Linköping Studies in Science and Technology

Licentiate Thesis No. 1270

Growth and Characterization of Ti-Si-N Hard Coatings

Axel Flink

LiU-TEK-LIC-2006:51

Thin Film Physics Division

Department of Physics, Chemistry, and Biology (IFM)

Linköping University, 581 83 Linköping, Sweden

ISBN: 91-85643-85-8 ISSN: 0280-7971

Printed by LiU-Tryck, Linköping, Sweden, 2006

Abstract

Metastable (Ti,Si)N alloy and TiN/SiNx multilayer thin solid films as well as SiNx/TiN surfaces have been explored. Cubic Ti1-xSixN (0�x�0.14) films deposited onto cemented carbide (WC-Co) substrates by arc evaporation exhibited a competitive columnar growth mode where the structure transforms to a feather-like nanostructure with increasing Si content as revealed by x-ray diffraction and transmission electron microscopy. X-ray photoelectron spectroscopy revealed the presence of Ti-N and Si-N bonding, but no amorphous Si3N4. Band structure calculations showed that phase separation of NaCl-structure Ti1-xSixN solid solution into cubic SiN and TiN phases is energetically favorable. The metastable microstructure, however, was maintained for the Ti0.86Si0.14N film annealed at 900 °C, while recrystallization in the cubic state took place at 1100 °C annealing during 2h. The Si content influenced the film hardness close to linearly, by combination of solid-solution hardening in the cubic state and defect hardening. For x=0 and x=0.14, nanoindentation gave a hardness of 29.9±3.4 GPa and 44.7±1.9 GPa, respectively. The hardness was retained during annealing at 900 °C. Nanostructured materials, e.g., nanocomposites and nanolaminates, are defined by internal interfaces, of which the nature is still under debate. In this work two-phase model systems were explored by depositing SiNx/TiN nanolaminate films, including superlattices containing cubic SiNx, by dual target reactive magnetron sputtering. It is demonstrated that the interfacial phase of SiNx onto TiN(001) and TiN(111) can be crystalline, and even epitaxial with complex surface reconstructions. Using in situ

structural analyses combined with ab initio calculations, it is found that SiNx layers grow

epitaxially, giving rise to strong interfacial bonding, on both TiN(001) and TiN(111) surfaces. In addition, TiN overlayers grow epitaxially on SiNx/TiN(001) bilayers in nanolaminate structures. These results provide insight into the development of design rules for novel nanostructured materials.

Preface This Licentiate Thesis is based on my research carried out with the Thin Film Physics Division at Linköping University in collaboration with SECO Tools AB in Fagersta, the Materials Science Department at University of Illinois at Urbana-Champaign, and the Department of Materials Chemistry at Uppsala University. The work is supported by SECO Tools AB, the Swedish Research Council (VR), and the Swedish Foundation for Strategic Research (SSF). Included Papers Paper I Influence of Si on the Microstructure of Arc Evaporated (Ti,Si)N Thin

Films; Evidence for Cubic Solid Solutions and their Thermal Stability A. Flink, T. Larsson, J. Sjölén, L. Karlsson, L. Hultman, Surf. Coat. Technol. 200 (2005) 1535-1542

Paper II Toward Understanding Interface Structure in Superhard TiN-SiN

Nanolaminates and Nanocomposites L. Hultman, J. Bareno, A. Flink, H. Söderberg, K. Larsson, V. Petrova,

M. Odén, J. E. Greene, I. Petrov, Submitted for publication

Other Papers by the Author Paper III Deposition of Ti2AlN Thin Films by Reactive Magnetron Sputtering T. Joelsson, A. Flink, J. Birch, L. Hultman, Manuscript in final preparation Paper IV MAX-Phase Ti2AlN Coatings by Arc Deposition

A.Flink, J. Sjölén, L. Karlsson, L. Hultman, In manuscript

Paper V Growth and characterization of single crystalline TiN/SiNx superlattice

films H. Söderberg, A. Flink, J. Birch, L. Hultman, M. Odén, In manuscript Paper VI Role of Carbon in Boron Suboxide Thin Films

D. Music, V. M. Kugler, Zs. Czigany, A. Flink, O. Werner, J. M. Schneider, L. Hultman, U. Helmersson, J. Vac. Sci. Technol. A21 (2003) 1355

Acknowledgements Many persons have contributed to my work and I would especially like to thank: Lars Hultman, my supervisor. Thank you for giving me the opportunity to do both my diploma work and PhD in the Thin Film Group. I am very grateful for all support you have given me so far. Lennart Karlsson, Jacob Sjölén, and Tommy Larsson at SECO Tools AB for fruitful collaboration. I always enjoy going to Fagersta and SECO for meetings and work. Javier Bareno, Vania Petrova, and Ivan Petrov from University of Illinois at Urbana-Champaign for taking care of me during my visits in Chambana. Hans Söderberg and Magnus Odén from Luleå University of Technology. Hasse, for nice collaboration and discussions. Magnus, for your efforts to always answer my endless list of questions regarding nanoindentation. Karin Larsson for introducing me and Javier to CASTEP. Per Persson, for helping me learning TEM, both by practice and discussions. Jens Birch for giving fast and accurate answers, and for increasing the already good spirit in the group. Hans Högberg, for always taking time to answer my questions, and for encouragement. Karl-Olof Brolin, Inger Eriksson, and Thomas Lingefelt for your endless helping spirit. Anders Hörling, for sharing your thoughts, both about science and life in general. Per Eklund, my old neighbor in Ryd, for our interesting football discussions, and for recommending me to apply for a diploma work at the Thin Film Physics Division. Anders E, Johan, and Timo for our unforgettable golf/poker trips. Erik, Fredrik, Martina, and Naureen for a fantastic week on Iceland! All friends and colleagues, both past and present, in the Thin Film and Plasma groups. I really have a lot of fun both at work and beside work with you! My family, of course, for all your support throughout these years (and the years before I started in graduate school).

