carbide and max-phase engineering by thin film synthesis163986/fulltext01.pdfcarbide and max-phase...

Comprehensive Summaries of Uppsala Dissertationsfrom the Faculty of Science and Technology 930

Carbide and MAX-PhaseEngineering by

Thin Film Synthesis

JENS-PETTER PALMQUIST

ACTA UNIVERSITATIS UPSALIENSISUPPSALA 2004

Printed in Sweden by Universitetstryckeriet, Uppsala 2004

-

Till Professor Baltazar

Carbide and MAX-Phase Engieneering by Thin Film Synthesis

Uppsala University i

List of publications This thesis is based on the following publications, chronologically numbered and referred to in the text by their Roman numerals.

I. Deposition of epitaxial ternary transition metal carbide films. J.-P. Palmquist, J. Birch, U. Jansson, Thin Solid Films 405 (2002) p. 122.

II. Magnetron sputtered W-C films with C60 as carbon source. J.-P. Palmquist, Zs. Czigany, M. Odén, J. Neidhart, L. Hultman, U. Jansson, Thin Solid Films 444 (2003) p. 29.

III. Epitaxial growth of tungsten carbide films using C60 as carbon precursor. J.-P. Palmquist, Zs. Czigany, L. Hultman, U. Jansson, Journal of Crystal Growth 259 (2003) p. 12.

IV. Magnetron sputtered epitaxial single-phase Ti3SiC2 thin films. J.-P. Palmquist, U. Jansson, T. Seppänen, P. O. Å. Persson, J. Birch, L. Hultman, P. Isberg, Applied Physics Letters 81 (2002) p. 835.

V. Kink formation around indents in laminated Ti3SiC2 thin films studied in the nanoscale. J. M. Molina-Aldareguia, J. Emmerlich, J.-P. Palmquist, U. Jansson, L. Hultman, Scripta Materialia 49 (2003) p. 155.

VI. New MAX phases in the Ti-Si-C system studied by thin film synthesis and ab initio calculations. J.-P. Palmquist, S. Li, P. O. Å. Persson, J. Emmerlich, O. Wilhelmsson, H. Högberg, M. I. Katsnelson, B. Johansson, R. Ahuja, O. Eriksson, L. Hultman, U. Jansson, Submitted to: Physical Review B (2003).

VII. Growth of Ti3SiC2 thin films by elemental target magnetron sputtering. J. Emmerlich, J.-P. Palmquist, H. Högberg, Zs. Czigány, Sz. Sasvàri, P. O. Å. Persson, U. Jansson, L. Hultman, In manuscript (2004).

VIII. Sputtering of Ti2AlC and Ti3AlC2 thin films. O. Wilhelmsson, J.-P. Palmquist, T. Nyberg, U. Jansson, Submitted to: Applied Physical Letters (2004).

IX. Electronic structure of the MAX-phases Ti3AC2 (A=Al, Si, Ge) investigated by soft x-ray absorption and emission spectroscopies. M. Magnuson, J.-P. Palmquist, S. Li, M. Mattesini, R. Ahuja, O. Eriksson, O. Wilhelmsson, P. Eklund, H. Högberg, L. Hultman, U. Jansson, In manuscript (2004).


ii Jens-Petter Palmquist

Publications not included in this thesis:

X. The formation and characterization of epitaxial titanium carbide contacts to 4H-SiC. S.-K. Lee, E. Danielsson, C.-M. Zetterling, M. Ostling, J.-P. Palmquist, H. Hogberg, U. Jansson, Materials Research Society Symposium - Proceedings 622 (2000) T691.

XI. Electrical characterization of TiC ohmic contacts to aluminum ion implanted 4H-silicon carbide. S.-K. Lee, C.-M. Zetterling, E. Danielsson, M. Ostling, J.-P. Palmquist, H. Hogberg, U. Jansson, Applied Physics Letters 77 (2000) p. 1478.

XII. Low resistivity ohmic titanium carbide contacts to n- and p-type 4H-silicon carbide. S.-K. Lee, C.-M. Zetterling, M. Ostling, J.-P. Palmquist, H. Hogberg, U. Jansson Solid-State Electronics 44 (2000) p. 1179.

XIII. Low temperature epitaxial growth of metal carbides using fullerenes. U. Jansson, H. Hogberg, J.-P. Palmqvist, L. Norin, J. O. Malm, L. Hultman, J. Birch Surface and Coatings Technology 142-144 (2001) p. 817.

XIV. Low resistivity ohmic contacts on 4H-silicon carbide for high power and high temperature device applications. S.-K. Lee, C.-M. Zetterling, M. Ostling, J.-P. Palmquist, U. Jansson, Microelectronic Engineering 60 (2002) p. 261.

XV. Structural characterization of epitaxial Ti3SiC2 films. T. Seppänen, J.-P. Palmquist, P. O. Å. Persson, J. Emmerlich, J. Molina, J. Birch, U. Jansson, P. Isberg, L. Hultman, Conference Proceeding: the 53rd annual Meeting of The Scandinavian Society for Electron Microscopy (SCANDEM), ISSN 1455-4518, Tampere, Finland, (2002), p. 142

XVI. Growth and characterization of MAX-phase thin films. H. Högberg, L. Hultman, J. Emmerlich, T. Joelsson, P. Eklund, J. M. Molina-Aldareguia, J.-P. Palmquist, O. Wilhelmsson, U. Jansson, Conference Proceedings: The 4th Asian-European International Conference on Plasma Surface Engineering (AEPSE), Jeju, Korea, (2003). (To be published in Thin Solid Film or Surface and Coatings Technology.)

Also not included in this thesis is the following patent.

XVII. Method of synthesising a compound of the formula Mn+1AXn, film of the compound and its use. J.-P. Palmquist, U. Jansson, J. Birch, L. Hultman, P. Isberg, T. Seppänen, International publication number: WO03046247 (5 June 2003).


Uppsala University iii

Comments on my participation In the papers and manuscripts that I am first author, Publication I,-IV and VI, I have performed the experimental work, most of the characterisation and written the major part of the manuscripts. In the deformation study in Publication V, I have deposited the films and taken part in planning and discussion. In Publication VI, I have taken part in planning and discussion of the theoretical study, but not performed the calculations. In Publication VII and VIII, I have taken part in planning, experimental work and some characterisation and contributed to the discussion of writing. Publication IX is a soft x-ray study on the MAX-phase films where I have made some of the films and taken part in the discussion and writing. Publication X-XVI, are mentioned since I have been part of the thin film depositions and also performed some characterisation. They also deal with an important possible application of epitaxial carbides as a contact material in SiC-based microelectronics. All electron microscopy studies are made in cooperation with the Thin Film Physics Group, IFM at Linköping University. The co-authors are greatly acknowledged for the opportunity I have had to work with you.


iv Jens-Petter Palmquist

Contents

List of publications.....................................................................................................iComments on my participation................................................................... iii

Chapter 1. Introduction .....................................................................................1 1.1 Motivation and background .........................................................11.2 Materials engineering...................................................................2

Chapter 2. Carbides and MAX-phases .............................................................42.1 Transition metal carbides.............................................................42.2 The Mn+1AXn phases ....................................................................6

Chapter 3. Thin films .........................................................................................93.1 Synthesis techniques ....................................................................93.1.1 Evaporation .............................................................................9 3.1.2 Sputtering ..............................................................................103.2 Film growth................................................................................113.2.1 Epitaxial growth....................................................................123.2.2 Pseudomorphic growth..........................................................14

Chapter 4. Characterisation ............................................................................164.1 X-ray diffraction ........................................................................164.1.1 Thin film diffraction..............................................................164.1.2 Reciprocal space mapping (RSM).........................................184.2 Transmission electron microscopy.............................................194.3 X-ray photoelectron spectroscopy .............................................204.4 Four point probe measurements.................................................214.5 Nanoindentation.........................................................................224.6 Soft x-ray spectroscopy..............................................................22

Chapter 5. Theoretical calculations ................................................................24

Chapter 6. Results and discussions .................................................................266.1 Deposition of transition metal carbides......................................266.1.1 Tungsten carbides..................................................................266.1.2 Ternary carbides....................................................................286.1.3 Epitaxial TiC growth by sputtering from Ti and C targets....306.2 MAX-phase thin film design......................................................31


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6.2.1 Ti3SiC2 thin film growth .......................................................326.2.2 New MAX-phases in the Ti-Si-C system..............................346.2.3 Tin+1(Al1-xSix)Cn and A-element bonding character ..............376.3 Engineering of mechanical properties........................................386.3.1 Precipitation and solid solution hardening ............................386.3.2 Hardness enhancement by phase stability tuning..................406.3.3 Mechanical properties of nanolaminate MAX-phases ..........416.3.4 Designed nanolaminate MAX-phase structures ....................426.4 Artificial thin film structures......................................................446.4.1 Epitaxial gradient films .........................................................446.4.2 MAX-phase multilayer structures .........................................45

Chapter 7. Afterthoughts .................................................................................47

Appendix 1 – Abbreviations and acronyms..........................................................48

Appendix 2 – A possible new phase: Ti2Si ............................................................49

Appendix 3 – Ti-Si-C ternary phase diagram ......................................................51

Appendix 4 – New intergrown MAX-phase structures........................................52

Acknowledgements -- Tack!...................................................................................54

References ...............................................................................................................56

Sammanfattning på svenska ..................................................................................61

Karbid och MAX-fas design med tunnfilmssyntes...............................................62Tunna filmer ...............................................................................................62Förångning, sputtring och filmtillväxt i vakuum ........................................62Kristallinitet och epitaxi .............................................................................64Karbider är hårda........................................................................................65Design av karbider med tunnfilmstillväxt ..................................................65Design av spänningsfria filmer och gradienter ...........................................65Titanstabiliserad epitaxiell volframkarbid ..................................................66Kontrollerad hårdhet...................................................................................66En MAX:ad lagerstruktur ...........................................................................66Forskningssamarbete mellan universitet och industri.................................67En ny MAX-fas och två nya MAX-strukturer ............................................68Forskningssamarbeten inom universitet .....................................................68

1

Chapter 1.

