synthesis and characterization of ternary carbide thin films170841/fulltext01.pdf · 2009. 2....

ACTA

UNIVERSITATIS

UPSALIENSIS

UPPSALA

2007

Digital Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 353

Synthesis and Characterization ofTernary Carbide Thin Films

OLA WILHELMSSON

ISSN 1651-6214ISBN 978-91-554-6991-7urn:nbn:se:uu:diva-8265

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List of publications

I. Deposition and Characterization of Ternary Thin Films within the Ti-Al-C System by DC Magnetron Sputtering

O. Wilhelmsson, J.-P. Palmquist, E. Lewin, J. Emmerlich, P. Eklund, P.O.Å. Persson, H. Högberg, S. Li, R. Ahuja, O. Eriksson, L. Hultman and U. Jansson J. Cryst. Growth 291, 290 (2006) II. Structural, Electrical, and Mechanical Characterization of Magnetron

Sputtered V-Ge-C Thin Films O. Wilhelmsson, P. Eklund, H. Högberg, L. Hultman and U. Jansson Submitted to J. Mater. Res. III. Electronic Structure and Chemical Bonding in Ti2AlC Investigated by Soft

X-ray Emission Spectroscopy M. Magnuson, O. Wilhelmsson, J.-P. Palmquist, U. Jansson, M. Mattesini, R. Ahuja and O. Eriksson Phys. Rev. B 74, 195108 (2006) IV. Intrusion-type Deformation in Epitaxial Ti3SiC2/TiC0.67 Nanolaminates O. Wilhelmsson, P. Eklund, F. Giulian, H. Högberg, L. Hultman and U. Jansson Appl. Phys. Lett. 91, 123124 (2007) V. Design of Nanocomposite Low-Friction Coatings O. Wilhelmsson, M. Råsander, M. Carlsson, E. Lewin, B. Sanyal, U. Wiklund, O. Eriksson and U. Jansson Advanced Funtional Materials 17, 1611 (2007) VI. Tribofilm Formation and Tribological Properties of TiC and Nanocompo-

site TiAlC Coatings M. Lindquist, O. Wilhelmsson, U. Jansson and U. Wiklund Submitted to Wear VII. Synthesis and Characterization of (Ti,Fe)C Films Deposited by DC Magne-

tron Sputtering O. Wilhelmsson, S. Bijelovic, M. Linquist, B. André, U. Wiklund, P. Svedlindh and U. Jansson In manuscript

Related Papers Thermal Stability of Ti3SiC2 Thin Films J. Emmerlich, D. Music, P. Eklund, O. Wilhelmsson, U. Jansson, J.M. Schneider, H. Högberg and L. Hultman Acta Mater. 55, 1479 (2007)

Structural, Electrical, and Mechanical Properties of nc-TiC/a-SiC Nano-composite Thin Films P. Eklund, J. Emmerlich, H. Högberg, O. Wilhelmsson, P. Isberg, J. Birch, P.O.Å. Persson, U. Jansson and L. Hultman J. Vac. Sci. Technol. B 23, 2486 (2005)

Electronic Structure Investigation of Ti3AlC2, Ti3SiC2, and Ti3GeC2 by soft X-ray Emission Spectroscopy M. Magnuson, J.-P. Palmquist, M. Mattesini, S. Li, R. Ahuja, O. Eriksson, J. Emmerlich, O. Wilhelmsson, P. Eklund, H. Högberg, L. Hultman and U. Jansson Phys. Rev. B 72, 245101 (2005) High-Power Impulse Magnetron Sputtering of Ti-Si-C Thin Films from a Ti3SiC2 Compound Target J. Alami, P. Eklund, J. Emmerlich, O. Wilhelmsson, U. Jansson, H. Hög-berg, L. Hultman and U. Helmersson Thin Solid Films 515, 1731 (2006) Nanocomposite nc-TiC/a-C Thin Films for Electrical Contact Applications E. Lewin, O. Wilhelmsson and U. Jansson J. App. Phys. 100, 054303 (2006) Growth and Characterization of MAX-Phase Thin Films H. Högberg, L. Hultman, J. Emmerlich, T. Joelsson, P. Eklund, J.M. Molina-Aldarequia, J.-P. Palmquist, O. Wilhelmsson and U. Jansson Surf. Coat. Technol. 193, 6 (2005)

My contribution to the papers Publication I: I have performed the experimental work and most of the characterization considering TEM, nanoindentation and XRD. I have done the calculations. I have written the manuscript. Publication II: I have performed the experimental work and the characteri-zation. I have also written the manuscript.

Publication III: I have deposited the films and carried out the X-ray charac-terization. I have taken part in discussion considering the writing of the manuscript. Publication IV: I have deposited most of the films and carried out part of the TEM analysis. I have written the manuscript. Publication V: I have written the manuscript, deposited part of the films, and carried out the XPS-analysis and XRD characterization. Publication VI: I have performed the TEM and XPS analysis. I have taken part in the discussion considering the writing of the manuscript. Publication VII: I have performed the experimental work considering film deposition, carried out characterization by XPS, XRD and Raman, and writ-ten the manuscript. I have taken part in the magnetic and tribological charac-terization.

Contents

Chapter 1 Introduction ..................................................................................11 1.1 Research objectives ............................................................................12

Chapter 2 Transition-metal carbides and related compounds .......................14 2.1 MC – Binary carbides ........................................................................14 2.2 Disordered solid solutions ..................................................................15 2.3 M3AC – Perovskite carbides ..............................................................16 2.4 Mn+1ACn – MAX-phases ....................................................................16

Chapter 3 Thin film synthesis .......................................................................19

Chapter 4 Thin film growth ..........................................................................21

Chapter 5 Characterization ...........................................................................23 5.1 X-ray diffraction.................................................................................23 5.2 X-ray photoelectron spectroscopy (XPS)...........................................24 5.3 Soft X-ray spectroscopy .....................................................................25 5.3 Raman spectroscopy...........................................................................25 5.4 Scanning electron microscopy (SEM)................................................26 5.5 Transmission electron microscopy (TEM).........................................26 5.6 Electrical characterization ..................................................................27 5.7 Mechanical characterization – Nanoindentation ................................28 5.8 Tribological characterization..............................................................29

Chapter 6 Theory ..........................................................................................31

Chapter 7 Results and discussion..................................................................32 7.1 Film Synthesis of MAX-phases in the Ti-Al-C and V-Ge-C systems32 7.2 Synthesis of nanocomposites in the Ti-Fe-C, V-Ge-C and Ti-Al-C systems .....................................................................................................37 7.3 Mechanical properties of the MAX-phases and nanocomposites ......41 7.4 Tribological properties of the films....................................................45 7.5 Electrical properties of the films ........................................................50 7.6 Magnetic wear-resistant Ti-Fe-C films ..............................................52

Future outlooks .............................................................................................54

Acknowledgement ........................................................................................55

Populärvetenskaplig sammanfattning ...........................................................56

References.....................................................................................................59

11

Chapter 1 Introduction

The importance of material for mankind can easily be understood by naming of eras as the Bronze Age and the Iron Age. Roughly speaking, material working at Bronze Age consisted of smelting ores containing copper and, e.g., tin and alloying those to cast bronze for weapons, tools and handicraft [1]. During the Iron Age the heat sources were improved and it became pos-sible smelting ores containing iron [2]. The advantages of iron compared to bronze at that time were larger resources, higher hardness and durability. The carbon content of the iron produced was 0.02-1.7 wt% with properties ranging from rather soft to brittle. For heavy duty applications as weapons and hammers the pure material was not so well-suited. However, by welding iron with low and high carbon contents followed by a repeated forging and folding a laminated multilayer-like material named Damascus steel with toughness and hardness was achieved. Recently, this material has also been shown to contain carbon nanotubes and cementite nanowires [3] (see Fig. 1.1). Without knowing, for more than 400 years ago, the blacksmiths had designed a very advanced nanoscaled composite material.

The basic principles considering alloying, forming nanoscaled structures and rein-force bulk material by whiskers are more or less the same today as for several hun-dred years ago. However, there is a sub-stantial difference since the materials fab-ricated today can be synthesized in a con-trolled manner and their physical and chemical properties can be characterized and also theoretically modeled. Further-more, the available techniques make it possible to optimize bulk and surface separately.

By applying an appropriate film on a bulk material, functional properties such as hardness, toughness, friction, catalysis,

bio-reactivity, oxidation resistance, electrical properties and thermal stability can be designed and tailored for a certain application, but there is no “uni-versal” film for all these features. However, there are examples of self-

Figure 1.1. TEM micrograph of steel of a Damascus blade showing cementite nanowires encapsulated by carbon nano-tubes (Nature 444, 206, 2006).

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adapting thin films, which change their characteristics to the prevailing envi-ronment. An illustrative example is nanocomposite films of WC/DLC/WS2 where carbon and WS2 contribute to low friction in humid atmosphere and vacuum, respectively [4]. Furthermore, nanocomposites, multilayers and gradient layers also offer a functional and structural design of thin films by combining different texture, phase composition, grain size, layer thickness, and matrix volume. Although, a “universal” film is not probable in the future and possibly not either desirable due to difficulties during, especially, the deposition process. A more likely development is films, which are highly optimized for its intended application.