Table of Contents

1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.1 Hard Coatings for Cutting Tools . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

2 The Ti-Si-N System. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

2.1 Phase Diagram. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 2.2 Titanium Nitride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 2.3 Silicon Nitride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 2.4 TiN/SiNx Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 2.5 TiN/SiNx Multilayers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 2.6 Ternary Solid Solutions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

3 Thin Film Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

3.1 Physical Vapor Deposition. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 3.2 Arc Evaporation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9 3.3 DC Magnetron Sputtering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11 3.4 Growth of Metastable Solid Solution Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 3.4.1 Low-temperature Synthesis. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13 3.4.2 Ion-induced Recoil Implantation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

4 Theoretical Modeling . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

4.1 Density Functional Theory. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15 4.1.1 Approximations for Many-body Interactions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16 4.1.2 Pseudo Potentials and Plane Waves . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16 4.1.3 Linear Muffin-Tin Orbital. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16

5 Thin Film Characterization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

5.1 X-ray Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19 5.2 Electron Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20 5.3 Nanoindentation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 5.4 Scanning Tunneling Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22 5.5 X-ray Photoelectron Spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

6 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

6.1 Ti1-xSixN Alloy Films. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25 6.2 TiN/SiNx Nanolaminate Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

Paper I . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33

Paper II . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45

2

1 Introduction

Ceramics is an interesting class of material in the sense of heat resistance, high hardness, thermal shock resistance, and producabililty as thin films. Ceramics can be defined as …“solid compounds that are formed by the application of heat, and sometimes heat and pressure, comprising at least two elements provided one of them is a non-metal or a non-metallic solid. The other element(s) may be a metal(s) or another non-metallic solid(s).”1 The properties of the ceramics make them attractive for thin film applications within the metal cutting industry. Many industries are dependent of tools for metal cutting applications; this has created a strong interest for developing new methods and materials for making the tools more efficient and cheaper.

1.1 Hard Coatings for Cutting Tools

During the 1970’s replaceable cutting inserts together with hard coatings were introduced. One of the first wear-resistant coating materials in the modern cutting tool industry used was TiN.2 It exhibits high hardness and stiffness which makes it suitable as a cutting tool coating. The main shortcoming is, however, limited stability due to oxidation at temperatures above 500 °C.3 Today the work temperatures of tools are typically between 800-1200 °C. Therefore, there is a growing interest concerning coating materials with improved thermal stability. The present design concept has been to employ ternary compounds, e.g. (Ti,Al)N4,5. (Ti,Al)N offers improved oxidation resistance, better thermal stability, defect (compressive residual stress) hardening, as well as newly discovered age hardening6. Together, these factors improve the life time of the tool, and provide the possibility to work at higher cutting speed. (Ti,Al)N has been used as a protective coating for cutting tools since the early 1990’s and is still a work horse for hard coatings.7 Today, (Ti,Al)N used in metal cutting industry are examples of metastable hard coatings that can be synthesized by arc evaporation at low temperature. Further on, the coating technology development and research on ternary compounds have expanded to cover a range of compounds based on the Ti-Al-N, Cr-Al-N, and most recently, Ti-Si-N systems (see Paper I). One example from the latter are the TiN/SiNx nanocomposites8, which exhibit thermal stability and very promising mechanical properties including superhardness. These nanocomposites consist of TiN nanocrystallites embedded in what is assumed to be amorphous SiNx. However, there is still the question about the nature of this tissue phase9 (see Paper II). Also, superhard

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TiN/SiNx multilayers, or nanolaminates, have been synthesized.10,11 Furthermore, recent work points to the possibility of fabricating (Ti,Si)N solid solutions12,13,14 (see Paper I) as well as epitaxially stabilized cubic-SiNx (see Paper II). The objective of this thesis is to explore the synthesis, structure, and properties of the novel (Ti,Si)N alloy and SiNx/TiN nanolaminate thin films. 1 M. W. Barsoum, Fundamentals of Ceramics, McGraw-Hill (1997) 2 P. O. Snell, Jernkontorets Anm. 154 (1970) 413 3 H. Ichimura, A. Kawana, J. Mat. Res. 8 5 (1993) 1093 4 O. Knotek, W. Bosch, T. Leyendecker, Proc. 7th Int. Conf. Vacuum Metallurgy, Linz, Austria 1985 5 W. –D. Münz, J. Göbel, Proc. 7th Int Conf. Vacuum Metallurgy, Linz, Austria 1985 6 A. Hörling, PhD Thesis (Linköping Studies in Science and Technology, dissertation no. 922, Linköping University, Sweden 2005 7 S. PalDey, S. C. Deevi, Mat. Sci. And Eng. A342 (2003) 58 8 S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64 9 S. Hao, B. Delley, C. Stampfl, Phys. Rev. B 74 (2006) 035402 10 H. Söderberg, M. Odén, J. M. Molina-Aldereguia, L. Hultman, J. Appl. Phys. 97 (2005) 114327 11 X. Hu, H. Zhang, J. Dai, G. Li, M. Gu, J. Vac. Sci. Technol. A23 (2005) 114 12 F. Vaz, L. Rebouta, B. Almeida, P. Goudeau, J. Pacaud, J. P. Riviere, J. Bessa Sousa, Surf. Coat. Technol. 120-121 (1999) 166 13 J. L. He, C. K. Chen, M. H. Hon, Mater. Chem. Phys. 44 (1996) 9 14 G. Pezzotti, I. Tanaka, Y. Ikuhara, M. Sakai, T. Nishida, Scr. Metall. Mater. 31 (1994) 403

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2 The Ti-Si-N System

In Paper I cubic Ti1-xSixN metastable solid solution coatings deposited by arc evaporation were studied. In Paper II TiN/SiNx multilayers were deposited by reactive DC sputtering with different thicknesses of the SiNx layers.