Introduction

1.1 Motivation and background Throughout history, mankind has been driven by an irresistible wish for new and improved solutions to technological applications. The need for better performance, higher profit or simply a curious human mind has pushed the technological development forward. To surpass previous levels of technology, sometimes, new ideas require new materials, sometimes, new materials breed new ideas. This thesis is a work of materials engineering with the purpose to explore the possibilities of controlled carbide film growth. Part of this work can be considered as a search for new materials sprouted from theoretical predictions, or simply wild ideas. Other parts are better viewed as finding and characterisation of new materials that are just waiting for the proper idea of application.

I would like to begin this thesis with a brief review of some important research, performed earlier within our group, which builds the foundation to this thesis and provided the tools for this study. Somewhere in the beginning (from my point of view) L. Norin studied the existence of transition metal fullerides (MexC60)1. It was found that C60 decomposes in the presence of metals and that they together form carbides. By the use of electron-beam evaporation of the metal with evaporated C60as carbon precursor, films of TiC, VC and NbC were grown1,2. Epitaxial TiC films could be deposited at the surprisingly low substrate temperature of 250 oC. Other authors had rarely reported epitaxial carbide growth earlier and it was therefore an unexplored field. Later, growth of epitaxial multilayers of TiC/VC pointed towards the possibility to design epitaxial carbide structures. TiC films on SiC were used to fabricate Ohmic contacts and Schottky diodes of excellent quality1,2. H. Högberg introduced magnetron sputtering of the metal and showed that it was possible to make epitaxial carbide superlattices and the temperature for epitaxial TiC film growth was further pushed down to 100 oC2.

When I started as a Ph.D. student, epitaxial carbide growth was well established in our group and I took up the task to further explore the possibilities of this controlled growth process. Could this approach be used to synthesise ternary carbides? Could

Chapter 1. Introduction

2 Jens-Petter Palmquist

the low synthesis temperature be used to stabilise high-temperature or metastable phases? What about tungsten carbide? And, what would happen with sputtered carbon instead of evaporated C60? There were also interesting theoretical papers that predicted random growth behaviour, possible phase-stabilisation by additives and chemically controlled hardness that I wanted to experimentally follow up3-5.

That is where it all begun, but not where this thesis ends. Yes, ternary carbides was a piece of cake, it was even possible to design epitaxial gradient films. The W-C system proved to be intriguing and full of delightful surprises, such as superhard tungsten-films and Ti-stabilised -WC1-x. And indeed, magnetron sputtering of carbon took the magic from C60 and proved to be an even better precursor for epitaxial carbide growth. Halfway through my Ph.D. student time a new compound emerged, Ti3SiC2 a so-called MAX-phase, with very persuasive arguments of useful properties6. This spurred my research towards the controlled design of epitaxial nanolayered MAX phase films. Both with the finding of new structures but also with controlled tuning of properties.

1.2 Materials engineering With this thesis, I want to put my publications into a larger perspective of materials engineering, which is what I have been doing these last years. A quick look in an encyclopedia gives the following definition of engineering:

Engineering is the application of science to the needs of humanity. This is accomplished through knowledge, mathematics and practical experience applied to the design of useful objects or processes.7

Design means, “to shape”. This implies that materials engineering is the ability to shape the material and the power to tune the properties. The properties of a material can be controlled by the microstructure with physical control over the process conditions or, on an atomic level, by changing the chemical bonding8. The physical aspect, cover topics like grain size control, artificial structures such as composites or laminates, it can also be precipitations by ageing, etc. The chemical aspect means control by phase composition, alloying, solid solutions, additives, etc. This classification is of course not stringent and is often implemented all together in engineering. A chemical additive may change the microstructure, or a precipitation can be seen as a change in phase composition. However, depending on the approach for materials design, this may be a relevant taxonomy.

I have used thin film synthesis as a tool for engineering. This approach is advantageous to use for controlled processing for several reasons. First of all, it gives the possibility to work with almost pure sources of the elements with very small influence from contaminants. The synthesis is a continuous growth process where the relative amount of the participating elements can be controlled individually and also changed during synthesis. The substrate temperature gives


Uppsala University 3

kinetics control over the process, and synthesis is possible far from thermodynamical equilibrium. Furthermore, the substrate can be used as a crystalline template for the film. All in all, this enables both chemical design and physical control over the microstructure in the synthesised material. Figure 1.1 gives a brief overview of how thin film synthesis have been used in this thesis to cover different aspects of engineering. Throughout the thesis, I will use the taxonomy introduced above and discuss the results from both a chemical and a physical point of view.

Figure 1.1. Mind map survey of this thesis work: Design of material properties can be performed from two viewpoints, physical control over microstructure or a chemical tuning on an atomic level. The recount of the results in chapter 6 is more of a jigsaw puzzle, rather than following a main theme. I hope this map may give the reader a safe road through the results. Roman numerals refer to each publication, where more can be found about the specific topics.

Physicalcontrol

Chemical control

Artificial structuresI Epitaxial gradient films VI Multilayers

Laminate structuresIV-VIII MAX-phases

PrecipitationII -WC1-x in W2C

matrix

Grain sizeII Controlled by

temperature

CompositesII -WC1-x in

amorphous carbon

AlloysTernary carbides: I (Ti,V)C, (Nb,Mo)C III (Ti,W)C VIII (Ti,Al)C

AdditivesIII -WC1-x stabilised

by Ti

Solid solutionII Martensite formation

due to solution of C in -W

Metastable growthI cubic MoC II -WC1-x, W2CVI Ti4SiC3VI Ti5Si2C3, Ti7Si2C5

Pseudomorphic growthIII Epitaxial -WC1-xIV-VIII epitaxial MAX-films

Phase compositionI Valence electron concentration

in (Nb,Mo)C films with random stacking of the Me-layers

VI Intergrown compound of known MAX-phases

Materialsengineering

4

Chapter 2.

Carbides and MAX-phases

2.1 Transition metal carbides Main part of this thesis work is made on transition metal carbides of group 4 to 6 in the periodic table. These are interstitial carbides where the structure can be viewed as metal atoms arranged in a closed-packed stacking with carbon occupying the octahedral interstitial sites9-11. The bonding in transition metal carbides is often described as a mixture of metallic, covalent and ionic3,12-14. This is also reflected in the properties that the carbides exhibit10,13. Due to a maintained metal-metal overlap, the carbide shows metallic properties such as good electric and thermal conductance. The formation of strong covalent bonds between the metal and carbon atoms gives stable compounds with high melting points and high hardness, but also chemical inertness and good resistance to oxidation. Since carbon is more electronegative than the metal, a net charge transfer from the metal to the carbon atoms gives the ionic contribution to the bonding.

Among the transition metal carbides there are three main crystal structures; the hexagonal Me2C, the cubic MeC1-x with rock salt (NaCl, B1) structure, and the hexagonal mono-carbide MeC, each of them represented in the W-C system as shown in figure 2.115. They differ in the stacking sequence of the Me-atoms and in C-occupancy, as indicated in the figure9. Some structural differences between the groups should be observed16. Group 4 elements only form the MeC1-x phase with extensive homogeneity ranges for the C occupancy. This phase is also observed in group 5, but with smaller homogeneity ranges. However, the MeC1-x phase is only stable at high temperatures in group 6. Group 5 and 6 (except Cr) elements also form the Me2C phase with high- and low-temperature modifications, distinguished by the arrangement of the carbon atoms. In group 6, both Mo and W form the stoichiometric MeC, which differs from the other carbides in that it has no homogeneity range. In addition, group 6 elements also form -phases, with Me4C3stoichiometry and there are other carbides with more complex crystal structures, especially for Mo and Cr. For example, the NiAs-structure, with the stacking AbAcBaBc and the -structure, with stacking AcBaCbAbCaBc, as in the figure, capital letters represents the Me-position and lower-case letter the C-position.