To deposit thin films, there are different methods at hand. One group of techniques is Physical Vapor Deposition (PVD) where, e.g., sputtering and cathodic-arc deposition are included. 1852 W.R Grove was the first person to study sputter deposition when he observed a “cathodic disintegration” of silver [5]. In 1877 A.W. Wright reported on the use of an “electrical deposi-tion apparatus” to form mirrors [6]. Since then the PVD techniques have become more versatile and is today applied in, e.g., tooling, electronic, deco-rative and medical industry. Within this thesis DC magnetron sputtering has been used.

1.1 Research objectives The aim with this thesis has been to characterize phase composition and microstructure of binary and ternary carbide thin films to understand how their physical properties and final performance can be tuned and optimized. To achieve this goal the sputtering process has been employed. This tech-nique gives a high control over individual elemental fluxes and synthesis far from thermodynamical equilibrium. In addition, to understand the final film performance a number of analytical and theoretical tools such as X-ray dif-fraction (XRD), transmission electron microscopy (TEM), pin-on-disc test and density functional theory (DFT) have been used. Figure 1.2 gives an overview of how the sputtering process has been used to synthesize different phases and structures in this thesis.

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Figure 1.2. Schematic overview of the studied phases and structures in this thesis.

14

Chapter 2 Transition-metal carbides and related compounds

In binary and ternary transition-metal carbides the early transition metals (M) constitute a closed-pack like arrangement with the non-metal atoms (C) in the octahedral interstitial sites or in the center of a trigonal prism [7]. The non-metal atom and the nearest metal atoms comprise a coordination poly-hedron. The polyhedrons share sides, corners or common edges. This gives rise to compounds such as monocarbides, perovskites and MAX-phases (see Fig. 2.1).

Figure 2.1. Crystal structures (from left to right) of cubic MC, M3AC (perovskite) and M3AC2 (MAX-phase). Coordination polyhedrons of M6C are shown. These share sides, corners or edges depending on crystal structure. Red = M, blue = A and black = C, e.g., M=Ti and A=Al.

2.1 MC – Binary carbides Thermodynamically stable binary metal carbides are found in group 4 to 6 in the periodic table. These are, e.g., known for their hardness (between Al2O3 and diamond), high melting point (2000-4000oC) and electrical conductivity [7]. Moving from group 4 to 6 the crystal structure changes. Group 4 (Ti, Zr and Hf) only form a NaCl-type crystal structure with common sides of the M6C polyhedrons. Group 5 (V, Nb and Ta) exhibits NaCl-type crystal struc-ture, but also a M2C hexagonal structure. Group 6 (Cr, Mo and W) forms several different cubic and hexagonal crystal structures.

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The bonds in monocarbides have a mixture of ionic, covalent and metallic character. These arise from a charge transfer from M to C, hybridization between C p and M d states, and pure M d state overlapping, respectively [8, 9]. The cohesive energy of monocarbides with a NaCl-crystal structure de-creases from group 4 to 6 and thereby the stability of the structure. This can be interpreted in terms of filling bonding, non-bonding and anti-bonding states arising from the, above mentioned, hybridization of C p states and M d states. In terms of bonds, the M-C and M-M strength decreases from group 4 to 6, while the interaction between C-C increases [8, 10]. Hence, when the C-C interaction increases the NaCl-crystal structure becomes metastable. As a result experimentally synthesized carbides always contain C vacancies. For example, the homogeneity range of TiCx varies between 0.47

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leads to a decrease grain size, which can influence tribological and mechani-cal properties.

Figure 2.2. Ternary phase diagrams of the Ti-Al-N (left) and Ti-Al-C (right) sys-tems [21]. Small regions with a solid solution of Al can be observed in both dia-grams. A solid solution gives a possibility to tune, e.g., mechanical properties and manipulate chemical bonds in the host lattice.

2.3 M3AC – Perovskite carbides In the perovskite structure the M6C polyhedron share corners with each other and form a simple cubic array (see Fig. 2.1). The A atoms, e.g. Al, Ge and Sn also form a cubic pattern surrounding the polyhedrons, i.e. one can say that the A atoms form an ordered solid solution. There is an interest in these non-oxide perovskites, since Ni3MgC recently has been shown to exhibit a superconductive transition at 8K [23]. Furthermore, from theoretical calcula-tions it has been shown that Al reduces the directionality of the bonding when it is introduced in TiC [24]. This can, e.g., change the mechanical properties.

2.4 Mn+1ACn – MAX-phases The first observation of the nanolaminated MAX-phases was done in the 60’s by Jeitschko et al. [25]. However, not until 1996 the real scientific breakthrough occurred. At this time Prof. M.W. Barsoum reported a unique mixture of ceramic and metallic properties for Ti3SiC2 and coined the aberra-tion “MAX-phase” [26]. As was mentioned above M stands for an early transition metal, while A is an element from group 13-16 in the periodic ta-ble and X is nitrogen or carbon. In the year 2000 there were about 50 known

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MAX-phases [27]. However, today previously unknown phases have been added to the list via thin-film synthesis and heat treatment of bulk samples [28-30].

In the MAX-phases the M6C polyhedrons share six of its twelve edges with adjacent octahedral forming M2C, M3C2 and M4C3 layers (i.e. Mn+1Cn). Above and beneath the Mn+1Cn-sheets tetrahedral voids are formed. At these sites A-atoms are positioned forming single layers. The whole structure can be seen as a nanolaminate. The general formula is written Mn+1AXn (n = 1, 2, 3) or shortly “211”, 312” and “413”. Figure 2.3 shows the hexagonal unit cells of the 211, 312 and 413 phases where the Mn+1Cn-sheets are interleaved by A-slabs. Except for the structures shown in Fig. 2.3 there are also exam-ples of so-called intergrown phases. For example, half a unit cell Ti3SiC2 and Ti2SiC sums up to Ti7Si2C5, which has a very long c-axis of ~60 Å [29].

Figure 2.3. Crystal structure of M2AX, M3AX2 and M4AX3 MAX-phases (211, 312 and 413). In the figure a different number of Mn+1Xn layers are interleaved by planar slabs of A-element for the different structures. Furthermore, for 312 and 413 there are two types of M-atoms; M1 that binds to X and A, and M2 that solely binds to X.

MAX phases share many of the properties of their binary relatives as electri-cal and thermal conductivity, and high melting points. However, mechani-cally they cannot be more different, since the MAX-phases are often machi-neable and soft. This is explained by their nanolaminated structure with strong M-X bonds and much weaker M-A bonds [31, 32]. Furthermore, in the “312” and “413” MAX-structures there are two types of M-atoms. One that bonds to both X and A (M1) and one that bonds solely to X (M2) (see Fig. 2.3). M1 and M2 have both a directional covalent character in their bonds but the former is, in similarity with the binary relatives, much stronger and has also an ionic contribution. Metallic properties are observed in the basal planes, while it is much weaker in the c-direction. The A-elements form a covalently bonded single layer. This bond configuration gives rise to deformation via basal-plane slip and kink formation and electrical transport in the MX-planes [32-35]. Not surprisingly the hardness of the MAX phases is different if measured perpendicular or parallel to the basal planes. Nickl et

18

al. reported that the hardness for single-crystal Ti3SiC2 is about 3 time higher normal to the basal planes (12-15 GPa) than parallel to them (3-4 GPa) [36]. The basal-plane slip, together with grain buckling, crack deflection and crack branching give rise to a damage tolerant bulk material [37].

The anisotropic structure and weak M-A bonds indicate that MAX phases may have low-friction properties. Myhra et al. have measured the friction coefficient on the nanoscale of the basal planes of Ti3SiC2 to

19

Chapter 3 Thin film synthesis

Physical Vapor Deposition (PVD) includes several different techniques such as sputtering, cathodic arc and electron-beam evaporation. Common for all PVD techniques is that a solid material is exposed to heat or momentum to transfer it into a vapor-like state of aggregation and to subsequently con-dense on a substrate forming a thin film. In this thesis direct current (dc) magnetron sputtering has been used.

A sketch of the sputtering process is shown in Fig. 3.1. Principally, Ar- gas is ionized in a vacuum chamber and accelerated into the target (cathode). The momentum of the ions is transferred to the target and atoms are sput-tered. The atoms are spread out in the chamber and a film starts to grow. This impact also generates secondary electrons. To increase the deposition rate and film quality, magnets can be introduced behind the target. This en-traps secondary electrons and thereby increases the plasma density in front of the target. This configuration is called dc magnetron sputtering. Further-more, the magnet configuration can be changed to modify the ion bombard-ment at the film surface (see Fig. 3.1).

Figure 3.1. Left; sketch of the sputtering process where Ar-gas is ionized and accel-erated into the cathode (target) sputtering target atoms. The atoms are transported to the substrate and condense into a film. Right; sketch of the magnetron where the magnets trap secondary electrons, which increase the plasma density in front of the target. This improves the sputtering rate and film quality. Furthermore, depending on the magnet configuration the ion bombardment at the substrate can be modified.