2.1 Phase Diagram

Fig. 2.1 The ternary phase diagram for Ti-Si-N at 1000 °C.1

The phase diagram at 1000 °C in Fig. 2.1 shows that Si3N4 is the only stable Si-N compound. TiN is stable over a wide stoichiometry range. Furthermore, there is no ternary phase present.

2.2 Titanium Nitride

TiN is a ceramic which is used in a wide field of thin film applications, from diffusion barriers to wear-resistant coatings to decorative coatings. TiN has a rocksalt structure (NaCl) with a unit cell consisting of 8 atoms; 4 Ti and 4 N. The lattice parameter is 4.24 Å.2 TiN exhibits high hardness, 20 GPa3 as single crystal thin film, and 26 GPa4 as polycrystalline, both grown by reactive DC magnetron sputtering.

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Fig. 2.2 Image illustrating the rocksalt, or NaCl, structure.

2.3 Silicon Nitride

Silicon nitride as a thin film material is mostly used within electronics. It exists as Si3N4 in three different polytypes, two hexagonal, �- and �, and one amorphous, a.5 There is also a high pressure, high temperature cubic phase.6 In a-Si3N4 the average binding distance is 1.74 Å.7

2.4 TiN/SiNx Nanocomposites

A nanocomposite can be defined as a composite structure whose characteristic dimensions are found at the nanoscale.8 The superhard nanocomposites9 of nc-TiN/SiNx

10

exhibit relatively good thermal and chemical stability. In the idealized case, this nanocomposite consists of crystalline TiN grains which are embedded in a tissue phase of amorphous Si3N4. A prerequisite to synthesize nanocomposites is a strong segregation tendency between the constituents in order to get a strong interface between the nanocrystals and the Si3N4 phase. This is the case for TiN and Si3N4, which have essentially no solid solubility; see the pseudo-binary phase diagram in Fig. 2.3. For the nanocomposites with the highest hardness, the grain sizes should be below 10 nm, and the tissue phase that separates the nanocrystallites on the order of 1-2 monolayers (ML) thick.11 The hardness enhancement is then explained by small crystallite sizes of TiN, which gives grain boundary hardening, together with inhibited grain boundary sliding and crack propagation from the Si3N4 phase.

N

Ti

4

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2.5 TiN/SiNx Multilayers

A multilayer thin film consists of alternating layers of two or more materials. The sum of two consecutive layers in a bilayer system is called multilayer period (�). A multilayer containing epitaxial layers is called superlattice. Also, multilayered structures with characteristic dimensions on the nanoscale are referred to as nanolaminates. Recently, publications regarding superhard TiN/SiNx multilayers have been published,2,12,13 where metastable c-SiNx have been epitaxially strained between TiN layers in an artificial superlattice structure. These hardness correlates to the thickness of the SiNx layer, and the hardness is highest for the case of superlattice with SiNx layer thickness of 1-2 ML. The high hardness is explained by hindering of dislocation motion.12

2.6 Ternary Solid Solutions

A solid solution can be defined as follows…“A solid solution is a solid-state solution of one or more solutes in a solvent. Such a mixture is considered a solution rather than a compound when the crystal structure of the solvent remains unchanged by addition of the solutes, and when the mixture remains in a single homogeneous phase. The solute may incorporate into the solvent crystal lattice substitutionally, by replacing a solvent particle in the lattice, or interstitially, by fitting into the space between solvent particles. Both of these types of solid solution affect the properties of the material by distorting the crystal lattice and disrupting the physical and electrical homogeneity of the solvent material.”14 It was stated in section 2.1 that there exist no thermodynamically stable ternary Ti-Si-N compounds. However, based on the use of our method for growth of metastable thin films in section 3.4, metastable (Ti,Si)N cubic solid solutions can in fact be realized. Figure 2.3 shows the pseudo-binary phase diagram for TiN and SiN and indicates a strong phase separation tendency from a solid solution into the binary phases. This implies, that for a metastable (Ti,Si)N solid solution, a phase separation into the binary phases may be expected during annealing. The region of �´+�´´ is the miscibility gap, where the cubic binary phases are preferred. The chemical spinodal is indicated by a dashed curve in Fig. 2.3. Within this curve the eventual decomposition is spinodal and outside of which towards the binodal (solid curve), phase separation by nucleation and growth would take place. Spinodal

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decomposition can briefly be described as ‘up-hill’ diffusion, in which atoms diffuse towards high-concentration regions.15

0,00 0,25 0,50 0,75 1,000

500

1000

1500

2000

2500

3000

3500

4000

0

500

1000

1500

2000

2500

3000

3500

Tem

pera

ture

(K)

x, Ti1-xSixN

Tem

pera

ture

(oC

)

δδδδ' + δ + δ + δ + δ''

δδδδ-Ti1-xSixN

TiN SiN

Fig. 2.3 Pseudo-binary phase diagram for TiN-SiN together with the chemical spinodal (dashed line), calculated down to 1727 °C (Liquid state not considered). From Paper I.

The pseudo-binary phase diagram was calculated in the following way. Gibbs free energy, G, of a system is defined by

G = H − TS Eq. 1

where H, T, and S are the system’s enthalpy, temperature, and entropy, respectively. The total energies were calculated by ab initio density functional theory (DFT) in Paper I for the Ti1-xSixN system at 0 K. This gives G = H in Eq. 1. To include the temperature and entropy dependence of G an ideal solution16 is assumed.

