black and grey circles, respectively [159] on a TiO 2 -terminated SrTiO 3 -substrate. To grow a p-type LAO/STO interface a SrO-layer was deposited on the SrTiO 3 -substrate prior to growth of LaAlO 3. The oxygen pressure is quite an important growth parameter in these systems that controls the oxygen stoichiometry and the underlying properties, e.g. STO alone can change from a wide band insulator to a metal with p O2 [20]. Nakagawa, Hwang and Muller [158] discussed recently that the p-type LAO/STO IF is ionically stabilized with an enhanced roughness attributed to oxygen vacancies while the n-type interface is electronically stabilized and hence sharp. In correlated materials with multivalent ions correlation driven charge order offers an additional degree of freedom to accommodate the charge imbalance. To this end LDA+U calculations predict a compensation mechanism by a charge disproportionation: a charge and orbitally ordered IF-layer is found for the LTO/STO and the n-type LAO/STO interface with Ti 3+ and Ti 4+ ordered in a checkerboard manner [159, 160]. Such a correlation driven compensation mechanism is not present e.g. at polar semiconductor interfaces. Moreover, although both LaAlO 3 and SrTiO 3 are nonmagnetic and LaTiO 3 is an antiferromagnet of G-type, the diluted layer of Ti 3+ in the IF layer has a slight preference to couple antiferromagnetically with a magnetic moment of 0.71 μ B. Recent experiments give first indication for localized magnetic moments at the n-type LAO/STO interface [161]. Thus the violation of charge neutrality at interfaces of transition metal oxides can be used to generate novel charge and magnetically ordered phases that do not exist in the bulk.