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The ions that strike the growing film surface will have an energy corre-sponding to the floating potential (Vf). A typical value of Vf is 2 Volts and higher [47]. This value can be compared with the activation energy (Eact) for diffusion of atoms on a surface. For example, the Eact for diffusion of Ni atoms on a Ni(100) surface is reported to 159 kJmol-1, which is equal to ~1.7 eV/Ni atom [48]. Correspondingly, this energy is of the same range as Vf. This “heating” is very different to a conventional heating source since the kinetic energy of the bombarding ions is transferred to very small areas and conveyed in the near vicinity. This heating is also accompanied by very fast cooling rates of ~1014 K/s [49]. The ion bombardment can be further in-creased by applying a bias to the substrate. This will strongly influence the film growth and the final properties. In summary, the sputtering process is working far from thermodynamically equilibrium. This makes it possible to deposited, e.g., metastable solid solutions.

In this thesis elemental targets have been used. This gives a high flexibil-ity and a wide range of phases can easily be deposited. However, in indus-trial applications it may be more convenient to use a compound target with a desired composition. This can be accompanied with some difficulties. At steady-state, the composition of the sputtered atoms that reach the growing film surface should be the same as in the bulk of the target. However, due to different mean-free paths and mass of the atoms there may not be a 1:1 situa-tion. This can be compensated by using, e.g., an extra target or manufacture a compound target that gives the accurate composition at the film surface. For further details about sputtering and details about PVD see, e.g., [6, 50].

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Chapter 4 Thin film growth

The growth of thin films usually proceeds through nucleation and different growth stages. Such stages involve adsorption, surface diffusion, bonding and nucleation [51]. All these processes are dependent upon the mobility of the deposited atoms on the surface, surface steps, lattice defects, and con-tamination. Each of them have their characteristic activation energy [50]. In the ideal case (no surface defects and interdiffusion) the initial stages of film growth proceeds via three modes. These are; Frank-van der Merve (layer-by-layer), Volmer-Weber (island growth) and Stranski-Krastanov (initially layer-by-layer followed by islands) [52]. In epitaxial growth the Frank-van der Merve is preserved throughout the film, while the Volmer-Weber is the initial stage of the polycrystalline growth.

In epitaxial growth the deposited film is crystallographically aligned to the substrate, both in-plane and out-of-plane. This kind of growth appears when single-crystalline substrates are used acting as templates. The direction is not arbitrary, but such that a minimal misfit between the substrate and film is achieved. For example, when a TiCx(lll) film is deposited on �-Al2O3(000l) there will be a 30o relative rotation of the two compounds. This is shown in Fig. 4.1 where the [101] direction of a TiCx(lll) coincide with the [210] direction of �-Al2O3(000l). With this orientation the misfit will be ~6%. Growing films can compensate for misfits by introducing misfit-dislocations in the interface or be tetragonally distorted.

Figure 4.1. Left; Al2O3(000l) surface. Red atoms = O and violet = Al. Right; TiC(lll) surface. Blue atoms = Ti and black = C. In MAX-phase deposition the [101] direc-tion of the TiC seed-layer coincide with the [210] direction of the Al2O3 substrate.

Polycrystalline thin film growth on amorphous or randomly oriented sub-strates proceeds via nucleation, island growth, coalescence of islands, devel-

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opment of a continuous film and thickening [53]. During the coalescence less energetic favorable grains will be consumed and a columnar growth may be observed. However, this is strongly dependent on, e.g, temperature, depo-sition time and process pressure and can be described in structure zone mod-els (SZM). A classical example is the model presented by Thornton for ele-mental sputtering of metals. In SZM the microstructure changes from ta-pered crystallites separated by voids to equiaxed large grains when the depo-sition temperature is increased or the Ar-pressure is reduced (see Fig. 4.2) [54].

Finally, from the understanding of thin-film formation and thin-film syn-thesis follows the possibility for microstructural and nanostructural engineer-ing to design material properties for specific technological applications.

Figure 4.2. Structural zone model proposed by Thornton for sputtered metal coat-ings. Different microstructures (zone 1-3 and zone T) can be identified depending on Ar-pressure and substrate temperature [54].

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Chapter 5 Characterization

To correlate microstructural properties to a macroscopic behavior a number of analytical tools are necessary, which may verify and complement each other. In this thesis analytical techniques have been used to characterize the microstructure, phase and chemical composition, electrical, mechanical and tribological properties of the deposited films. The aim of this chapter is to describe some basic characteristics of the instruments and what kind of in-formation that can be attained.

5.1 X-ray diffraction X-ray diffraction is a powerful technique for phase analysis, studying epi-taxial relations as well as preferential orientation and measure residual stress [55]. The grain size is easily extracted from X-ray measurements since the peak broadening is dependent on the crystallinity. For high crystalline films the diffraction peaks are sharp, while they become broader with increasing structural disorder. The grain size can be estimated by applying Scherrer’s formula [56].

To obtain the information, mentioned above, different instrumental setups and scan types are needed. For example, for phase analysis of polycrystalline films so-called Grazing Incidence (GI) scans are well suited. This method will increase the diffraction volume from the film and decrease the substrate signal. Furthermore, by varying the incidence angle the probe depth can be varied and information of, e.g., elemental surface segregation can be ex-tracted. The GI-setup has been used in, e.g., [V-VII]. In [I] and [II] epitaxial MAX-phase films where deposited. These films have only one out-of-plane orientation and �-2� scans are necessary for phase analysis. For such films the residual strain can be measured by reciprocal space mapping (RSM) [57]. Furthermore, by mapping the position of asymmetric and symmetric peaks, the in-plane and out-of-plane cell parameters can also be calculated [58]. This method has been used in [I] and [II].

A well established technique to quantify the strain level in polycrystal-line-thin films is the sin2�-method where d is plotted versus the tilt angle (�) [59]. For a biaxial case the residual strain (��) follows,

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� � �� EE2sin1 2 � (5.1)

where E is Young’s modulus, � Poisson’s ratio and � the in-plane stress. The lattice spacing d can be attained according to,

0

0

ddd �

�� (5.2)

where d� and d0 is the measured and the strain-free lattice-plane distance, respectively. With Eq. 5.1 and 5.2, and knowledge about the elastic con-stants of the analyzed material can be calculated and a stress-free cell pa-rameter obtained. This method has been used in [VI].

In conclusion, the X-ray diffraction technique is very well suited for thin film analysis since several film properties can be measured in a fairly simple way.

5.2 X-ray photoelectron spectroscopy (XPS) X-ray Photoelectron Spectroscopy (XPS) is a surface sensitive (probe depth ~50 Å) analysis method for quantitative, qualitative and chemical analysis. By irradiation of the sample with X-rays of well-defined energy (h) core electrons are excited by the photoelectric effect, escape into vacuum and can be analyzed with respect to their specific kinetic energy (EK) [60]. In the XPS spectrum the intensity is plotted versus the binding energy, which is calculated from the energy conservation law. Beneficial with the technique is that the binding energy of an element may change if the chemical environ-ment is modified. For example, the binding energy of C1s for C-C and Ti-C is ~284 eV and ~281 eV, respectively. Interpretation difficulties may, how-ever, arise in the chemical analysis of C1s. For example, several authors have noticed that an extra peak appears in the spectrum of nc-TiC/a-C thin films [61, 62]. This feature is related to the nanocomposite structure and the intensity is enhanced by Ar-depth profile sputtering.

The XPS instrument used in this thesis is a PHI Quantum 2000. This in-strument is equipped with a secondary electron detector (SXI) and the analy-sis spot size can be varied between 10-150 �m. This makes it possible to analyze small structures such as scratches. In [V-VII] the instrument has been used to study the chemical composition and chemical state of the film in wear tracks or the adhered material after pin-on-disc tests.

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5.3 Soft X-ray spectroscopy In soft X-ray spectroscopy, the sample is irradiated by tunable synchrotron radiation and the emission of photons (XES) or absorption is measured (XAS) [63]. With these methods unoccupied and occupied electronic states can be probed. Figure 5.1 shows schematically XAS and XES processes as well as the XPS process for comparison. As was discussed in sec. 5.2 in XPS measurements, absorption of X-rays can result in an excitation of core elec-trons into vacuum. However, core electrons can also be excited into an un-occupied state (so-called resonant excitation). If the absorption of the incom-ing X-rays is measured, an XAS spectrum can be collected. It can be noted that the absorption of the incoming X-rays are indirectly measured by prob-ing the Auger electrons. After the excitation, a spontaneous radiative decay can occur. By detection of the emitted photons an XES-measurement is car-ried out. A non-resonant XES spectrum can be interpreted in terms of partial valance band DOS (PDOS) and compared with theoretical results. Soft X-ray measurements have been carried out on Ti2AlC MAX-phase and TiC0.67 thin films in [III] to study the bond characteristics of the two phases.