1 S. Sambasivan, W. T. Petuskey, J. Mater. Res. 9 (1994) 2362 2 Powder Diffraction Files, JCPDS International Center for Powder Diffraction Data, Swarthmore, 1989, card 6-642 3 H. Ljungcrantz, M. Odén, L. Hultman, J. E. Greene, J. –E. Sundgren, J. Appl. Phys. 80 (1996) 6725 4 H. Ljungcrantz, C. Engström, M. Olsson, X. Chu, M. S. Wong, W. D. Sproul, L. Hultman, J. Vac. Sci. Technol. A16 (1998) 3104 5 I. Tomaszkiewicz, J. Thermal. Anal. Cal. 65 (2001) 425

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6 A. Zerr, G. Miehe, G. Serghiou, M. Schwarz, E. Kroke, R. Riedel, H. Fuess, P. Kroll, R. Boehler, Nature 400 (1999) 340. 7 Aylward, A & Findlay, T. in Store (ed.) Si Chemical Data, 3rd Edition, Wiley & Sons Milton, 1994 8 http://www.uspto.gov/go/classification/uspc977/defs977.htm 9 S. Veprek, M. G. J. Veprek-Heijman, P. Karvankova, J. Prochazka, Thin Solid Films 476 (2005) 1 10 S. Veprek, S. Reiprich, Thin Solid Films 268 (1995) 64 11 A. Niederhofer, T. Bolom, P. Nesladek, K. Moto, C. Eggs, D. S. Patil, S. Veprek, Surf. Coat. Technol. 146/147 (2001) 183 12 H. Söderberg, J. M. Molina-Aldereguia, L. Hultman, M. Odén, J. Appl. Phys. 97 (2005) 114327 13 X. Hu, H. Zhang, J. Dai, G. Li, M. Gu, J. Vac. Sci. Technol., A 23 (2005) 114 14 http://en.wikipedia.org/wiki/Solid_solution 15 A. Hörling, PhD Thesis (Linköping Studies in Science and Technology, dissertation no. 922, Linköping University, Sweden 2005 16 D. A. Porter, K. E. Easterling, Phase Transfromations in Metals and Alloys, Chapman & Hall, 2nd ed. (1992) 14-15

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3 Thin Film Deposition

3.1 Physical Vapor Deposition

Thin solid films can be synthesized by physical vapor deposition (PVD) techniques. Generally, the coating material is vaporized in vacuum from a solid material, target. The vapor will eventually condense onto a substrate surface. Next follows descriptions of arc evaporation and DC magnetron sputtering, the PVD methods employed in this work.

3.2 Arc Evaporation

Arc evaporation has been widely used because of its promise of an efficient source of highly ionized material for producing dense, adherent coatings having a wide range of compositions.1 A cathodic arc can be described as a low voltage, high current plasma discharge between two electrodes. The evaporation process of target material is a consequence of the very high local surface temperature in an arc spot. This creates a molten pool from which evaporation of the cathode (or target) material and electron emission occurs. The electrons are then attracted by an electric field and will collide and ionize evaporated atoms; this is called the ionization zone, see Fig. 3.1. The ions are transported to the substrate surface where they condensate and react with a reactive gas (if present) from the surrounding. In Paper I, N2 was utilized as reactive gas. However, the molten pool also emits macro particles. To synthesize high quality thin films, the importance of plasma ionization2,3,4 should be emphasized. Arc evaporation, in contrast to sputtering, provides highly ionized plasmas and can therefore be manipulated with electric and magnetic fields5. Other ion induced effects are acceleration of the nucleation stage,6 enhanced adhesion,7 modification of crystal structure,3 film stress,8 densification, and in the case of deposition from a compound target, changes in stoichiometry.

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Fig. 3.1 Schematic illustration of particle flux at the arc spot.9

The changes in stoichiometry appear when arc evaporation is applied to a compound target where the different materials cause different degree of ionization. This gives the ions different acceleration towards a negatively biased substrate. Therefore, ions with higher degree of ionization will impinge on the surface with higher energy and thus penetrate deeper into the film compared to an element with lower degree. This will cause preferential resputtering of the surface near material and the film will contain a higher concentration of the material with higher degree of ionization.1 This phenomenon was apparent in Paper I, where a slightly higher Ti:Si ratio was observed in the film compared to the target. The average charge state during arc evaporation for Ti and Si is typically +2.1 and +1.4, respectively.10 However, during reactive arc evaporation, the average charge will be somewhat less.

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Fig. 3.2 Image of an arc evaporation deposition chamber at SECO Tools AB.

3.3 DC Magnetron Sputtering

The process of sputtering starts by introducing a sputtering gas, preferably inert, into a vacuum chamber. A high voltage is applied to the target; this creates a visible glow discharge, often referred to as plasma, by ionization of the inert gas. The gas ions will be accelerated towards, and eventually collide, with the negatively charged target. If the kinetic energy of the incoming ions is higher than the binding energy of the target surface atoms, the atoms will be ejected, sputtered. The ejected target material will vaporize and travel through the plasma to the substrate. Depending on the kinetic energy of the incoming coating material and the temperature of the substrate, ad-atoms may or may not migrate on the surface until they occupy an energetically favorable position. As the ions collide with the target they will also cause emission of secondary electrons. Since the electrons are negatively charged they will be repelled from the target and instead collide with other atoms and ions to free electrons. This will create positively charged ions to maintain the process. In Paper II reactive DC magnetron sputtering was used to deposit TiN/SiNx multilayers from two separate targets, where the reactive N2 gas was mixed with the inert Ar gas. In this process the vaporized target species in the plasma mix and react with N2

before they migrate on the substrate. In Fig. 3.3 a schematic view of the interior of the sputtering chamber is shown. The magnetrons were of unbalanced types which together

Target, cathode

Plasma

Substrates Trigger

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13

with magnetic coupling induced by a coil around the substrate holder enhances the ionization near the growing film. The ions will then be accelerated towards the negatively biased substrate, in this case -50 V.