3000

2600

2200

1800

1400

1000

At% C0 10 20 30 40 50 60

oC

W-C system

bcc -W -WC1-x

-W2C

WC

Interstit ial siteStacking:

AcBaCbAcBaCb

Stacking: AbAbStacking: ABcABc

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At% C0 10 20 30 40 50 60

oC

W-C system

bcc -W -WC1-x

-W2C

WC

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At% C0 10 20 30 40 50 60

oC

W-C system3000

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oC

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At% C0 10 20 30 40 50 60

oC

W-C system

bcc -W -WC1-x

-W2C

WC

Interstit ial siteInterstit ial siteStacking:

AcBaCbAcBaCb

Stacking: AbAbStacking: ABcABc

Figure 2.1. In the W-C system, one finds the three main occurring carbide structures, Me2C,MeC1-x and MeC. For each structure, the stacking of atoms is indicated; capital letters represent the Me-position and lower-case letter the C-position. Pure tungsten has bcc (body-centered cubic) structure, and can only solve very limited amounts of carbon atoms at interstitial sites. Dark spheres are the Me-atoms and light spheres are the C-atoms.

The chemical and structural similarities among the transition metals allows for substitutions of the metal in the carbides. Ternary carbides, quaternary carbides and even higher combinations are frequently studied17. Complete solid solutions are found between compounds with the same crystal structure18. However, carbides of different crystal structure often only shows limited solubility18. Furthermore, carbon is commonly substituted for nitrogen19, boron20 or oxygen. This opens up for an unlimited chemistry of materials design with the possibility to tune the desired properties4,13,17.

The binary carbides are generally too hard and brittle to use as a construction material. However, their attractive properties can be used in combination with other compounds. One example is the cemented carbides (hard metal), which consist of fine-grained tungsten carbide in a matrix of cobalt metal16. Other carbides are also used in powder metallurgy as alloy components. There is also the iron-carbon system21. However, early transition metal carbides (and nitrides) are above all used as wear protective coatings22. It has long been known that a thin film of these materials can enhance the performance of the coated object. For example, the working lifetime of cutting tools can be extended with orders of magnitude23.Historically, TiC coatings was one of the first, soon followed by Ti(C,N) and TiN coatings (the one that looks like gold on drills and bits). Today (Ti,Al)N is also commonly used as wear protective coatings. However, the coating must of course be adapted for the specific application to obtain optimal performace24,25. This has lead to the development of numerous coating techniques and the design of optimised complex coatings. This includes concepts like multilayers26,27, superlattices28,29,functionally graded coating30-32, low-friction diamond-like carbon (DLC) films33 and so on. These concepts are all presented in literature and also widely implemented by the industry.

Chapter 2. Carbides and MAX-phases


The aim with the carbide studies within this work has been to further develop the present low-temperature deposition technique for epitaxial carbide film growth. The use of substrate temperatures far from thermodynamical equilibrium opens up for deposition of metastable phases. It can also facilitate deposition of ternary carbides and artificial structures since the low temperature limits segregations and phase separations. In particular, the W-C system has been studied, since it has been connected with unsolved problems, where deposition of phase pure films, especially for the WC phase, has rarely been reported. Single crystal substrates have been used to promote high crystalline film quality for further characterisation of structure, hardness, electrical properties and bonding character, etc. A low-temperature process could initiate deposition on substrates earlier not available for carbide coatings. This could be usable in new wear protective applications or give way for epitaxial carbide films in microelectronic devices, where the continuing miniaturisation process leads to higher demands on the participating compounds.

2.2 The Mn+1AXn phases Barsoum et al. introduced the family name MAX, referring to compounds with the general formula Mn+1AXn, n=1, 2 and 334. Where M is an early transition metal, the A-element is mostly from group 13-14* in the periodic table, and X is either C and/or N. Henceforth, based on the stoichiometry, these phases will be referred to as 211, 312, 413 etc. for n=1, 2, 3… respectively. Figure 2.2 shows where in the periodic table the MAX-phases can be found. There are about 50 different thermodynamically stable MAX-compounds34.

Sc Ti

Zr

Hf

V

Nb

Ta

Cr

Mo

C N

Cd

Ga

I

Tl

Al

Ge

Sn

Pb

Si

As

P S

H

Li

Na

K

Rb

Cs

Fr

Mg

Ca

Sr

Ba

Ra

Be

Y

Lu

Lr Unq Unp

W

Unh

Mn

Tc

Re

Uns

Fe

Ru

Os

Uno

Co

Rh

Ir

Une

Ni

Pd

Pt

Cu

Ag

Au

Zn

Hg

B

Sb

Bi

Se

Tl

Po

O

Br

I

At

Cl

F

Kr

Xe

Rn

Ar

Ne

He= early transition element= group 13-15 element= C or N

Periodic Table of Mn+1AXn Phases

M2AX (211): Ti2AlC, Nb2AlC, Ti2SnC…M3AX2 (312): Ti3SiC2, Ti3AlC2, Ti3GeC2M4AX3 (413): Ti4AlN3, Ti4SiC3

MAX

Sc Ti

Zr

Hf

V

Nb

Ta

Cr

Mo

Sc Ti

Zr

Hf

V

Nb

Ta

Cr

Mo

C NC N

Cd

Ga

I

Tl

Al

Ge

Sn

Pb

Si

As

P S

Cd

Ga

I

Tl

Al

Ge

Sn

Pb

Si

As

P S

H

Li

Na

K

Rb

Cs

Fr

Mg

Ca

Sr

Ba

Ra

Be

Y

Lu

Lr Unq Unp

W

Unh

Mn

Tc

Re

Uns

Fe

Ru

Os

Uno

Co

Rh

Ir

Une

Ni

Pd

Pt

Cu

Ag

Au

Zn

Hg

B

Sb

Bi

Se

Tl

Po

O

Br

I

At

Cl

F

Kr

Xe

Rn

Ar

Ne

HeH

Li

Na

K

Rb

Cs

Fr

Mg

Ca

Sr

Ba

Ra

Be

Y

Lu

Lr Unq Unp

W

Unh

Mn

Tc

Re

Uns

Fe

Ru

Os

Uno

Co

Rh

Ir

Une

Ni

Pd

Pt

Cu

Ag

Au

Zn

Hg

B

Sb

Bi

Se

Tl

Po

O

Br

I

At

Cl

F

Kr

Xe

Rn

Ar

Ne

He= early transition element= group 13-15 element= C or N

Periodic Table of Mn+1AXn Phases

M2AX (211): Ti2AlC, Nb2AlC, Ti2SnC…M3AX2 (312): Ti3SiC2, Ti3AlC2, Ti3GeC2M4AX3 (413): Ti4AlN3, Ti4SiC3

MAX

Figure 2.2. Periodic table of the Mn+1AXn-phase elements. M is an early transition metal, A is an A-group element mainly from group 13-14, and X is carbon or nitrogen.

* In the American nomenclature for the periodic table, group 13-14 is referred to as IIIA and IVA, thus the naming of the A-element.



The nominator for all MAX-compounds is the nanolaminate structure with a matrix of MX slabs interleaved by single layers of A-element. As can be seen in figure 2.3, all three compositions have a hexagonal unit cell and they all belong to space group

mmcPD h /6346 . For n=1, we recognise the H-phases

*; M2AX, characterised by Jeitschko, Nowotny, Toth and co-workers during the 1960-ies35-37. The H-phases have a layered hexagonal structure with M2X layers interleaved with single layers of A-element35. n=2 gives M3AX2, first identified in Ti3SiC238 and Ti3GeC239, that are structurally related to the H-phases, but having M3X2 layers between the A-layers. For the composition M4AX3 (n=4), only one compound has so far been described, namely Ti4AlN340,41 found by Procopio et al. 1999. This phase have layers of M4X3separating the A-layers.

= M

= A

= X

211 312 413

Mn+1AXn= M= A

= X

= M

= A

= X

211 312 413

Mn+1AXn

Figure 2.3. The crystal structure of all MAX-phases is made of transition metal carbide or nitride (MX) sheets that are interleaved with single layers of pure group 13-14-element (A). The composition is Mn+1AXn, where n=1, 2 or 3. Based on the relative number of M, A, and X-atoms in each structure, they fall into three classes referred to as 211, 312 or 413 compounds.

Ti3SiC2 is the best-characterised MAX-phase to date, but detailed data is also available for other representatives34. Thermally, elastically, chemically and electrically they share many of the advantageous properties of their respective stoichiometric binary carbides (or nitrides) due to the strong covalent M-X bond. However, mechanically, the MAX-phases distinguish themselves from their hard and brittle related MX compounds due to a relatively weak M-A bond. The anisotropic structure gives the MAX-phases a very special deformation behaviour where dislocations and glide only occur along the basal plane, i. e., the A-layer42,V.This makes the MAX-phases machinable, relatively soft, resistant to thermal chock

* H-phase = Hägg-phase. Originally, the Hägg-phases referred to transition metal compounds with a small atom on interstitial sites. However, during time, in some literature, the meaning of “H-phases” has been altered. H-phases are commonly used in the MAX-phase community as a family name for the M2AX compounds.

Chapter 2. Carbides and MAX-phases


and unusually damage tolerant for a compound with otherwise ceramic properties6.They are the only polycrystalline solids that deform by a combination of kink and shear band formation, together with delaminations of individual grains34.Furthermore, lateral force microscopy measurements show extremely low coefficient of friction with less than 0.005 on cleavage faces along the basal plane43. It is also interesting to note that Ti3SiC2 exhibit excellent cyclic oxidation resistance at elevated temperatures44. Most research on MAX-phases has been concentrated to bulk studies.