Figure 5.1. Schematic sketch of XPS, XAS and XES where the kinetic energy of photoelectrons (EK), adsorption of incoming X-rays (h2) and emitted photons are measured (h3), respectively.

5.3 Raman spectroscopy Raman spectroscopy relies on inelastic scattering of monochromatic light by molecules. Usually, light from a laser in the visible, near infrared or near ultraviolet range is used. The laser light is focused and the spot size is ~1 �m. Consequently, small structures can be studied. The interaction between the incoming light (photons) and the sample induces transitions between vibrational-energy states of the molecule. These transitions correspond to excitation of electrons into a virtual state followed by the relaxation into a vibrational excited state. The excitation of electrons can be seen as a polari-

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zation of the chemical bond and therefore are all materials not Raman active. An example is metals. For Raman active materials, however, the transitions are very specific and chemical information can be attained. In [VI-VII] the method has been used to study the sp2 and sp3 hybridization of carbon and the oxide type in wear tracks after pin-on-disc measurements.

5.4 Scanning electron microscopy (SEM) The first commercial SEM was supplied by Cambridge Instruments 1964 [64]. In a simple way the technique gives topographic and chemical informa-tion. Both elastic and inelastic electrons can be detected to create a contrast in the image. In the first case so-called backscattered electrons are collected forming an atomic contrast in the image. The cross-section of the process is dependent of atomic number (Z) – the higher Z the brighter appears the ana-lyzed area. Secondly, by detection of inelastic-scattered electrons, i.e. secon-dary electrons (SE), SEM-images with a three-dimensional appearance can be created. These electrons have a low energy ( 5° can be selected. The scattering angle depend on the atomic number of the sample and are used to form so called Z-contrast images [66].

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Microscopes may also be equipped with energy-filter systems, which allow recording of inelastically scattered electrons. With these electrons, electron energy loss spectra (EELS) or energy-filtered images (EF-TEM) can be ob-tained, enabling chemical information of the sample. In this thesis a post-column system (Gatan image filter, GIF2002) has been used, which consists of a magnetic prisms that bends the electrons 90o off the initial optic axis [67]. Electrons that have lost different amount of energy will be separated into an energy-dispersive spectrum. For elemental mapping an energy select-ing slit is used to choose a certain energy range. Figure 5.2 shows an exam-ple of EF-TEM images for Ti, Fe and C where the elemental distribution is seen. Chemical information can also, in similarity with SEM/EDS, be at-tained by scanning-TEM/EDS analysis. Compared to EF-TEM the resolution is worse, but the demand on the sample is less. TEM analysis has been used in [I-II], [IV] and [VI].

Figure 5.2. EF-TEM image of heat-treated (Ti,Fe)C film where regions of Fe, Ti and C has formed. The brighter contrast, the higher is the concentration of the element.

5.6 Electrical characterization The resistivity of thin films can be measured by four-point probe where four equally spaced tungsten tips are made to contact the film surface under measurement (see Fig. 5.3). A electric current is supplied to the outer probes and the voltage drop is measured between the middle probes to determine the resistivity. The measured quantity is the sheet resistivity (Rs) in �/square. With knowledge about film thickness the bulk resistivity can be calculated. This method has been used in [I-II].

Figure 5.3. Schematic sketch of the four-point probe setup with a current flowing between the outer probes simultaneously measuring the voltage drop between the two inner tips.

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5.7 Mechanical characterization – Nanoindentation To quantify the hardness and Young’s modulus of thin films, the intention is to obtain a fully developed plastic zone in the film. This is achieved by forc-ing an indenter into the material. However, the depth of the plastically de-formed zone under the indenter is ~10 times the indenter depth. Therefore, traditional micro-hardness measurement is not applicable for thin films since the substrate would influence the measurement. By nanoindentation the maximum indentation depth can be reduced. During the indentation the ap-plied load (P) and the penetration depth (h) is recorded and presented as a loading-unloading curve (see Fig. 5.4). The slope (dP/dh) of the initial part of the unloading curve is indicative for the stiffness (S) of the whole contact, i.e. indenter tip and film. By applying the Oliver-Pharr method on this part the reduced Young’s modulus (E) can be calculated [68]. The hardness (H) is defined as the maximum load (Pmax) over the area of residual imprint.

Figure 5.4. Left; Load-unloading curve for nanoindentation. Right; imprint after indentation [69].

The nanoindentation method is straight forward to use. However, the me-chanical response of the material may influence the final result. An example is formation of kinks and pile-up under indentation of MAX-phases (see Fig. 5.5). This may change the real residual imprint area. Furthermore, pop-in events may occur. This corresponds to a sudden penetration of the indenter and is accompanied by a reduction of hardness and Young’s modulus as the actual indentation depth increases (see Fig. 5.5).

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Figure 5.5. Left; persistent kink formation after indentation of Ti3SiC2 MAX-phase [70]. Right; loading-unloading curve where pop-in events are seen. The pop-in is accompanied by a sudden increase in penetration depth and is typically observed during indentation of MAX-phases.

5.8 Tribological characterization Friction takes place when two surfaces with a relative motion against each other come into contact. The friction law states that the friction coefficient (�) is the ratio between the measured friction force (FL) and the normal force (FN). The friction coefficient can be measured in several ways. In [V-VII] the pin-on-disc method has been used where FL is continuously measured for a certain FN [71]. Figure 5.6 shows a schematic picture of the setup. With this method, � can be calculated, and the wear rate and wear mechanism can be studied.

Dry sliding friction can often be separated into an adhesive part and a ploughing part. The adhesive part is associated with the formation of prefer-able easily sheared tribochemical layers, while ploughing partly arise due to film asperities [72]. The tribological response is deduced by the whole sys-tem, e.g., temperature, material combination, substrate, applied load and atmosphere. Consequently, any experiment to characterize the tribological behavior of a film only depicts part of the reality. During tribological tests a film can suffer from, e.g., adhesive wear, which is a result of intrinsic plastic shear or cracking of the film. A low friction coefficient, high hardness and low solubility of the counter surface are desirable to reduce this wear. How-ever, it should be noted that a high hardness is not necessarily the prime requirement for wear resistance, but also a low Young’s modulus (E). Such combination results in wear resistant and fracture toughness [73]. Except for the normal film wear, spallation by interfacial failure may also occur. This is explained by a poor film adhesion to the substrate. The adhesion can be quantified by scratch tests where a diamond stylus is slid over the film sur-face. During the test the applied normal load is increased continuously until

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the film detaches. The scratch test also gives an opportunity to study crack formation (see publication VI).

Figure 5.6. Schematic sketch of the setup for pin-on-disc measurements. The fric-tion force (FL) is continuously measured for a certain applied load FN and rotation speed. A wear track will form in the film and wear debris will adhere to the counter surface.

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Chapter 6 Theory

To theoretically study real solids a large number of electrons have to be han-dled (1023), but still the Schrödinger equation has to be solved. However, this is not feasible for all electrons and therefore approximations and simplifica-tions have to be made. A useful theory to study real solids is the ab initio density functional theory (DFT) based on electron charge densities [74]. The theory was initially introduced by Hohenberg and Kohn, and later improved by Kohn and Sham [75]. This method implement approximations as Born-Oppenheimer and especially the local density approximations (LDA), which makes it possible to solve the Schrödinger equation with local effective po-tentials [76].

In [I], [III] and [V] DFT-theory has been used. These calculations has made it possible to discuss the absence of intergrown MAX-phase structures in the Ti-Al-C system, compare and interpret experimental soft X-ray emis-sion spectra and study the effect of a solid solution of weak-carbide forming element in TiC. Figure 6.1 shows an example from [V] where a DOS curve for TiC is shown together with its charge density map. The DOS curve gives an idea about the electron distribution for a certain band. In the DOS plot, two regions can be identified; one below the Fermi level and one above. These two regions correspond to bonding and antibonding states, respec-tively. For example, for TiC C2s has bonding states at -10 eV, which are strongly correlating with Ti2p located at the same level. This designates a strong hybridization. The chemical interaction can, however, be more closely studied by crystal orbital overlapping population (COOP) curves (see [III]) or charge density plot (see Fig. 6.1). Such plots display chemical bonds and the electron distribution in real space.

Figure 6.1. Calculated density of states (DOS) for TiC. The inset shows the charge density plot for the 110 plane of TiC where the electron density is evenly distributed.

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Chapter 7 Results and discussion

In the periodic table shown in Fig. 7.1 the elements used in this thesis are highlighted. The constituting elements have been an transition metal (M = Ti, V and Fe), an element from group A (A = Al, Si and Ge), and carbon (C). Depending on the chosen M-A-C system different binary and ternary carbide thin films have been synthesized and characterized.

Figure 7.1. Periodic table where the elements used in this thesis are highlighted.