Fig. 3.3 Schematic interior view of the deposition chamber.

3.4 Growth of Metastable Solid Solution Films

In Paper I metastable (Fig. 3.4) ternary solid solutions was synthesized by arc evaporation. Two mechanisms are described below for depositing metastable solid solutions.

Ene

rgy

Atomicarrangement

Ea

I

II

Ene

rgy

Atomicarrangement

Ea

I

II

Fig. 3.4 Schematic illustration of metastable (I) and thermodynamically stable (II) states. Energy Ea is needed to activate transformation from state I − to − II.

Substrate

Coil

Magnetrons

Shutter Ar inlet

N2 inlet

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3.4.1 Low-temperature Synthesis

Low temperature deposition is preferable from an industrial point of view, since it allows a wider range of substrate materials with respect to their thermal stability and cost. Utilizing lower temperatures also decrease the production time. Low temperatures can be at or below 500 °C; a typical process condition in arc evaporation. In addition, low-temperature film synthesis (being far from the thermodynamical equilibrium of the deposition material) induces kinetic limitations which, e.g., allow for synthesis of metastable phases.11 This is the first mechanism for the metastable ternary solid solution synthesized in Paper I.

3.4.2 Ion-induced Recoil Implantation

When a negatively biased target (or substrate in the thin film deposition case) is bombarded by the energetic ion beam, atoms will be redistributed in the target.12 These collisions also cause atomic recoils, which more precise are energetic ions traversing the target and goes on to generate a cascade of secondary atoms. The collisions induced by the ion bombardment cause recoil implantation, wherein target atoms are knocked downstream by collisions and ion-mixing. This process resembles thermal diffusion, but is driven by collisions rather than by thermal motions.13 The ion-induced recoil mixing is proposed as a second mechanism to form the metastable ternary solid solutions in Paper I.

1 R. L. Boxman, D. M. Sanders, P. J. Martin, J. M. Laferty, Handbook of Vacuum Arc Science, Fundamentals and Applications, Noyes Publications, New Jersey, 1995 2 D. Dobrev, Thin Solid Films 92 (1982) 41 3 J. M. E. Harper, J. J. Cuomo, H. T. G. Hentzell, J. Appl. Phys. 58 (1985) 550 4 E. Key, F. Parmigiani, W. Parrish, J. Vac. Sci. Technol. A6 (1988) 3074 5 A. Anders, Surf. Coat. Technol. 120-121 (1999) 319 6 M. Marinov, Thin Solid Films, 46 (1977) 267 7 J. E. Griffith, Y. Oiu, T. A. Tombrello, Nucl. Instrum. Methods, 198 (1982) 349 8 J. Cuomo, J. M. E. Harper, C. R. Guarnieri, D. S. Yee, L. J. Attanasio, J. Angilello, C. T. Wu, R. H. Hammond, J. Vac. Sci. Technol. 20 (1982) 349 9 R. L. Boxman, S. Goldsmith, Surf. Coat. Technol. 52 (1992) 39 10 I. G. Brown, IEEE transactions on plasma science 19 (1991) 713 11 I. Petrov, P. B. Barna, L. Hultman, J. E. Greene, J. Vac. Technol. A21(5) (2003) 117 12 S. M. Myers, Nucl. Instrum. Methods 168, 265 (1980) 13 I. Manning, Phys. Rev. E, B42 16 (1990) 9853

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4 Theoretical Modeling

To receive further understanding in why the materials of the present thesis behave as they do, theoretical modeling is used to investigate phase stability as well as structural and elastic properties. Modeling is an area that has expanded tremendously during the last decade due to increased computer power and more efficient program codes. However, in order to provide reliable results with physical meaning, a deep understanding and knowledge is necessary. In this thesis density functional theory was used to determine lattice parameters and total energy for Ti1-xSixN solid solutions (0�x�1) and for the determination of surface reconstructions for SiNx onto TiN.

4.1 Density Functional Theory

Today the density functional theory (DFT) formalism is the most used ab initio method in computational material science and solid-state physics. The reason for this is the high computational efficiency combined with high accuracy. Ab initio is a Latin term that means first principles. This implies that the calculation relies on basic and established laws of nature without additional assumptions or special models. In the 1960s Hohenberg and Kohn1 presented and proved two theorems, which became the fundament for DFT. The first theorem states that the external potential in which the electrons move, is a unique functional of the ground state electron density. This means that the systems are fully determined by the electron density. Hence, the total energy of the system can be expressed as a functional of the density. The second theorem states that the ground state electron density minimizes the total electronic energy of the system. The theory was then further developed by Kohn and Sham2 who used these theorems to derive the Kohn-Sham equations:

( ) ( ) ( ) ( ) ( )rrrrrr

rnnnxcext Evvd

ne

mψψ =�

�

���

�++

−+∇− � 'r

''

222� (Eq. 2)

where the external potential, vext(n(r)), is determined from the electronic density, n(r), instead of from the electron wave functions as for the general Schrödinger equation. �(r) is the time-independent wave function.