There are several areas where MAX-phase films could be of technological importance, for example as low-friction gliding contacts, self-lubricating wear surfaces, or as corrosive protective coatings. However, previous there was no deposition technique available for MAX-phase film growth. A significant part of this thesis focus on the development of a process for MAX-phase film deposition, with the presentation of a patented and successful approach for epitaxial growth at (fairly) low temperatures. The initial development was focused on Ti3SiC2 film growth. But, the studies were taken further to search for and design new MAX-phases in the Ti-Si-C system by taking advantage of the developed method. The Ti-Al-C system was chosen as second model system. Ti2AlC and Ti3AlC2 show similar mechanical properties as Ti3SiC2, but with superior high-temperature stability due to alumina formation on the surface instead of silicon oxide. Furthermore, we also wanted to establish a foundation for further studies of MAX-phase solid solutions in the Ti-(Al-Si)-C system. Even though the MAX-phases were discovered 30 years ago, it is only recently that they have attracted attention. Therefore, there are still areas of chemistry and bonding, as well as mechanical properties that is not quite understood or fully investigated. The possibility of high quality single-crystal MAX-phase films can be used for further investigations to gain deeper knowledge, especially when experimental results can be combined with theoretical treatment.

9

Chapter 3.

Thin films

3.1 Synthesis techniques There are numerous coating techniques, all with their special advantages depending on type of material to be deposited and on the application. However, for carbide thin film synthesis the two main approaches are chemical vapour deposition (CVD) and physical vapour deposition (PVD). In carbide CVD, the film is grown from gas-phase precursors of the reactants (typically a metal-chloride and a hydrocarbon gas), which often requires high temperatures (800-1400 oC) for the synthesis. In PVD, the film is grown from atoms in vapour-phase. The vapour can be produced by many techniques, but the most common are evaporation or sputtering. PVD typically utilises much lower temperatures than CVD, since the plasma yields energetic species and no precursor reactions need to be thermally activated. In this work, I have used two evaporation techniques, electron beam evaporation of metals and thermal heating of C60 in a Knudsen effusion cell. Furthermore, I have used DC magnetron sputtering, with and without the co-evaporation of C60. For all my studies, the PVD processes were performed in ultra-high vacuum (UHV) chambers, out-baked to a base-pressure in the range of 10-10 Torr (except publication IV,which was made in an high vacuum (HV) chamber, having a base-pressure of ~10-8

Torr).

3.1.1 Evaporation In publication I-IV, C60 was used as a virtually pure carbon precursor for the carbide film growth. The C60 molecule remains intact when evaporated from solid phase and it has a sufficient high vapour pressure at 400-500 oC45. A resistively heated Knudsen effusion cell was used to evaporate C60 to obtain a good process control over the C60 flux. The temperature was used to control the vapour pressure of C60 in the cell and thereby also the relative C60 flux that reached the substrate. For a more detailed description about the set-up the reader is referred to earlier publications from our group1,46.

Chapter 3. Thin films


In publication I, the atomic metal flux was provided by electron beam evaporation of the pure element in the vacuum chamber. The evaporation is performed by accelerating electrons over a large potential drop (typically 0.7-1.9 kV) towards the tip of a metal rod. This will induce heating and eventually melting of the tip and the metal evaporates. Ti and V were easy to evaporate i. e., did not require too high powers and were stable over long time. It was also relatively easy to manually control the individual metal fluxes by the applied power. In contrast, the attempts to evaporate Nb and Mo were connected with major problems. The strong bonding character of these elements gives extremely high melting points and low vapour pressures. This had of course consequences for the depositions, very high powers had to be applied to evaporate Nb and Mo. The equipment could not stand those high temperatures for any prolonged depositions and a shortcut was normally encountered after a few minutes. Therefore, only very thin films of (Nb,Mo)C could be made.

3.1.2 Sputtering DC magnetron sputtering of metals together with co-evaporated C60 was used in publication II-IV. In publication V-IX, C60 has been replaced with sputtering of carbon. In sputtering, ions are accelerated towards a target of the coating material and the collision ejects target atoms towards the substrate. The principle for DC magnetron sputtering47,48 is described in figure 3.1.

Figure 3.1. Schematic sketch of a DC magnetron sputter chamber illustrating the basic principle of sputtering. The applied voltage accelerates Ar+ ions that knock out atoms from the target, which are deposited on the substrate. Secondary electrons, trapped in the magnetic field from the magnets placed behind the target, ionises the Ar gas and the process is self-sustained as long as the voltage is applied.

Ar+

e-e- e-

e-

Magnets

Magnetron

Target

Ar-Plasma

Film

Substrate Heater

Sample-Holder

Target atoms

DC+

_



The sputtering gas, normally a noble gas, is introduced in the vacuum chamber and a high negative voltage is applied over the target. This accelerates the few naturally occurring ions in the gas towards the target. The impact of the ion with the target transfers the momentum of the ion to the target atoms. This has two results: First, if the momentum is high enough, target atoms can be ejected, sputtered away form the target. Second, some of the transferred energy from the ion will cause emission of secondary electrons from the target. The electrons are accelerated from the target in the applied potential and can ionise more gas atoms, which will sustain the process. This process can be highly optimised with strong magnets placed under the target. The magnetic field will, together with the applied electric field, trap the secondary electrons to move in a circle orbit above the target. This enhances the ionised sputter gas plasma close to the target and yields more effective sputtering.

3.2 Film growth In PVD, the atoms arrive from the gas phase to the substrate. The atoms adsorb on the surface. Depending on the energy available, the adatom can diffuse on the surface before it finds an energetically stable position, or it may also re-evaporate. See figure 3.2 for typical processes that may occur during film formation48.

(a)

(g)

(b)

(c)

(d)

(e)

(f)

(a)

(g)

(b)

(c)

(d)

(e)

(f)

Figure 3.2. Simplified, but typical atomistic processes that occurs during film growth. (a) The atom is deposited from the gas phase. Once on the surface, the adatom can diffuse (b), meet other adatoms and form a dimer or attach to an existing island (c). Diffusion can occur along the island edges (d) and adatoms on top of an island can form a new layer (e) or diffuse to the edge (f). Some atoms may desorb from the surface (g).



There are three different growth modes that are commonly used to describe the nucleation and following film growth48-50.

(i) Frank-van der Merwe (FM), 2D* layer-by-layer growth. The adatoms form a complete monolayer before growth is initiated on a second layer. This growth is observed when the binding between adatom and substrate is stronger than between the adatoms. It also requires high adatom surface mobility.

(ii) Volmer-Weber (VW), 3D* growth. Small clusters of the adatoms nucleate on the substrate surface and continue to grow into 3D islands. The islands can meet and grow together. This growth occurs when the adatoms from strong bonds and the growth kinetics do not allow enough surface mobility.

(iii) Stranski-Krastanov (SK), combined 2D and 3D growth. Initially a 2D layer-by-layer growth followed by re-nucleation to a 3D island growth**.

These growth models can be used to discuss and understand the resulting thin film microstructure, which can show large variations depending on the circumstances during deposition. Typically, growth-rate and substrate temperature have major influence on the observed growth mode.

3.2.1 Epitaxial growth The crystalline surface of the substrate can work as a template for epitaxial film growth. To illustrate the concept of epitaxy, I would like to use an example from publication II. Epitaxial means that the film has a crystallographic orientation relationship to the substrate. Figure 3.3 shows a transmission electron microscopy image with atomic resolution of a W film (with 7 at% C) grown on MgO(100) substrate. The epitaxial relationship between film and substrate is clearly visible in the interface.

Epitaxial growth is described by the out-of-plane relationship: film(hkl)//substrate(hkl) and in-plane directional relationship: film[uvw]//substrate[uvw], where (hkl) is a crystalline plane, [uvw] is a direction and “//” means that they are parallel48. Figure 3.4 shows the in-plane relationship between bcc W and MgO. This is a kind of “cube on cube” growth, where the W unit cell is rotated 45 degrees and has the diagonal aligned along the unit cell axis on the MgO substrate. This gives an epitaxial growth of W(001)//MgO(001) and the in-plane relationship W[110]//MgO[100].

* 2D and 3D means 2 and 3 dimensional, respectively ** Differences in interfacial surface tensions between substrate, film and vapour phase will be determining factors if the layer-by-layer growth is sustained or if renucleations occur, as described in standard text books48,50.



Figure 3.3. Epitaxial W(001) film grown on MgO(001) that quite nice illustrates the topic of this chapter. The mismatch between film and substrate induces misfit dislocations, encircled by white circles. The film shows a tetragonal distortion and the strain is relaxed through the insertion (and termination) of extra half-planes, encircled by black circles.

W

O

Mg

W(001)

MgO[100]

W[110]

MgO(001)W

O

Mg

W

O

Mg

W(001)

MgO[100]

W[110]

MgO(001)

Figure 3.4. It is possible to grow epitaxial W(001) films on a MgO(001) substrate. The cubic appearance of the oxide surface works as a good template for the cubic tungsten lattice. The W unit cell is rotated 45o compared to the MgO unit cell, which gives the in-plane directional relationship W[110]//MgO[100].