7.1 Film Synthesis of MAX-phases in the Ti-Al-C and V-Ge-C systems

Epitaxial MAX-phase films of Ti3AlC2, Ti2AlC and V2GeC were deposited by DC magnetron sputtering on single crystal Al2O3(000l) substrates (see [I] and [II]). Figure 7.2 and Fig. 7.3 show the �-2� diffractograms of the Ti3AlC2 and V2GeC temperature deposition series. As can be seen the MAX-phase formation is strongly temperature dependent. Ti3AlC2 is formed above 800 oC, but not at lower temperatures. In similarity with Ti3AlC2, also Ti2AlC vanish below 800 oC. In contrast to the Ti-Al-C system, the V2GeC MAX-phase structure is formed in a wider temperature range from 450 to 850 oC. The critical nucleation temperature of the MAX-phases can partly be explained by diffusion requirements. The c-axes of MAX-phases are very long and the unit cell is large. To form such a structure a rather high surface diffusion rate (i.e. a high deposition temperature) is needed. For example, Högberg et al. have observed a slightly higher temperature for the growth of Ti4GeC3 (c~22.7 Å) compared to Ti3GeC2 (c~17.9 Å) [77]. For the growth

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of Ti3AlC2 (c = 18.58 Å) and Ti2AlC (c = 13.60 Å) in [I] and [II], however, no clear difference in nucleation temperature was observed, in spite of the longer c-axes of the “312”-phase.

Figure 7.2. �-2� diffractograms of Ti3AlC2 films deposited between 300-900 oC [I].

Figure 7.3. �-2� diffractograms of V2GeC films deposited between 300-850 oC [II].

The c-axis of the V2GeC phase was measured to 12.28 Å, which is shorter than for both the studied Ti-Al-C phases. At a short glance this could possi-bly explain the much lower nucleation temperature of V2GeC. However, the elemental composition has also to be considered. For example, Walter et al. have deposited Cr2AlC (c = 12.80 Å) at 450 oC [45]. In comparison with our reported results on Ti2AlC the lower deposition temperature is consistent with the shorter c-axis of Cr2AlC. However, at the same time the number of d-electrons of the M-element has been increased when replacing Ti for Cr. This suggests that the deposition temperature for MAX-phases may be re-duced when using M-elements from the third period in the periodic table

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with more d-electrons than Ti. However, further studies including theoretical modeling are required to determine the influence of d-electrons on the MAX-phase stability.

The stability of the MAX-phases seems to be reduced at too high tem-peratures. This can be seen in Fig. 7.3 where the peak intensity of the mono-carbide increases when the deposition temperature is changed from 700 oC to 850 oC. This can be explained by the evaporation of Ge from the growing film surface, which makes it more favorable to form a larger fraction of the monocarbide. The effect can be attributed to the melting point of the A-element and is not influenced by the M-element. For example, it has been shown that there is no significant difference in the evaporation temperature of Ge in the Ti-Ge-C and V-Ge-C systems [28] [II]. This is also supported by the decomposition studies of epitaxial thin Ti3SiC2 films where it was demonstrated that Si starts to evaporate at ~1160 oC in high-vacuum [78]. Furthermore, theoretical calculations show that the extra electron of Ge compared to Al contributes to antibonding states in the A-layer of Ti2GeC [79]. This decreases the s-p hybridization between Ge atoms, which may favor the tendency for Ge-segregation as observed in the Ti-Ge-C system.

The MAX-phases are highly ordered structures but we have in [I] demon-strated that the there is an energy cost for the insertion of A-element planes between the carbide slabs. This energy cost is larger in the Ti-Al-C system than in e.g. the Ti-Si-C system which can be attributed to the weaker Ti-Al bond (see Fig. 7.4). Furthermore, the average energy required to insert a Si-layer in Ti-Si-C is independent of the number of layers. This suggests that a random stacking or more complex stacking sequences could be formed. In the Ti-Si-C and Ti-Ge-C systems such intergrown structures of Ti7Si2C5 and Ti5Ge2C3 was actually observed together with random stacking sequences [28, 29]. In the Ti-Al-C system, however, the slope of the curve in Fig. 7.4 is initially steeper. These results suggest that the energy requirement for ran-dom stacking and intergrowth is less favorable in Ti-Al-C. This was also the case and only an occasionally stacking sequence of Ti5Al2C3 was observed by TEM-analysis (see Fig. 7.5).

Figure 7.4. Cohesive energy from ab initio calculations for the Ti-A-C (A = Si or Al) system as a function of inserted A-layers. The curves are only a guide for the eye [I].

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In addition to the energy cost for the A-plane insertion, the diffusivity of the A-element has to be considered. For example, Högberg et al. have observed that Ti3GeC2 and Ti5Ge2C3 are formed in Ti2GeC films when the deposition temperature was increased from 850 oC to 1000 oC [80]. From the discussion in the previous paragraph it was concluded that the surface segregation and evaporation of the A-element is independent on the M-element. Therefore, it is most likely that intergrown structures also can form in the V-Ge-C system at temperatures higher than 850 oC where evaporation and segregation is further increased. In the present study the highest achievable deposition tem-perature was only 850 oC.

Figure 7.5. High-resolution cross-sectional TEM of an epitaxial Ti3AlC2 film. Within the squares a rare Ti5Al2C3 stacking sequence can be seen (black circle = Ti, white circle = Al atom) [I].

In the diffractograms in Fig. 7.2 peaks from the monocarbides TiCx and VCx/V8C7 are seen. They originate from two sources: a deposited seed layer and/or spontaneously formed carbide inclusions in the film. In previous stud-ies it has been shown that a TiCx seed-layer improves MAX-phase crystallin-ity [29]. We also observed an improved quality for the Ti-Al-C MAX-phase films using such a TiCx seed-layer. The transition from TiCx seed-layer growth to MAX-phase growth is not instantaneous. Instead, we observed a delayed nucleation (see Fig 7.6). This can be explained by the fact that a critical surface concentration of the A-element is required for MAX-phase nucleation. These results are similar to those observed in the Ti-Si-C system [44]. In contrast to the Ti-Al-C system, the quality of V2GeC films was not enhanced by an initial layer of VCx. Instead a direct nucleation on the �-Al2O3 substrate was observed. Therefore, the seed-layer was excluded from the synthesis process (see Fig.7.7).

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Figure 7.6. Cross-sectional TEM micrograph of Ti2AlC MAX-phase film deposited on a �-Al2O3 substrate. In the figure a delayed nucleation and formation of inclusion can be seen [I].

Figure 7.7. Cross-sectional TEM micrograph of an epitaxial V2Ge(000l) film with VCx inclusions deposited on a �-Al2O3 substrate [II].

The formation of carbide inclusions can also be explained by small fluctua-tions in the sputter flux of the A-element. When the surface concentration of the A-element temporarily decreases, a pure binary carbide will start to grow coherently on the MAX-phase. As soon as the critical surface concentration has been reached again the MAX-phase growth will start again and the bi-nary carbide will appear as a coherent inclusion in the film. Figure 7.6 shows an inclusion formed in the Ti2AlC film already at an early stage in the depo-sition process. In Fig. 7.6 a step-like appearance of the surrounding Ti2AlC film can also be seen. This feature is explained by the lateral growth mecha-nism of the MAX-phases suggested by Emmerlich et al. [44]. This step-like appearance is also observed for the inclusion formed during the deposition of the V2GeC film (see Fig. 7.7). In the V-Ge-C system the inclusions consist of both V8C7 and VCx (V8C7 is a superstructure of VCx where the carbon vacancies have been ordered). In Fig. 7.8 a HAADF image of the V2GeC film deposited at 700 oC is shown. Dark and bright areas can be observed,

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which correspond to VCx/V8C7 and V2GeC, respectively. As can be seen in Fig. 7.8 the width and horizontal elongation of the carbide inclusions varies. From Fig. 7.7 it is also evident that some inclusions are formed immediately in the film growth. The MAX-phase structure consists of [lll] oriented MX-slabs and consequently a coherent relation can be expected for the MAX-phase and the formed inclusions. From X-ray pole figures a coherent relation between V2GeC and VCx/V8C7 was confirmed. This crystallographic rela-tionship has also been observed for other MAX-phase systems [81, 82].

Figure 7.8. Left; HAADF micrograph of V2GeC film with MAX-phase structure deposited at 700 oC. The bright areas correspond to MAX, while the dark regions are inclusions of VCx/V8C7. Right; EDX point analysis of region A and B [II].

7.2 Synthesis of nanocomposites in the Ti-Fe-C, V-Ge-C and Ti-Al-C systems

The nanolaminated MAX-phases are only obtained at fairly high deposition temperatures. At lower temperatures the microstructure and phase composi-tion changes and a cubic carbide phase are usually formed. For carbon-rich compositions a nanocomposite is formed with nanocrystalline carbide grains in an amorphous carbon matrix. The composition of the carbides and the relative amount of the matrix can be controlled by e.g. deposition tempera-ture, target currents and also as discussed below by solid solution of a third element. Different types of carbides and nanocomposites were observed in the Ti-Al-C, Ti-Fe-C and V-Ge-C systems. At low carbon contents a pure carbide was usually observed. With the addition Al and Fe, the grain size of the carbide was reduced and free carbon was usually formed. For higher carbon concentrations nanocomposites are almost always formed. In the V-Ge-C system the addition of Ge induced an amorphization of the carbide [II]. In metallic alloys Ge is known for a glass-forming ability and this may explain the observations.