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4.1.1 Approximations for Many-body Interactions

The expression for the energy contribution from many-body interactions, which are captured by the exchange-correlation energy vxc(r), is unknown. However, different approximations have been developed; the generalized gradient approximation (GGA) and the local density approximation (LDA) are perhaps the most commonly used ones. In LDA the exchange-correlation energy is taken from known results of the many-body interactions in a uniform electron gas. This apparently easy approximation works surprisingly well for most applications3 and requires relatively short computational time. However, for more rapid changes in the electron gas, LDA seems coarse and great effort has been made to find better approximations. To overcome this problem gradient correction were included to the exchange-correlation potential. This, together with constraints on the exchange-correlation functions led to the implement of GGA. In a comparison between LDA and GGA, the latter tends to improve total energies4, atomization energies2,5,6, energy barriers, and structural energy differences7,8. GGA also expands and softens bonds6, an effect that sometimes corrects9 and sometimes overcorrects10 the LDA prediction.

4.1.2 Pseudo Potentials and Plane Waves

Cambridge serial total energy package (CASTEP) is a powerful simulation package from Accelrys Inc.11 In this package either GGA or LDA can be utilized. CASTEP uses pseudo potentials, which are approximations where the core electrons are treated as frozen. Since the computational time is heavily dependent on the number of electrons, this decreases the computational time dramatically. One type of pseudo potential method is the ultra soft pseudo potentials12, these were used in Paper II.

4.1.3 Linear Muffin-Tin Orbital

The full-potential linear muffin-tin orbital (FP-LMTO) method within LDA was used for the calculation in Paper I. Here, the unit cell is divided into non-overlapping muffin-tin spheres around the atoms. The potentials and charge densities in the crystal can have any, and not necessarily spherical, shape.

1 P. Hohenberg, W. Kohn, Phys. Rev. B, 136 (1964) B864 2 W. Kohn, J. Sham, Phys. Rev. A, 140 (1965) A1133

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3 W. Kohn, Rev. Modern. Phys. 71 5 (1999) 1253 4 J. P. Perdew, J. A. Chevary, S. H. Vosko, K. A. Jackson, M. R. Pederson, D. J. Singh, C. Fiolhais, Phys. Rev. B 46 (1992) 6671; 48 (1993) 4978 5 A. D. Becke, J. Chem. Phys. 96 (192) 2155 6 E. I. Proynov, E. Ruiz, A. Vela, D. R. Salahub, Int. J. Quantum Chem. S29 (1995) 61 7 B. Hammer. K. W. Jacobsen, J. K Norskov, Phys. Rev. Lett. 70 (1995) 3487 8 D. R. Hamann, Phys. Rev. Lett. 76 (1996) 660 9 V. Ozolins, M. Körling, Phys. Rev. B 48 (1993) 18304 10 C. Filippi, D. J. Singh, C. Umrigar, Phys. Rev. B 50 (1994) 14947 11 M. D. Segall, P. J. D. Lindan, M. J. Probert, C. J. Pickard, P. J. Hasnip, S. J. Clark, M. C. Payne, “First principles simulation: ideas, illustrations and the CASTEP code”, J. Phys. Cond. Matt. 14(11) (2002) 2717-2743 12 D. Vanderbilt, Phys. Rev. B 41 (1990) 7892

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5 Thin Film Characterization

To obtain necessary information about material properties, several characterization techniques are needed. In this work, investigation about composition, morphology, microstructure, thermal stability and mechanical properties using the techniques described below were performed.

5.1 X-ray Diffraction

In x-ray diffraction (XRD) an x-ray beam is scattered, or diffracted, by atoms in the investigated material. X-ray diffractograms are in principle created from the condition of constructive interference from Bragg’s law

2d·sin�=n� (Eq. 3)

which requires that the path difference between the traveled x-rays is equal to an integer number of wavelengths. In Eq. 3, d is the atomic plane spacing of the investigated crystal, � the wavelength of the x-rays and 2� is the angle of diffraction.

Fig. 5.1 Schematic of diffraction according to Bragg’s law.

XRD is a powerful analytical technique for microstructural investigations and can be used for characterization of crystal structure, phase transformations, residual stress, thickness measurement etc.

d 2 �

X - ray source

Detector

d 2 �

X - ray source

Detector

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5.2 Electron Microscopy

Transmission electron microscopy (TEM) is an invaluable analysis technique due to its ability of getting both a physical image and an electron diffraction pattern. In this thesis an FEI analytical TEM, Technai G2 UT FEG operating at 200 keV, equipped with scanning transmission electron microscope (STEM), energy loss electron spectroscopy (EELS), and energy dispersive x-ray spectroscopy (EDX) has been used for high resolution imaging and analytical analysis. A Philips EM 400T (120 keV) and a Philips CM 20 UT (200 keV) were used for the overview images and electron diffraction. A TEM is in many ways similar to a light optical microscope. Both types are built up with an illumination and an image part, where the first illuminates the sample and the second creates the image. However, instead of photons, electrons are irradiating the sample in TEM; the sample has to be viewed in vacuum in order to increase the mean free path of electrons. Also different lenses are used, instead of ordinary optical lenses, electromagnetic lenses are employed. Since electrons are accelerated with, in this thesis, 120 and 200 keV, respectively, their wavelength is on the 10-12 m scale, which can be compared to wavelength of photons, 10-7 m scale. The output is better resolution. In STEM the electron beam been focused to a small probe, i.e. a convergent beam, which scans over an area of the sample. Some imaging modes in STEM supply information that cannot be obtained in a conventional TEM, e.g. micrographs containing mainly z-contrast (described below). Due to the small probe size, chemical analysis can be performed in different fashions, e.g., EDX line-scan. TEM and STEM requires careful sample preparation in order to obtain an electron transparent thin area (<100 nm thick). Depending on the purpose of the analysis the sample are studied in cross-section or plan-view. The cross-sectional samples in this thesis were prepared by gluing two small slices cut out from the sample face-to-face and mount them in a titanium grid. This was continued by mechanically grinding until a thickness of about 50-60 µm were achieved. Finally, the sample was etched by a 2-5 keV Ar+ ion beam at a 5° incident angle in a precision ion polish sputter (PIPS) until electron transparency was achieved. TEM and STEM combined with the analytical techniques, EELS and EDX, give possibilities to measure chemical composition on the nanoscale with a point-to-point resolution in the subnanometer range.