Normally, when the film and substrate is not of the same material (hetero-epitaxial growth) there will be a difference in atomic plane distances (d) at the film/substrate

MgO(001)

-W(001)

2 nm

MgO(001)

-W(001)

2 nm



interface48. This is the so-called mismatch ( ) between the film and substrate, which is calculated from the in-plane relationship:

Sub

SubFilm

ddd

The mismatch between W (dW(110)=2.238Å) and MgO (dMgO(100) = 2.1056Å) is 6.3%, which is quite large. It is often energetically favourable for the film to adjust its atomic lattice to match the substrate. If the film has a larger lattice than the substrate, it will experience a coercive uniform strain in the plane. This can lead to a tetragonal distortion as shown in figure 3.5.

Figure 3.5. A mismatch between film and substrate can lead to a strained film with tetragonal distortion of the film lattice when the film matches the substrate lattice at the interface. The strained unit cell of the film can be described by the in-plane axis (af or a//) and out of plane axis (cf or a ).

The film can also compensate for the mismatch by introducing misfit-dislocations in the film substrate interface to reduce the strain48. Such misfits can be seen in figure 3.3, encircled by white circles. However, even with misfit dislocations, diffraction showed that this film was strained with a tetragonal distortion. The film cannot maintain the strained growth, but will relax the compressed lattice during growth. This can also be seen in the TEM image in figure 3.3, where extra half-planes are generated during growth, encircled by black circles.

3.2.2 Pseudomorphic growth PVD processes often utilise low substrate temperatures far from thermodynamical equilibrium. This gives kinetic limitation for bulk diffusion, atomic re-arrangement and phase decomposition. However, the adsorbing species are often energetic enough (especially in sputtering) to allow high surface mobility and since the synthesis is a continuous growth process, it is therefore possible to obtain high crystalline quality in the film. These circumstances during growth leads to that films of metastable (or high-temperature) phases are commonly observed using PVD. When the film adopts a crystal structure and lattice constants that is not the

a0

a0

a//

a

af

cfaf

af

a0

a0a0

a0

a//

a

af

cfaf

af

a0

a0

Relaxed lattice Tetragonal distortion of the film lattice



equilibrium structure for the material, but which match coherently to the underlying substrate, it is referred to as pseudomorphic growth50,51. A metastable film structure is often stabilised by epitaxial growth or in multilayers49,52. The metastable growth can be persistent if the nucleation barrier is sufficiently high, otherwise a more stable structure will soon be adopted in the film49. The TEM image (from publication II) in the Swedish Summary on page 64 shows a typical example of Stranski-Krastanov (SK) growth in a tungsten carbide film. The first monolayers exhibit epitaxial pseudomorphic 2D growth (FM) imposed by the substrate. However, as the film grows thicker it becomes more favourable to form islands with following renucleations and 3D nanocrystalline growth (VW). It should be noted that it is possible to control the occurring growth mode. This will be further discussed in section 6.1.1.

16

Chapter 4.

Characterisation

The principal techniques I have used to gain structural and chemical information of the films are x-ray diffraction, transmission electron microscopy imaging, and x-ray photoelectron spectroscopy. Four-point probe and nanoindentation measurements have been performed to determine electrical and mechanical properties.

4.1 X-ray diffraction The theory of x-ray diffraction (XRD) is based on the diffraction of x-rays against crystalline planes in the radiated specimen. The angle of diffraction follows Bragg’s law:

n =2dsin

Where n is an integer, is the wavelength of the x-rays, d is the inter-planar spacing in the crystal and is half of the diffracting angle. By recording the diffraction angles from a sample with a known x-ray wavelength, the crystal structure can be identified from the plane distances and peak intensities. For a more detailed description of XRD theories, I refer the reader to a standard textbook in XRD53. In this chapter, I briefly present and discuss the different thin film diffraction techniques used in the publications. I would also like to present an intuitive interpretation of reciprocal space mapping (RSM) in order to allow non-specialists to better follow the arguments and discussion about RSM presented in the publications. For a deeper understanding of the theories and physics of RSM I recommend the thesis of J. Birch54 as the textbook on the topic.

4.1.1 Thin film diffraction Thin film diffraction is often performed with Bragg-Brentano geometry where both the x-ray source and the detector are aligned on a circle. This can be utilised with a focusing geometry, where the divergence is controlled by slits, or with a defocusing geometry with parallel-beam set-up with x-ray mirrors and/or parallel-plate



collimators. Standard measurements on films are -2 diffraction, which is also referred to as "locked couple" because the detector follows exactly the doubled incoming angle. This is a symmetric scan and diffraction will only occur from the crystal planes parallel to the sample surface. If there is a preferred orientation (due to texture or epitaxy) in the film, those peaks will contribute most to the diffractogram. A problem with this set-up is that the substrate peaks can dominate in intensity. This can be overcome by making a scan with a small offset ( ) on . Normally referred to as "unlocked couple" since the scan is performed as an -2 scan, where =( - ).The offset affects the diffraction conditions and reduces the peak intensity. Since the substrate is generally of better crystalline quality than the film, its intensity will be reduced more effectively than the film peak, see figure 4.1. For polycrystalline films, it is advantageous to perform gracing incidence (GI) scan, using a de-focusing geometry with a fixed low incoming angle and scan the detector over 2 . The low incoming angle gives high x-ray intensity from the film and also enables diffraction from planes that are not parallel with the sample surface.

2

F(002)

S(002)

LC

Inte

nsity

2

F(002)

S(002)

Inte

nsity

RC

F

18 21

Inte

nsity

RC S

18 212

UnLC

S

F

Inte

nsity

2

F(002)

S(002)

2

F(002)

S(002)

LC

Inte

nsity

2

F(002)

S(002)

LC

Inte

nsity

2

F(002)

S(002)

Inte

nsity

RC

F

18 21

Inte

nsity

RC

F

18 21

Inte

nsity

RC S

18 21

Inte

nsity

RC S

18 212

UnLC

S

F

Inte

nsity

2

UnLC

S

F

Inte

nsity

Figure 4.1. Principal sketch of how x-ray diffraction (XRD) reciprocal space mapping (RSM) of a symmetric peak can be related to other more common thin film XRD measurements. Locked couple (LC) is a standard -2 scan and unlocked couple (UnLC) is a scan with a small offset in . The -direction in the RSM corresponds to the rocking curve (RC) -scan. The map shows the intensity distribution of the peak and gives a 2D plot of the diffraction conditions.

The -scan, or rocking curve (RC), gives the degree of preferred orientation in the film. The 2 value is locked on a certain reflection and the incident x-ray angle is

Chapter 4. Characterisation


varied around . The full width at half maximum (FWHM) of the peak from the -scan can be used to quantify the crystalline quality when comparing different films. The epitaxial relationship between film and substrate can be determined by a -scan.The incoming angle, and the Bragg-angle, 2 are chosen for a plane that is not parallel to the sample surface and the sample is rotated around its normal, giving a

-scan. The -scan in figure 4.2 exemplifies a typical “cube-on-cube” growth in an epitaxial film where the (113) peaks in both film and substrate are detected at the same -angle.

4.1.2 Reciprocal space mapping (RSM) The reciprocal space map is a three dimensional plotting of a series of unlocked scans performed with different offsets. This gives a -2 map plotted with 2 on the x-axis and on the y-axis, see figure 4.1. The scan where = (2 )/2 = represents an ordinary -2 scan as can be seen in inset to the left in figure 4.1. Along the -direction the map can be viewed as a series of rocking curve -scans, as can be seen in the insets under the map. An unlocked scan has also been added to illustrate the influence of a small offset, which reduce the substrate peak and gives a better view of the film peak. A reciprocal space map should preferably be plotted in reciprocal lattice units (rlu) (or; Q = 2 /d [Å-1]). However, even if plotted in degree space, the map still contains useful information. The reciprocal space can be viewed as the inverse of real space. Therefore, a broadening in reciprocal space will correlate to a short distance in real space. This correlation gives structural information about the film from the shape (or appearance) of the map, as discussed in publication I andIII.

The RSM in figure 4.1 is of a symmetric peak, for example a (002) peak from an epitaxial MeC(001) film on MgO(001) substrate. In comparison, the RSM in figure 4.2 is a map of asymmetric peak, for example (113) peak from an epitaxial MeC(001) film on MgO(001) substrate. The most pronounced broadening that originates from the substrate peak is the wavelength broadening (WLB), going from left to right in the maps in figure 4.1,2. The WLB depends on the degree of monochromatization of the x-ray source in the measurement set-up. The Bragg’s law explains this; since the d-value for the diffraction plane is fixed, will variations in give a broadening in 2 . The -broadening have already been identified above with the RC, see figure 4.1. This is a lateral broadening and originates from short in-plane distances, for example from mosaicity or misfit-dislocations. It should be noted that the map of the asymmetric peak should be tilted, plotted with rlu, the WLB would point towards origin of the reciprocal space. There are also other kind of broadenings, for examples X-shaped, V-shaped or noses, these are discussed in more detail in publication I and III. If the film experiences a tetragonal distortion, it shows in the positioning of peak in the asymmetric RSM as a shift away from the WLB, as illustrated in figure 4.2. A peak from a relaxed film will coincidence with the WLB (assuming a cubic substrate), due to symmetry of Bragg’s law. This makes it possible to calculate the tetragonal distortion of the film as described by J. Birch et al.55.