As was discussed in chapter 3 and 4 the sputtering process is an inher-ently metastable method that works far from equilibrium. With the tech-

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nique, substitional solid solutions can be formed that widely exceeds the maximum limit expected from equilibrium phase diagrams of bulk material. An example is shown in Fig. 7.9 where a Ti-Fe-C deposition series is indi-cated in the ternary equilibrium phase diagram for the bulk material. It can be seen that the maximum solubility of Fe into TiC, i.e. the formation of (Ti,Fe)C, is limited to ~2 at% Fe. However, for the deposited films a maxi-mum solubility of ~30 at% Fe was attained.

Figure 7.9. Left; equilibrium phase diagram of bulk Ti-Fe-C at 1000 oC [83]. Indi-cated in the figure is the Ti-Fe-C deposition series (A). Right; GI diffractograms of the films in series A [VII].

The formation of a solid solution of Fe into TiC (i.e. (Ti1-xFex)C) was evident from the X-ray diffractograms shown in Fig. 7.9. As can be seen the peaks are shifted towards higher 2� due to a smaller unit cell. This is expected when the smaller Fe atom replaces Ti in a substitional solid solution. In Fig. 7.9 it can also be seen that the texture and grain size is influenced by the Fe concentration. Similar results were obtained in the Ti-Al-C system in [I] when the deposition temperature was lowered beneath the nucleation point for MAX-phase growth. A maximum solubility of ~25 at% Al was observed for the Ti-Al-C system with a corresponding cell parameter of 4.26 Å [84]. The solid solution of Al into the TiC-lattice induced, in similarity with the Ti-Fe-C system, a decrease of the grain size, but also a formation of C-C carbon and a formation of a nanocrystalline structure. Figure 7.10 shows a TEM image of a (Ti,Al)C film from [VI] where the nanocrystalline structure can be observed. It should be noted that the high concentrations of substitu-tionally dissolved metals into the carbide creates a metastable state. Upon annealing these materials will transform into a thermodynamically more stable state. This can be either formation of free carbon or by formation of metal precipitates. For example, in the Ti-Fe-C system, annealing will lead to the formation of �-Fe precipitates [VII].

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Figure 7.10. Cross-sectional TEM micrograph of nanocrystalline (Ti,Al)C film [VI].

We have discovered that a very important factor in the nanocomposite syn-thesis is the ability of the transition metal to form carbides. It is well known that the stability of carbides decreases with the number of d-electrons in the metal and it is actually energetically unfavorable for transition metals to form carbides after group V in the periodic table [10]. In [V] the concept of weak-carbide forming (WCF) metals as a mean to control nanocomposite formation was introduced. The idea with this concept is that a solid solution of a WCF-metal into a stable monocarbide such as TiC favors the formation of free carbon or graphite. The weak-carbide forming ability of Al was theo-retically studied as a model system. The low tendency of Al to form bonds in the TiC structure can be seen in the DOS-curves and charge density plots in Fig. 7.11.

Figure 7.11. The density of states for (Ti0.5Al0.5)C (upper panel) and TiC (lower panel). In the insets, the calculated charge densities are shown [V].

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From the DOS curves it is concluded that the Ti-d and C-p hybridization in (Ti0.5Al0.5)C is only preserved for C atoms with Ti as nearest neighbor. C atoms that have Al as nearest neighbor only exhibit a broad band. The results of the DOS curves are also shown in real space charge density maps (see inset in Fig. 7.11). From these highly directional bonds are evident for Ti and C in TiC. For (Ti0.5Al0.5)C there is no accumulation of charge at the Al sites, instead the Al charge is smeared out over the whole structure. Conse-quently, the Al-C bond in this structure is weak. This makes it more favor-able to either remove Al from the system or to remove carbon under the formation of graphite. Removal of Al requires solid state diffusion according to a substitutional mechanism, which always requires higher activation en-ergy than interstitial carbon diffusion. Therefore, the most likely process to reduce the total energy is by diffusion of carbon to the surface and formation of free carbon or graphite. The formation energy for this process was deter-mined by calculating the following energy difference,

� �� CAlTiEgraphiteyECAlTiEE xxyxx �� 111 The result is shown in Fig. 7.12 where it can be seen that for an Al-content higher than 12.5% (x>0.25) it is energetically favorable to spontaneously form graphite.

Figure 7.12. Energy change for graphite formation from a (Ti1-xAlx)C structure. The energy difference is plotted against the Al content (x) for two different amounts of released carbon, y. For negative values there is a spontaneous formation of graphite [V].

These results were also verified experimentally. Figure 7.13 shows experi-mental C1s spectra of TiC0.5, and as-deposited and heat treated (Ti0.3Al0.7)C0.5 films. As can be seen Al induces a broadening of the C1s peak towards higher binding energies corresponding to C-C bonds. Furthermore, by a heat treatment step the amount of surface C-C can be increased. In similarity with the theoretical calculations a x>0.25 is needed for a spontaneous formation

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of C-C. In sec. 7.4 the WCF concept will be discussed in terms of low-friction films.

Figure 7.13.a) Typical C1s spectra of 1) as-deposited TiC0.5 2) as-deposited (Ti0.3Al0.7)C0.5, and 3) (Ti0.3Al0.7)C0.5 after heat treatment in high vacuum at 600 oC. b) Change in the amount of amorphous carbon (C-C) as a function of Al content (x). With increasing x the amount of C-C increases (triangle), while the amount of car-bide (Me-C) decreases (square). A post annealing can further increase the amount of C-C (hourglass) [V].

7.3 Mechanical properties of the MAX-phases and nanocomposites

In Fig. 2.3 the crystal structure of the 211, 312 and 413 MAX-phase was shown. As can be seen there are two types of M-atoms in the 312 and 413 phases; one that bind to X and A (M1), and one that bind solely to C (M2). This implies that the M1 and M2 atoms have different roles in the chemical bonding and gives rise to an anisotropic chemical bond in the MAX-phases. There are several theoretical and experimental studies on the electronic structure of MAX-phases in the literature. In [III] we present soft X-ray spectroscopy and theoretical calculations for Ti3AlC2 and Ti2AlC MAX-phases. From the areas in the BCOOP-curves of Ti2AlC, Ti3AlC2 and TiC shown in Fig. 7.14 it is obvious that the Ti 3d – C 2p bond is stronger than the Ti 3d – Al 3p bond in both Ti2AlC and Ti3AlC2. Furthermore, the Ti-C bond for Ti1 is stronger than for Ti2 in the case of Ti3AlC2. In Ti2AlC where only Ti1 is present the Ti-C bond is stronger than in pure TiC. Finally, the Ti-Al bond is slightly weaker in Ti2AlC than in Ti3AlC2. The anisotropy in the chemical bonds is the foundation for the observed basal-plane slip and kink formation in MAX-phases.

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Figure 7.14. Calculated balanced crystal overlap population (BCOOP) of TiC, Ti2AlC, and Ti3AlC2 [III].

The mechanical properties of the deposited Ti2AlC, Ti3AlC2 and V2GeC MAX-phase films where studied by nanoindentation. The Ti2AlC and Ti3AlC2 films exhibited similar hardness and Young’s modulus of ~20 GPa and ~250 GPa, respectively. These values are in the range of other MAX-phase systems under basal plane indentation [44]. In our theoretical results, discussed above, it was shown that the Ti-Al bond is somewhat weaker in Ti2AlC than Ti3AlC2. Therefore, a lower Young’s modulus can be expected for Ti2AlC compared to Ti3AlC2. However, this is difficult to deduce from the experimental nanoindentation due to a scattering of the data points. Fur-thermore, in [I] the bulk modulus of Ti3AlC2 and Ti2AlC to 187 GPa and 161 GPa, respectively was calcualted. These values are in agreement with the literature and are indicative for the hardness of materials [79] and suggests that Ti3AlC2 may have a higher hardness than Ti2AlC. However, the hard-ness of a compound is determined by the resistance of bonds for distortions (i.e. the bulk modulus) and how easily dislocations can move in the system. Therefore, it is not surprising to observe a deviation between theoretical results and experimental film where, e.g., inclusions (see Figs. 7.5 and 7.6) are present and interfere with the dislocation motions.

In contrast to the Ti-Al-C system, the hardness and Young’s modulus of the V2GeC film were lower and measured to ~7 GPa and ~190 GPa, respec-tively (see [II]). However, Sun et al. have reported on the bulk modulus of V2GeC (~215 GPa) and Ti2AlC (166 GPa) [79, 85]. The higher bulk

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modulus of V2GeC is explained by a higher valence electron population for V and therefore more charge is placed in the V-C bond than in the Ti-C bond. Considering the error range in the experimental measurements the hardnesses of V2GeC and Ti2AlC are in the same range. Furthermore, as was mentioned above and will be further discussed in the following paragraph the inclusion content also influences the mechanical measurement of V2GeC.