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50nm

Figure 5.2 a) Cross-sectional scanning transmission electron micrograph, and b) transmission electron micrograph, of Ti0.86Si0.14N deposited onto WC(Co) and annealed at 1100 °C. From Paper I.

The STEM image in Fig. 5.2 a) is obtained by collecting high-angle scattered electrons collected with a high angle annular dark field detector (HAADF)1 at a low camera length. This configuration provides mainly z-contrast (thus, very little diffraction contrast) which gives bright contrast from heavy elements. In Fig. 5.2 a) the heavy elements at the grain boundaries correspond to W and Co according to a line-scan measurement with EDX/STEM. Compare also with the bright field TEM image in Fig. 5.2 b), were Z-contrast and diffraction contrast are present. In TEM the heavy elements have dark atomic number contrast. Note that some grains appear black due to diffraction contrast.

5.3 Nanoindentation

In nanoindentation an indenter deforms a material on a very small scale. During the indent displacement and indent load are continuously recorded and by using the Oliver and Pharr2 method, hardness – a material’s ability to resist deformation upon a load3 –, and Young’s modulus can be evaluated. To avoid influence from substrate the penetration depth should not exceed ~10%4 of the film thickness. It is important to take into account that the acquired hardness data is affected by the material’s microstructure (grain size and boundaries, voids etc.), loading orientation (for non-isotropic materials), and environment, i.e. temperature, humidity etc. The hardness is

a) b)

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thus a system response and property that should be evaluated using a statistical approach, i.e., each samples hardness should be evaluated from multiple indents.

5.4 Scanning Tunneling Microscopy

Scanning tunneling microscopy (STM) delivers images of a solid surface by moving a sharp conductive tip in a very precise and controlled manner across the sample surface and recording the electron tunneling current between the tip and sample as a function of position. The tip edge ideally consists of only one atom in order to provide atomic resolution. Tunneling is a quantum mechanical effect in which electrons from one conductor penetrate through a classically impenetrable potential barrier (for STM, vacuum) into a second conductor. The phenomenon arises from the leaking of the respective wave functions into the vacuum and their overlap within classically forbidden regions. This overlap is significant only over atomic-scale distances and the tunnel current depends exponentially on the distance between the conductors. Since an image of the surface is obtained in STM, the analysis has to be carefully performed, keeping in mind that there is always a risk for unwanted features, like for example double tip image, vibrations, and electrical noise. A double tip image forms when the tip picks up contamination from the sample surface or by other structural modification with features to which tunneling can take place. The image created from a double tip will contain doublets of the true surface features. Fig. 5.3 shows a micrograph created from a double tip. The Material Research Laboratory at the University of Illinois at Urbana-Champaign has a deposition system, equipped with one sputtering target and one sublimation* target, together with variable temperature scanning tunneling microscope (VT-STM) and low-energy electron diffraction (LEED)†; all in the same ultra high vacuum (UHV) system. Hence, new possibilities of in situ surface studies where depositions of the order of sub-monolayer directly can be probed by means of STM and LEED. This system was used in Paper II.

* Sublimation is a PVD process where the target is heated to make the target material evaporate from the surface. † LEED works in principle as ordinary electron diffraction, but is utilized at lower energies which makes it very surface sensitive and therefore can provide information from the topmost atomic layers.

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Fig. 5.3 Scanning tunneling micrograph from SiNx deposited on TiN scanned by a double tip. Note that all features appear twice in the image.

5.5 X-ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS) is a surface sensitive technique, which can provide information about elemental composition and chemical bonding. X-ray photons, typically Al or Mg K�, travels incident on a surface and eject valence or core electrons from a depth of 0-10 nm. The kinetic energy, Ekin, from the ejected photoelectrons are detected and the electron binding energy, EB, is calculated from

Ekin = h�−EB−� (Eq. 4)

where h� is the incident photon energy, and � a work function that corresponds to the x-ray source and the sample. The binding energy of a core electron does not only depend on the element but also on the surrounding atoms. The change in binding energy is often referred to as chemical shift. This information is of great interest since it could provide information about the surrounding atoms, e.g. if a Si atom is located in an octahedral (like atoms in a NaCl lattice) or in a tetrahedral (like atoms in a ZnS lattice) position. The

20nm

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25

energy resolution of the XPS also plays a central role to resolve these shifts. In Paper I XPS was used to determine the Si bond character.