Figure 4.2. A -scan of a asymmetric peak in film and substrate gives the epitaxial in-plane relationship. It is also possible to make a RSM of the asymmetric peak. The appearance of the asymmetric map in degree space will look tilted compared to the correct plotting with reciprocal lattice units. The broadening and shape of the peaks as well as the position of the peak gives information about the film as discussed in the text.

4.2 Transmission electron microscopy Transmission electron microscopy (TEM) is a very useful technique to gain deeper understanding into the structure of matter. TEM has been utilised in my studies to verify epitaxial relationships, tetragonal distortion, determine grain sizes, make phase characterisation, study nucleation processes, imaging of dislocations, to study deformation behaviour and answer all kinds of questions regarding microstructure. TEM can be resembled with an optical microscope, but since the electron have a shorter wavelength than light, the resolution is much better. The resolution obtained can be in the Å range, with the possibility of imaging atoms. TEM gives a physical image of the sample, but the obtained image is a projection of the three-dimensional sample. Besides imaging, TEM can also be used for electron diffraction (selected area electron diffraction, SAED) to gain crystal information about the phases in the film. The sample must be very thin in order for the electron beam to penetrate through it. For high-resolution (HR) images, this means less than 100 Å thick samples.

NaCl-structure(113)-plane

Strained and tetragonal distorted film2

F(113

S(113)Relaxed film

2

F(113)

S(113)

(o)0 360

-scan

F(113)

S(113)



It should be noted, that comparing the TEM images in the publications included in this thesis there are variations in quality of the reproduction of the MAX-phases. One problem with the MAX-phases studied with TEM is the build up of charges, which induces drift in the sample during imaging. This problem is overcome by another microscopy with high-angle annular dark field (HAADF) in 1.4Å probe scanning TEM mode imaging, which could be used for some samples. The A-layers can be identified in the HR images by looking for two second-neighbour M-atoms placed on top of each other, the A-layer lies between them, as sketched in figure 4.3.

Figure 4.3. The identification of the A-layers in high-resolution TEM images of the MAX-phase can be troublesome. One trick is to look for the stacking of the M-atoms, normally it is possible to find the twin boundary separated by the A-layer, as indicated in the sketch.

4.3 X-ray photoelectron spectroscopy X-ray photoelectron spectroscopy (XPS) is also known as ESCA, which means electron spectroscopy for chemical analysis*. The latter name says more about the usefulness of this analysis technique. The sample is irradiated with x-ray photons of a known energy (h ) and electrons leave the atom with a kinetic energy (Ekin) equal to the photon energy minus the binding energy (EB) of the electron and the work function**. The Ekin is measured and intensity of the electrons is plotted against the EB. Since each element has its own specific electron levels with well-known EB, the spectra can be used as a fingerprint to identify the elements in the sample. Furthermore, it is possible to detect chemical shifts in the EB, which can be related to the bonding character to the neighbouring atoms. The relative intensity of the peaks from each element can be used to determine the elemental composition. This requires sensitivity factors that describe how the intensity of the photoelectrons varies for each element. To be able to gain as correct data as possible I have used standard samples of TiC, WC, Ti3SiC2, Ti:Al and Ti2AlC with known composition to determine S for the compositional analysis. The standard samples were also used to establish if there was any preferential sputtering during depth profiling with Ar-sputtering.

* Kai M. Seigbahn was awarded with the Nobel Price in Physics in 1981 for his contribution to the development of high-resolution electron spectroscopy56.** Albert E. Einstein earned the Nobel Price in Physics in 1921 for his discovery of the law of the photoelectric effect56.

M

A

X

M

A

X



4.4 Four point probe measurements Four-point-probe was used to determine the resistivity of the films (see figure 4.4 for a sketch of the set-up). Current is passed through the outer electrodes through the film and the inner electrode measure the voltage drop. The sheet resistance (RS) can be used to calculate the bulk resistivity ( ) if the film thickness (t) is known, with the following equations57:

IURS 532.4 and tRS .

Figure 4.4. Principle sketch of four-point probe measurement.

For MAX-phase films on a TiC seed layer, the conducting TiC contributes to the measured sheet resistance. The films can be viewed as two parallel-coupled resistors. In publication VI, was the TiC seed layer omitted. The resistance in TiC is about one order of magnitude higher than for the MAX-phase and also less than 20 % of the total film thickness. This means that the TiC contribution to the total sheet resistance is more or less negligible. However, in publication VII and VIII,MAX phase films are also grown at lower substrate temperatures, which reduces the crystallinity and also the conductance. Therefore, the TiC contribution to the measured sheet resistivity has to be considered when the bulk resistivity is calculated. This can be done with following formula, deduced from parallel-coupled resistors:

STiCS

STiCSMAXS RR

RRR

)(

)()( and MAXMAXSMAX tR )(

Where, RS(TiC) is the measured sheet resistance for a TiCx film deposited with same parameters as the seed-layer.



4.5 Nanoindentation Hardness and Young's modulus have been measured with nanoindentation for some of the films within this thesis. Two instruments have been used. In publication IIwe used a NanoIndentor II instrument with a three-sided diamond Berkovich tip. For the MAX-phase films in publication V-VIII a Triboscope , Hysitron Inc. was used with both Berkovich tip, and with a sharper cube-corner tip. The properties were calculated from the load and displacement curve of the indent using the Oliver and Pharr method58. The use of very small loads (0.1-3.0 mN for MAX-phases and 0.2-20 mN in W-C system) gives rise to certain problems. Great care has to be taken, not to measure the properties of the substrate and the indent should normally not be deeper than 1/10:th of the film thickness to give accurate data. Furthermore, surface roughness, and surface oxidation can obstruct the measurement. In this work, all measured films have been at least 0.5 m thick and a series of indents have been made for each load to provide as reliable results as possible.

4.6 Soft x-ray spectroscopy Publication IX investigates and compares the electronic structure and inter-atomic bonding relationship of Ti3AlC2, Ti3SiC2, Ti3GeC2 and TiC using soft x-ray absorption (SXA) and emission (SXE) spectroscopies. This is a useful measurement technique to investigate the valence-band structure especially since it is a bulk sensitive technique (in contrast to photoelectron spectroscopy, which is very surface sensitive). In soft x-ray spectroscopy the sample is irradiated with synchrotron radiation of tuneable energy and the x-ray emission or absorption is measured. The x-ray emission and absorption transitions involve both valence and core levels in the atom, see figure 4.5. Therefore, it is possible to probe each element separately by tuning the excitation energy to the appropriate core edge. The SXA spectrum gives the transition energies and core edges for the elements in the probed sample. These energies are then used for the SXE measurement to probe each state of the elements. In SXE the energy of the emitted x-ray is measured, it gives the difference between the binding energy of the valence band electron and the core level. The SXE spectra can simply be interpreted in terms of partial density of states (DOS) (and thus be compared with the calculated spectral function by projection of the DOS on the transition probabilities59, as described in publication IX).



Figure 4.5. A schematic view of how soft x-ray absorption (SXA) and emission (SXE) spectroscopy measurements correlates to core level and valence band energies. The x-ray photoelectron (XPS) ejection is also indicated. The inset shows how each spectrum from SXE, SXA and XPS relate to each other. (The figure is reprinted with permission from T. Wiell60.)

24

Chapter 5.

Theoretical calculations

The main part of the work presented in this thesis is experimental studies. However, a deeper understanding of matter can be achieved by a theoretical treatment. Abinitio calculations can predict phase stabilities and trends in the periodic table. It is also possible to discuss bonding mechanisms and study the density of states (DOS) from the calculations. Furthermore, bulk modulus and Young’s modulus, high-pressure behaviour and other physical properties can be estimated61. An ab initio (or first principle) calculation means that the theory does not include any input from experimental data, besides the structure. The theory is based on the interaction of charges and the electron distribution that gives the charge (electron) density. This approach is known as the Density Functional Theory (DFT)62,63, which awarded Walter Kohn the Noble prize in Chemistry 199856. The basic theories have been adapted and implemented in various methods distinguished by how the electron density is described and on the approximations made.

In publication VI we study the MAX-phases in the Ti-Si-C system using the Vienna Ab-inito Simulation Package (VASP)64. The VASP calculations were based on the local density approximation (LDA)65 and uses plane-wave pseudo-potentials to describe the periodic electron density within the crystal. This method is known to be a fast computational method that gives accurate and reliable results. Therefore, it is also very popular and commonly used. Higher accuracy can be achieved by the use of more complex methods to describe the electron density3. One of the most accurate computational methods available is the Full-Potential Linear Muffin-tin Orbital (FP-LMTO) method. However, calculations using this approach take quite long time and cannot be used on larger systems. Normally, the results from the VASP calculations are validated by re-calculation of some data points with the FP-LMTO method. The reader is referred to publication VI or the licentiate thesis of S. Li66 or the thesis of H. W. Hugosson3 and references therein, for more details concerning the calculation methods and the theories.