The nanolaminated MAX-phases clearly have a very special deformation behavior with kink formation as was shown in Fig. 5.5. An intriguing ques-tion is if this mechanism can be suppressed. In [IV] a new type of intrusion deformation in multilayered Ti3SiC2/TiC0.67 films was demonstrated. By interleaving the Ti3SiC2 MAX-phase with the less ductile TiC0.67 into a mul-tilayered epitaxial structure kink formation could be suppressed. Figure 7.15 shows a TEM micrograph of such a multilayered structure with a period of ~20 nm after indentation.

Figure 7.15. STEM image after indentation of a multilayered Ti3SiC2/TiC0.67 film with a period of ~20 nm. In comparison to Fig. 5.5 no clear kinks are observed. Instead the whole structure ha collectively bent. In similarity with Fig. 5.5 delamina-tion cracks are running in parallel with the substrate [IV].

As can be seen that the film has bent collectively and there is no pronounced kink or pile-up formation. However, material has flown from the indenter zone towards the sides of the indenter. It was concluded that the dominating deformation mechanism in Ti3SiC2 still is the basal-plane slip. This gave rise to a feather-like appearance where atomic sheets were forced into each other. (see Fig. 7.16).

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Figure 7.16. a) Unfiltered cross-sectional TEM image of Ti3SiC2/TiC0.67 multilayer. In the figure the Ti3SiC2 layers exhibit a feathering appearance from the intrusion deformation. Persistent slip lines are seen in the TiC0.67 layers. b) Deformed region shown in a) at higher magnification [IV].

Dark bands are also seen in the TiC-layers. These features arise from a per-sistent slip along the primary slip system (110) in the TiC compound. Edge dislocations were also observed in the TiC-layers (see Fig. 7.17).

Figure 7.17. Cross-sectional TEM image of TiC0.67 layer interleaved by Ti3SiC2 showing two domains in the TiC0.67 separated by an 18o deformation induced dislo-cation wall boundary as well as individual edge dislocations on the inclined (111) planes (marked with T symbols).

However, as evident from Fig. 7.16 a gross slip through out the film was hindered by the interleaving Ti3SiC2 layers and instead the multilayer deform via an intrusion-type mechanism evident from the feather-like appearance in Fig. 7.16. The hardness and Young’s modulus were measured to ~15 GPa and 250 GPa, respectively for all bilayer period. These values correlate with the rule of mixture for TiC0.67 and Ti3SiC2. Although the ambition with [IV] was to suppress the kink formation and study the deformation mechanism it

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is interesting to consider the lack of any hardness enhancement. The prob-able explanation is as follows: the dislocation motion observed in the TiC-layers is hindered at the Ti3SiC2 interface. However, this builds up a stress concentration at the interface, which nucleates new dislocations in the Ti3SiC2 layer. Since the basal planes are the active slip system in the Ti3SiC2, which also are parallel to the substrate, they will never encounter an interfer-ing TiC0.67 interface and the mechanical data of the multilayers will correlate with the rule of mixture.

Mechanical properties have also been studied for a epitaxial Ti3AlC(lll) perovskite and (Ti,Al)C nanocomposites (see [I] and [VI]). The microstruc-ture and crystallinity of these materials are different. However, they are in-teresting to relate to each other in order to consider the influence of Al. The crystal structure of Ti3AlC is shown in Fig 2.1. The structure can be seen as an ordered solid solution of Al into TiC. For the perovskite the hardness was measured to ~11 GPa. This is considerably lower than the measured value of ~24 GPa for epitaxial TiC0.67(lll). The lower hardness of Ti3AlC in compari-son to TiC0.67 is explained by the higher amount of C vacancies and the higher amount of weak Al-C bonds. The low tendency for Al-C bond forma-tion is seen in the electron charge density plots shown in Fig. 7.11. Also, in the nanocomposite (Ti,Al)C a lower hardness in comparison to TiC was observed. This is explained by the weakening of the bond as described above. However, also the nanocomposite structure with (Ti,Al)C crystals embedded in amorphous carbon contribute to the mechanical properties. Such composites can deform via grain boundary rotation and grain boundary sliding depending on elemental composition and volume ratio of the crystal-line and amorphous phase. This deformation gives rise to tough materials. In comparison with TiC, the elastic recovery of (Ti,Al)C was about 10% higher (see [VI]).

7.4 Tribological properties of the films For low friction properties it is beneficial with an easily sheared material in the tribological contact. The nanolaminated structure of the MAX-phases with weak M-A bonds suggests that these materials may exhibit very good low-friction properties. However, this was not observed in pin-on-disc tests of, e.g., Ti3AlC2 MAX-phase films [84]. Furthermore, also other studies of nanolaminated MAX-phases have shown that these materials have a limited wear-resistance [43]. Possibly these properties can be improved by, e.g., an appropriate interlayer but this remains to be demonstrated.

On the other hand, nanocomposites are a group of materials that is known to exhibit good tribological properties. In section 7.2, we demonstrated that a solid solution of a weak carbide-forming element into the carbide structure would give a metastable system with a driving force for diffusion of carbon

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out from the carbide under the formation of free carbon or graphite. This means that such “alloyed” carbides have a strong tendency to form graphitic surface layers as a response to heat and/or external pressure. These graphitic overlayers can act as an easily sheared material and reduce friction and al-loyed carbides with weak carbide formers could therefore be excellent low-friction materials. This was also demonstrated in a series of papers ([V], [VI] and [VII]) and this concept has also been included in a patent.

Figure 7.18 shows the evolution of the friction coefficient of TiC and nanocomposite (Ti,Al)C films, and their ability to reform a tribofilm. As can be seen in Fig. 7.18 the friction coefficient of the Al-containing films is con-tinuously decreasing to final values of ~0.1 and ~0.05 depending on the C/Me-ratio. This is explained by the formation of a carbon-rich tribofilm where surface graphitization plays an important role. In comparison, the TiC films with similar C/Me-ratios exhibit a friction coefficient of ~0.22.

Figure 7.18. Evolution of friction coefficient as a function of the number of revolu-tions for A = TiC1.5, B = TiC0.7, C = (Ti0.5,Al0.5)C0.7 and D = (Ti0.5,Al0.5)C1.2. The films should be compared in groups of two with similar C/Me-ratio, i.e., A is com-pared with D and B is compared with C. The ball was replaced with a new one every 3000 revolution and re-positioned in the wear track [V].

Figure 7.19 shows C1s XPS spectra of the wear track of (Ti0.5Al0.5)C0.7 and TiC0.7.

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Figure 7.19. XPS C1s spectra from the tribofilms in the wear tracks from a) TiC0.7 and b) (Ti0.5Al0.5)C0.7. All spectra were recorded after 90 s Ar-ion pre-sputtering [VI].

As can be seen the amount of C-C bonds at ~284 eV has increased signifi-cantly for the Al-containing film compared to both the as-deposited film and to the wear track of TiC0.7. From the curve fit areas the C-C content of the Al-containing film had increased from ~20 to 55% in the wear track. Also a minor C-C signal was observed in the wear track of TiC0.7. The C-C ob-served by XPS in the wear tracks was indeed of graphitic nature, which was evident from micro Raman analysis. Figure 7.20 shows Raman spectra of different regions of the (Ti0.5Al0.5)C0.7 and TiC0.7 films. Clear D and G bands are observed from the analysis of the adhered material in the wear track of the Al-containing film. Also the wear track region without adhered material showed an increase in graphite-like content in comparison to the as-deposited film surface. In contrast, the adhered material of the wear track of the TiC0.7 film only exhibited a weak signal attributed to graphite-like mate-rial. XPS analysis also showed that the oxidized Ti and Al had been formed in the wear track.

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Figure 7.20. Raman spectra for as-deposited and wear track tribolayer and from additional adhered material for a) TiC0.7 and b) (Ti0.5Al0.5)C0.7. Solid lines indicate the position of the D and G bands [VI].

The addition of Al to TiC0.7 also strongly influenced the microstructure. For pure TiC0.7 a highly crystalline material was observed, while the addition of Al changed the film microstructure to nanocrystalline. Such change in struc-ture has also a strong effect on the mechanical properties where nanocrystal-line materials may be tough and withstand overloads without catastrophic failure in, e.g., tribological contacts. The hardness was reduced from ~36 GPa to ~23 GPa by the addition of Al. The Young’s modulus also decreased from ~294 GPa to ~230 GPa. The change in mechanical properties and mi-crostructure by the addition of Al also improved the elastic recovery. We observed an elastic recovery of ~35 % for the pure TiC0.7 film, while it in-creased to ~45 % for (Ti0.5Al0.5)C0.7. The differences in mechanical response were also mirrored in the crack formation of the films where Al increased the scratch resistance and made the (Ti0.5Al0.5)C0.7 film less brittle than the TiC0.7. Consequently, the Al containing film should withstand overloads to a

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higher extent. In the pin-on-disc measurements film surfaces will inevitably suffer from wear. For the TiC0.7 and (Ti0.5Al0.5)C0.7 films we observed a simi-lar wear rate, although the large difference in mechanical properties. For both films worn material was observed outside the wear tracks of the films as-well as on the counter surfaces. Figure 7.21 shows, as an example, an EF-TEM image of the counter surface that has run against the Al-containing film. In Fig. 7.21 dark and bright areas are seen in the tribomaterial. The bright regions are amorphous material, which according to STEM/EDX analysis consist of a significant amount of free C-C. It is likely that this bright areas consist of a graphite-like material, but this will be further ex-plored in a near future.