An important output of a DFT calculation is the total energy of the phase or structure. The cohesive energy (Ecoh) of the phase/structure can be obtained by subtracting the energy of the elements from the calculated total energy61. The energy difference ( E) between two (or more) phases or structures can be used to predict the relative stability. This was performed in publication VI, where Ecoh for suggested MAX phases was compared with competing equilibrium phases at



corresponding composition. A negative E indicates a stable phase. However, such conclusions have to be taken with a grain of salt. The calculations only give the change in enthalpy ( H) for such a reaction at 0K. To really determine the stability at higher temperatures one would have to consider the sign of G for the assumed decomposition reaction. This requires knowledge of the temperature dependence on the enthalpy values and, in addition, entropy effects would have to be included. But still, the energy difference gives valuable information, especially when studying trends in the periodic table or comparing competing reactions.

26

Chapter 6.

Results and discussions

I have divided the presentation of the results into several sub-chapters. First, depositions of carbides, followed by the design of epitaxial MAX-phase films, thereafter, compounds with engineering potentials for controlled mechanical properties by tuning of the composition, and finally, some results of artificial thin film structures. I will also put the results in a larger context and discuss engineering possibilities and technological applications of these thin films.

6.1 Deposition of transition metal carbides The attractive properties of high melting points, high hardness, considerably low friction, and good oxidation resistance have made the carbides useful as wear protective coatings. But since the carbides also have high electrical and thermal conductivities, they have potential for a variety of other applications in electronic devices such as electrically conducting diffusion barriers, contacts and diodes, etc13.However, the use of carbide films is generally limited by the requirement of high process temperatures, which limits the possible substrate materials both in electronic and wear protection applications. The development of low-temperature carbide PVD processes for high quality crystalline films has therefore been one of the aims with the present thesis. Using a low temperature opens up, both possibilities for combination of carbides that under thermal equilibrium condition would separate or decompose, and it facilitate the use of substrates that normally not can be used in carbide deposition processes.

6.1.1 Tungsten carbides Thin film growth in the W-C system is connected with many unsolved questions. For example, regardless of deposition technique (CVD or PVD), the use of hydrocarbons as carbon precursor would most certainly yield phase mixtures of W, W2C, -WC1-x and WC or nanocrystalline or nearly amorphous films of -WC1-x67.The hexagonal stoichiometric WC phase, cannot be deposited without very special conditions68,69. It has been discussed that incorporation of hydrogen or hydrocarbons



during growth de-stabilises the WC-phase. Furthermore, it has also been proposed that the commonly observed -WC1-x phase is stabilised by a small grain size since large-grained microstructures have not been reported. A low-temperature process without hydrocarbons for a controlled growth in the W-C system is therefore most interesting.

The phase formation in the W-C system was examined in publication II over a wide range of compositions using sputtering of W together with co-evaporated C60 as carbon source. Thin films were deposited on single crystal substrates to promote a controlled growth and the temperature was varied from room temperature (RT) to 800 oC. The observed phase formations are shown in figure 6.1, which is an experimental phase diagram derived from the study.

-W -W2C -WC1-x

oC

25

800 WC

20 35 50at% C

-W -W2C -WC1-x

oC

25

800 WC

20 35 50at% C

Figure 6.1. Experimental phase diagram from publication II. Thin films were deposited by sputtering of W with evaporated C60 as carbon source over a wide C60/W flux ratio and from room temperature up to 800 oC. All phases in the W-C phase diagram (Fig. 2.1) could be grown, although the hexagonal WC phase was only observed in phase mixtures, as indicated by the hatched lines.

The hexagonal WC phase was only observed in phase mixtures at high temperatures (800 oC) and could neither be synthesised as single-phase films nor epitaxially grown. This result indicates that, even without hydrocarbon precursors, growth of WC phase is connected with problems. Furthermore, re-sputtering of C from the films during growth by back-scattered Ar-neutrals was observed. It was suggested that the formation of WC phase was hindered by the fact that it exhibits no homogeneity range and thus it is likely to be more sensitive for the growth conditions than the other W-C phases. This conclusion is in agreement with the observation by A. H. Cottrell that the WC phase is very sensitive for vacancy formation, which causes the formation of a second phase, such as W2C12.

However, metastable film growth of single phase W2C and -WC1-x could be established on various substrates at very low temperatures. Epitaxial films of the high-temperature phase W2C was successfully grown on MgO(111) at 400 oC even though W2C only is stable above 1250 oC15. On other substrates, poly-crystalline films were deposited. The -WC1-x phase was found to prefer a nanocrystalline growth, see diffractogram (a) in figure 6.2. The grain size could be controlled by the

Chapter 6. Results and discussion


substrate temperature from 15-20Å at 100 oC up to 80Å at 800 oC. Furthermore, composite films of very fine grained -WC1-x (nearly x-ray amorphous) in an amorphous carbon matrix could be deposited using high C60/W flux ratios that gave more than 40-45 at% C in the films, diffractogram (b) in figure 6.2. The -WC1-xphase is also a high-temperature phase only stable above 2535 oC15, but the general occurrence of this phase in low-temperature PVD processes indicates that a fine-grained structure can act stabilising on this phase.

d=2.4Å

Cou

nts

(arb

. uni

ts li

near

) (111)

(200) (220) (113)

a)

b)

(222)

d=2.4Å

Cou

nts

(arb

. uni

ts li

near

) (111)

(200) (220) (113)

a)

b)

(222)

d=2.4Å

Cou

nts

(arb

. uni

ts li

near

) (111)

(200) (220) (113)

a)

b)

(222)

Figure 6.2. GI-XRD of -WC1-x films deposited at 400 oC with different carbon content. (a) Nano-crystalline microstructure formation with the NaCl structure in films with 40-45at% C. (b) Films with more than 45 at% C is more or less x-ray amorphous with only one weak and broad peak at 37.5 o (d=2.4Å).

Even though the -WC1-x prefers a random nanocrystalline growth, two different approaches were successfully used to stabilise epitaxial deposition on both MgO(100) and MgO(111) substrates at 400 oC (publication III). First, a very low deposition rate of only 6 Å/min was used to control the growth kinetics. This gave the adsorbed atoms more time for surface diffusion and a pseudomorphic epitaxial 2D FM-growth could be established. At higher deposition rates, the epitaxial growth was interrupted after a few monolayers followed by re-nucleation and random nanocrystalline growth (SK-growth mode), as can be seen in the figure on page 64 in the Swedish Summary. However, by using a few percent of Ti as additive, epitaxial growth could be established at a deposition rate of 40 Å/min. Calculations show that Ti stabilises the cubic -WC1-x (B1) phase70. This could explain why the re-nucleation is suppressed and a stable epitaxial growth is achieved. The results in publication III shows for the first time that large-grained -WC1-x can be synthesised under proper conditions. It is not only a small grain size that stabilises the -WC1-x phase, the growth kinetics are also a determining factor for the resulting microstructure.

6.1.2 Ternary carbides As mentioned in chapter 1, epitaxial growth of binary carbides, such as TiC and VC films were well established with C60 as carbon source. The similarity among the



transition metal carbides enables solid solutions and the prospect of combining two, or more, carbides can be utilised to design new materials with the possibility to tune in the desired properties. Publication I presents the results from a study of ternary carbide growth using two model systems: First the TiC-VC system, which share the same cubic NaCl-crystal structure and shows complete solid solubility, as shown in the ternary phase diagram in figure 6.318. NbC-MoC was chosen as the second model system since they represent carbides of different crystal structures without complete solubility. Both carbides have the cubic NaCl-structure, but the MoC is a high-temperature phase, stable only above 1800 oC.

Figure 6.3. Ternary phase diagram of Ti-V-C at 1000 oC, after T.F. Fedorov et al18.

Epitaxial Ti1-xVxCy(001) films was deposited within the whole range from 0



6.1.3 Epitaxial TiC growth by sputtering from Ti and C targets The role of C60 in epitaxial carbide growth has been discussed and questioned and the mechanism for epitaxial growth has not been completely investigated or fully understood1,2,72. It should be noted, however, that in literature, no papers report epitaxial carbide growth by other techniques at such low temperatures that was observed with C60. Therefore, it was assumed that the very low temperatures could be assigned to the presence of fullerenes during growth. To learn more about the influence of the fullerenes, C60 was replaced as carbon source with sputtering from a C-target. To our surprise, epitaxial TiC growth was obtained in the first deposition experiment and thus the special behaviour of C60 had to be re-evaluated. The absence in literature of low temperature epitaxial carbide growth can be explained by several factors. One reason can be that most CVD or reactive PVD carbide processes uses hydrocarbon gases as carbon precursor, which requires high temperatures to break the C-H bond. Hydrocarbon precursors can also favour fine-grained microstructures since the surface of the growing film will be covered by adsorbed hydrocarbons with limited surface mobility, which can act as re-nucleation points. Furthermore, hydrogen incorporation in the film can also affect the microstructure.

Epitaxial TiC growth by DC magnetron sputtering from two targets of Ti and C has been established on MgO(100), MgO(111), 6H- and 4H-SiC(0001) and Al2O3(0001). Good crystalline and epitaxial films have been grown at substrate temperatures from 1100 oC down to 100 oC. Without assisted heating, polycrystalline TiC films are formed. The influence of the vacuum-quality in the chamber has also been examined for TiC film growth. The purity and film quality decreases, but epitaxial growth is possible up to at least a base pressure of 2x10-5

Torr (in an unbaked UHV chamber). XPS analysis shows that epitaxial and single phase TiCx films with 0.5



substrates but on Al2O3 it shows twinning and domain formation due to different initial nucleation of ABC and BCA stacking of the Ti-planes. This growth behaviour was also confirmed with TEM as seen in figure 6.4 from publication VII.

TiC ABC

TiC BCA

Twin-type Moiré

Twin-type Moiré

Al2O3(0001) 3 nm

TiC ABC

TiC BCA

Twin-type Moiré

Twin-type Moiré

Al2O3(0001)

TiC ABC

TiC BCA

Twin-type Moiré

Twin-type Moiré

Al2O3(0001) 3 nm

Figure 6.4. High Resolution TEM image of twin formation in epitaxial TiC(111) film on Al2O3(0001) substrate. Above the scale bar, there is a terrace step on the substrate, which may lead to different stacking, ABC or BCA, during initial nucleation.

6.2 MAX-phase thin film design When we first heard and read34 about the MAX-phases we immediately wanted to study thin film growth of MAX phases as we recognize a material with unlimited possibilities for materials design (for reasons discussed in chapter 2.2). However, thin film deposition of MAX-phases had not been performed earlier, beside a few reports from the early 70’s of CVD on Ti3SiC273. Early attempts to deposit Ti3SiC2films by sputtering from compound target74,IV or using laser ablation75 resulted in poorly crystalline films and phase mixtures of Ti3SiC2, TiC, TiSix, SiC and other phases. This was explained by nucleation problem on the substrates and phase separation during growth. To overcome these problems, an approach was developed for single crystal and epitaxial MAX-phase film growth. Together with ABB Ltd., this method has been patented (publication XVII) as an international application and published under the patent cooperation treaty (PCT).

The method, which is simple and very straightforward, is developed from our method for epitaxial carbide film growth72 and includes a seed layer of the MX matrix elements that promotes the MAX-phase nucleation. With an epitaxial MX



seed layer a epitaxial stabilised MAX-phase film can be deposited on, for example, oxide substrates, as seen in figure 6.5. This method has been demonstrated using fourXVI successful approaches to deposit single crystal and epitaxial MAX-phase films: (i) sputtering of Ti and Si with C60 as carbon sourceIV, (ii) sputtering from a Ti3SiC2 compound targetIV, (iii) sputtering from three individual targets of Ti and C together with SiVI,VII, AlVIII or Ge76, and (iv) reactive sputtering from an alloy 2Ti:Al target in an mixed Ar/N2 discharge77,78. Figure 6.5 shows a single crystal Tin+1SiCnMAX-phase film nucleated on a TiC seed layer deposited with method (iii).

Figure 6.5. Cross sectional TEM image of a single crystal and epitaxial Tin+1SiCn(0001) film deposited at 900 oC from three elemental targets. The patented method to use a TiC seed layer to promote the MAX-phase nucleation works excellently.

6.2.1 Ti3SiC2 thin film growth The process for epitaxial Ti3SiC2 film growth was first presented in publication IVand was developed, and further studied in publication VII. Best results are obtained on a (thin) TiC(111) seed layer to promote the Ti3SiC2(0001) nucleation and growth (see for example diffractogram (a) in figure 6.9, which shows a single phase epitaxial 312 film on TiC seed layer on Al2O3(0001) substrate). Epitaxial Ti3SiC2films can also be grown without TiC seed layer on oxide substrates, but with slightly lower quality. Interestingly, the Ti3SiC2 nucleation is delayed after that the Si-flux is added to the Ti- and C-flux. A thin Si-deficient incubation layer of epitaxial TiC initiates the MAX film growth, both on the TiC seed layer and on a pure substrate, as can be seen in figure 6.6. Furthermore, the studies indicate that there seems to be a lower temperature limit around 700 oC for epitaxial Ti3SiC2 film growth. Below

100 nm MgO(111)

TiC(111)

MAX(0001)

100 nm MgO(111)

TiC(111)

MAX(0001)



700 oC down to 500 oC, an epitaxial TiC(111) growth is observed even with relatively high amounts of Si incorporated in the TiC film.

Figure 6.6. High resolution TEM image that shows the delayed nucleation of the Ti3SiC2(0001) film on the TiC(111) seed layer. The nucleation is initiated on different levels in the film. The arrow indicates a starting (or termination) point for the lateral growth of the A-layers in the MX-matrix to form the MAX compound.

Films deposited on a TiC(100) seed layer show that the Ti3SiC2 nucleates epitaxially with the (1015) orientation in four different directions on the four-fold-symmetric TiC(100) surface. But the observed preferred growth direction is still along the c-axis and the growth of the basal planes of the MAX-phase occurs on the TiC{111}planes. This leads to overgrowth of smaller grains with voids in the film and occasionally the TiC seed layer is consumed by the growth of Ti3SiC2 grains into the seed layer, see figure 6.7 from publication VII.

Figure 6.7. Cross sectional TEM image of an epitaxial but columnar Ti3SiC2(1015) film nucleated on a TiC(001) seed layer on MgO(001) substrate.

TiC(111)

Ti3SiC2(0001)

TiC(111)

Ti3SiC2(0001)

Basal plane

MgO(001)Void

TiC(001)

Ti3SiC2(10-15)

200 nm

Si onset

Basa

l plan

eBasal plane

MgO(001)Void

TiC(001)

Ti3SiC2(10-15)

200 nm

Si onset

Basa

l plan

e



The observations made in the TEM studies of the MAX-phase films suggest a complex nucleation process. The observed Si-deficient incubation layer indicates the presence of Si-segregation during growth and that there need to be a critical amount of Si present on the surface to initiate the Si-layer formation. However, the temperature study also shows that thermally activated Si-diffusion in the TiC is a requirement for (large-grained) Ti3SiC2 film growth. This is also supported by the Ti3SiC2(1015) film that penetrates into the TiC(001) seed layer along the closed packed TiC{111} planes. It is evidently that Si-diffusion occurs along the basal planes during growth. Figure 6.8 shows a high-resolution atomic force microscopy (AFM) image of the surface of a Ti3SiC2 film. Flat terraces and surface steps of one half and one full unit cell height are present*. The topographical appearance suggests a lateral growth mode by propagation of half or full unit cells. This is consistent with the TEM observations of the nucleation process.

Figure 6.8. Atomic force microscopy (AFM) image of the surface of a Ti3SiC2(0001) film grown at 900 oC. The step heights on the terraced surface correspond to half and full unit cells.

6.2.2 New MAX-phases in the Ti-Si-C system Publication VI is dedicated to new MAX compounds in the Ti-Si-C system. Earlier high temperature bulk studies, only revealed the Ti3SiC2 phase79. Furthermore, Ti3SiC2 is the only reported stable MAX-phase with Si as A-element34. Considering the variety of other Mn+1AXn phases, there is an underlying question here about whyTi3SiC2 is such a lonely fellow in a large family of closely related compounds. It could be possible to stabilise metastable Tin+1SiCn compounds other than 312 by taking advantage of our patented low-temperature process. Considering that the epitaxial growth of the MAX phase film occurs in the basal plane direction, the film

* It would be interesting to know if the top-most layer of the terraces is terminated by Si-layers or TiC?



cannot really know how many TiC layers it grows on top of the latest Si-layer before the next Si layer forms*. Therefore, a series of deposition were made with ~50 at% Ti and Si and C concentration to match different Tin+1SiCn compositions. Figure 6.9 from publication VI summaries the XRD results from this study and Table 6:I shows the observed cell parameters.

Figure 6.9. -2 diffractograms of epitaxial Tin+1SiCn(000l) MAX-phase thin films grown on TiC(111) seed layer. (a) Ti3SiC2 single-phase film. (b) The new MAX-phase, Ti7Si2C5 grown in a multilayer film of 312/725/413. (c) The new MAX-phase, Ti5Si2C3 in a phase mixture film of 312, 523 and Ti5Si3Cx (marked with an arrow). (d) Ti4SiC3 single-phase film. The films in (a,b) and (c,d) were deposited on MgO(111) and Al2O3(0001) substrates, respectively.

Table 6:I. Unit cell parameters of the observed Tin+1SiCn MAX-phases, determined with XRD and RSM on epitaxial films grown on TiC(111) seed layer. The c-axis for 523 and 725 is given assuming a P-centred hexagonal unit cell, which requires three repetition of each 211/312 or 312/413 sequence to complete the c-axis, see discussion in publication VI and appendix 4. (The observed variations in c-axis in different films for each phase were in the order of 0.05 Å.)

a-axis c-axisn Phase (Å) (Å)1 Ti2SiC2 Ti3

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