Figure 7.21. Cross-sectional TEM micrograph of the counter surface (Ti0.5Al0.5)C0.7. The brighter regions correspond to carbon enriched amorphous material.

In the micro-Raman analysis only a graphite-like material was observed on the counter surface that had run against the Al-containing film, while also oxides were present for the counter surface to TiC0.7 (see Fig. 7.22). In sum-mary, the Al-containing film generates more C-C in the contact in the pin-on-disc tests than the TiC0.7 film.

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Figure 7.22. Raman spectra of different regions on the counter surfaces for (Ti0.5Al0.5)C0.7 (bottom) and TiC0.7 (top) [VII].

Also the Ti-Fe-C system has been studied with the emphasis to also to add magnetic properties to the films. In similarity with the Ti-Al-C system a formation of graphite-like material was observed. Friction values as low as for the nanocomposite (Ti,Al)C films were not reached due to a bad film/substrate adhesion. However, by optimization of the deposition process it is likely that low friction values will be attained.

In general, nanocomposite films are beneficial for tribological contacts since the possibilities of materials design are large. By the addition of weak-carbide forming element self-lubricating mechanisms can be achieved as-well as a wide range of mechanical properties. Consequently, such films can be optimized for the prevailing contact situation, but also have adaptive properties.

7.5 Electrical properties of the films The electrical properties of MAX-phases depend on the stoichiometry and materials system. The conductivity is strongly metallic and anisotropic. In the MAX-phases there are two types of M-atoms with different chemical surroundings. In 312 and 413 one M-atom bonds to X and A (M1), and one

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bond solely to C (M2). In contrast the 211 phase only have M-atoms of M1 type. The different chemical surrounding of the M-atoms give them different properties. For example, the M1 atom has a greater portion of metallic bonds that contribute to the electrical conductivity. Since the 211 phase only have M1 bonds it implies that its resistivity is lower than for the 312 and 413 phases. One can say that the 211 phase is more “metallic-like”. The lower resistivity of 211 phases has indeed been observed experimentally. For ex-ample, Ti3GeC2 and Ti2GeC have a resistivity of 50 ��cm and 15��cm, respectively [86]. The resistivity can also be modified by changing the M-atom. For example, if Ti is replaced by Cr in Ti2AlC to form Cr2AlC the number of d-electrons are increased and thereby the number of states at the Fermi level [87]. Also the A-element influences the electrical properties. For example, is the density of state increased from 3.72 to 4.65 states/(eV cell) by replacing Al for Ge in Ti3AC2 (A = Al or Ge) [88].

The resistivity of the Ti3AlC2, Ti2AlC and V2GeC MAX-phase films have been measured to 51 ��cm, 44 ��cm and 21 ��cm, respectively (see [I] and [II]). The higher resistivity of Ti3AlC2 in comparison to the Ti2AlC is consistent with the greater portion of metallic bonds in Ti2AlC. The even lower resistivity of V2GeC is understood from the larger number d-electrons in V compared to Ti. As described above this increases the number of states at the Fermi level, which improves the conductivity. Emmerlich et al. have shown that inclusions in the MAX-phase constricts the current flow [86]. In the V2GeC films inclusions of VCx/V8C7 were observed (see sec. 7.1). Pos-sible a lower resistivity can be achieved if the amount of these can be re-duced.

The good electrical properties of the MAX-phases make them interesting as contact material. However, the deposition temperature is typically of 450-1000 oC, which is too high for engineered contact materials. On the other hand at low-temperatures (~300 oC) nanocrystalline materials can be grown. Such materials may also exhibit good electrical properties. We have, e.g., measured the resistivity for V2GeC films deposited at 300 oC to 98 ��cm. As was discussed in sec. 7.2 this film consists of X-ray amorphous VCx and V2C together with polycrystalline Ge. On the other hand films deposited under same conditions but without Ge constitutes of polycrystalline VCx. For such films the resistivity was measured to 180 ��cm. It is a surprising result that an amorphous matrix including semiconducting particles can exhibit a lower resistivity than a more crystalline structure with conductive carbides. Further studies are required to clarify this observation.

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7.6 Magnetic wear-resistant Ti-Fe-C films The concept of doping with a WCF metal can be extended to include not only good tribological films. By the addition Fe or Ni also magnetic films can be synthesized. The magnetic characterization of these films shows that the properties can be modified by tuning the composition of the as-deposited films or carry out a heat-treatment step. Figure 7.23 shows the field depend-ence of the magnetization at 10 K for the as-deposited and heat-treated (Ti1-xFex)Cy films studied in [VII].

Figure 7.23. Field dependence of magnetization for a) as-deposited films and b) corresponding results after heat treatment. The inset shows a blow of the region close to the coercive field [VII].

As can be seen that the saturation magnetization increases for decreasing carbon contents. Furthermore, by a post heat treatment step the magnetiza-tion saturation can be increased for higher carbon contents. This increment is explained by the formation of �-Fe particles in the matrix of (Ti1-xFex)Cy. The magnetic characterization also shows that we have an ordered magnetic (Ti1-xFex)Cy phase for the lowest carbon contents.

A formation of graphite-like material was indeed observed in pin-on-disc measurements. However, due to a high wear, explained by a bad film adhe-sion friction coefficients as low as in the the (Ti,Al)C system were not at-

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tained. These results can most probably be improved by a proper pre-etch step and by optimizing the interlayer.

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Future outlooks

The papers included in this thesis give answers and gives rise to many new questions. For example, the temperature dependence for the nucleation of MAX-phases has been discussed, but there is no clear general trend what actually determines the lowest deposition temperature. A better knowledge can probably be attained if the M- and A-elements are changed in a system-atic way. For example, it would be interesting to compare the nucleation temperature Ti2GeC and V2GeC in comparison to Ti2AlC and V2AlC. Such deposition series would give knowledge about the influence of more d-electrons (i.e. changing from Ti to V) for different A-elements (i.e. Al and Ge). Furthermore, intergrown MAX-phase structures are probably easily achieved in the V-Ge-C by increasing the deposition temperature from 850 oC to 1000 oC.

The multilayered Ti3SiC2/TiC0.67 would be thrilling to study further. It would, e.g., be interesting to deposit films where one of the bilayers are thicker and explore their mechanical properties. Furthermore, their tribologi-cal performance would be interesting to explore.

Films with an X-ray amorphous carbide structure were shown to exhibit good electrical properties. However, today it is not clear why such films have lower resistivity than more well-crystallized carbide films. It would be interesting to further explore the influence of the amorphous microstructure on the electrical properties.

Furthermore, very good tribological properties were achieved for nano-composite (Ti,Al)C films by the formation of C-C carbon. In the literature there are examples of other composite system such as TiN/Cu where very high hardness values have been achieved. Therefore, it would be interesting to further explore our nanocomposite (Ti,Al)C films to see if the volume ratio of the carbide and amorphous phase can be optimized to achieve high hardness values. Furthermore, in combination with a good tribological per-formance, also the magnetic properties can be modified. This was shown for (Ti,Fe)C films. The magnetic properties were changed by the C/Me-ratio, composition and post heat treatment. By the heat treatment particles of �-Fe was formed in the matrix. These particles are randomly oriented. However, if the formation (shape, elongation and position in the film) of these particles can be controlled, even better magnetic properties may be achieved.

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Acknowledgement

Ett stort och varmt tack till min handledare Prof. Ulf Jansson som med sin oändliga tålmodighet, kreativitet och kunskap, inspirerat och motiverat mig under åren vid Ångström. Tack! Mattias, jag ser framemot din disputationsdag då vi kan börja umgås utan ha pågående forskningsprojekt hängade över oss. Jag hoppas vi behåller vår vänskap livet ut. Erik L, jag har uppskattat att ha dig som arbetskollega och likaså alla de gånger du dykt in på mitt rum och diskuterat forskning och dryftat mer världsliga ting. JP tack för introduktionen i tunnfilmernas värld. Martin M, tack för alla intressanta och fina mätresultat från MAX-lab. Er-nesto och Hans, tack för all konsulthjälp vid TEM’et. Tack till Janne och Anders för all er kunskap och hjälp med verkstadstekniska problem. Tack till Olle E, Mikael R, Biplab och övriga teoretiker för trevligt samarbete och en djupare förståelse av forskningsresultat. Tack till Peter S och Stojanka för all magnetisk karakterisering. Tack till Mikael O för hjälp med av röntgenut-rustningen. Tack till alla på institutionen för Materialkemi. Tack till alla trevliga och hjälpsamma doktorander och seniorer på institutionen för Mate-rialvetenskap. Under mina år vid Ångströ

synthesis and characterization of ternary carbide thin films170841/fulltext01.pdf · 2009. 2....

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