1 E. M. James and N. D. Browning, Ultramicroscopy 78 (1999) 125 2 W. C. Oliver, G. M. Pharr, J. Mater. Res. 7 (1992) 1564 3 C. –M. Sung, M. Sung, Mater. Chem. Phys. 43 (1996) 1 4 R. Sburlati, Journal of Mechanics of Materials and Structures 1 (2006) 554

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6 Results

6.1 Ti1-xSixN Alloy Films

In Paper I Ti1-xSixN (0�x�0.14) thin films were deposited onto cemented carbide (WC-Co) substrates by arc evaporation at ~500 °C. Cross-sectional scanning electron microscopy showed that all films exhibited a columnar structure. A closer investigation using XTEM revealed that within each column for the x=0.14 film, a feather-like domain structure could be revealed. This is an effect originating from point defects and dislocations yielding low-angle grain boundaries. TEM together with XRD showed that the obtained films phases were of cubic NaCl type with a lattice parameter close to the reference value of �-TiN at 4.24 Å. Calculations by FP LMTO gives that the lattice parameter of NaCl Ti1-xSixN (x=0, 0.25, 0.5, 0.75, 1) are very similar (within 1%) of each other. This implies that it can be hard to distinguish SiNx from TiN in nanocomposites or multilayers or even from a (Ti,Si)N solid solution using XRD. Furthermore, the calculations in Paper I suggests that a NaCl-lattice is more favorable than a ZnS-lattice for Si-content x<0.67. The synthesis of the (Ti,Si)N alloy films is made possible by the kinetic limitations for atom mobilities at the chosen low substrate temperature (~500 °C) combined with the employed high-flux low-energy metal ion bombardment, which induces recoil mixing. Effects of N content on the phase stability and properties of SiNx polytypes, however, were not investigated. The microstructure of as-deposited solid solution films with a Si content of x=0.14 was retained up to an annealing temperature of at least 900 °C/120 min (see Paper I). This absence of phase separation −despite the deep miscibility gap− is likely due to a limited driving force for the nucleation of a Si3N4 phase due to molar volume mismatch (Si3N4 has a larger unit cell than the NaCl-structured solid solution). When annealing at 1100 °C/120 min, interdiffusion of W and Co occurs in the film grain boundaries which transforms the film into a nanocrystalline cellular microstructure. Mechanical properties of the films were investigated with nanoindentation. The measured hardness increased close to linearly with increasing Si content and is retained up to 900 °C. At 1100 °C the hardness decreased to below 30 GPa due to the Co and W diffusion, which weakens the grain boundaries. Also, initial deposition experiments were performed for the Ti1-xSixN system in the wider composition range (0�x�0.22) onto cemented carbide (WC-Co) substrates by arc evaporation. The as-deposited Ti0.78Si0.22N film exhibited a dense fine grained and

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columnar two-phase structure consisting of a defect-rich crystalline cubic phase and a nanocrystalline − to − amorphous structure as seen by cross-sectional and plan-view transmission electron micrographs, see Fig. 6.1. The crystalline structure is a saturated (Ti,Si)N solid solution similar to the Ti1-xSixN x�0.14 alloys. Due to the supersaturation of N2 gas in the deposition, the amorphous phase is likely to by rich in nitrogen as for the thermodynamically stable a-Si3N4. However, no z-contrast was obtained for the structure using XSTEM at low camera length (90 mm) which indicates that the two phases have a similar composition, i.e., a-Si3N4:Ti and c-(Ti,Si)N, respectively. The maximum hardness was reached for Si content of 0.135�x�0.175 and was of the same order as for the Ti0.86Si0.14N film in Paper I. Higher Si contents resulted in a leveling out or slight reduction of the hardness.

a-Si3N4

2 nm 200 nm

a) b)

a-Si3N4

2 nm 200 nm

a) b)

Fig. 6.1 Transmission electron micrographs from a Ti0.78Si0.22N film deposited by arc-evaporation, a) plan-view containing amorphous area of Si3N4 between defect-rich (Ti,Si)N crystallites, b) cross-sectional bright-field overview image showing diffraction contrast indicating location of amorphous and crystalline phases, respectively. A. Flink, T. Larsson, J. Sjölén, L. Karlsson, L. Hultman, unpublished.

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6.2 TiN/SiNx Nanolaminate Films

The nc-TiN/a-Si3N4 nanocomposites can exhibit extraordinary mechanical properties and thermal stability. The interface between TiN and SiNx is of great importance to understand since the atoms involved occupy a relatively large volume fraction of the nanocomposite. However, there is little knowledge about the bonding at the interface and for any tissue phase forming or effects of any segregated contamination. In Paper II, an interface study was performed with three different approaches. 1) Investigation of the interface between TiN and SiNx by using in situ STM to probe sub atomic layer coverages of SiNx deposited onto TiN(001)/MgO(001) and TiN(111)/MgO(111), 2) multilayers of TiN/SiNx deposited onto SiOx and MgO(001) which reduces the interface to a two dimensional, instead of a three dimensional, problem, and 3) surface calculations of different SiNx coverages placed onto TiN surfaces by using ab initio DFT methods. In the in situ UHV STM and LEED study, epitaxial TiN(001)/MgO(001) was deposited by magnetron sputtering at 700 °C. These substrates were then inserted in the vacuum deposition system equipped with LEED and VT-STM. The deposition system contains one Ti magnetron sputtering source and one Si evaporation source and N2 and Ar gas supplies. Here, a TiN(001) template layer containing atomistically flat terraces was sputtered. Then, different SiNx surface coverages, �SiNx, was deposited onto the TiN surface by sublimation of Si followed by annealing in N2 atmosphere during 12 h at temperatures ranging from 600 to 800 °C. Several different crystalline reconstructions were found for different �SiNx, containing rows in the <110> directions. The multilayers in Paper II were deposited by reactive magnetron sputtering at 500 °C onto MgO(001) substrates. The TiN layers were kept at fixed thickness, 40 Å, while the SiNx layer thicknesses (lSiNx) were varied between 3 and 25 Å. For a SiNx layer thickness up to 5 Å, both the TiN and SiNx layers were epitaxial, forming a superlattice. Increasing the SiNx thickness to 13 Å yielded a polycrystalline TiN layer structure alternating with SiNx layers that are initially crystalline to a thickness of 5-13 Å before becoming amorphous. The epitaxial stabilization of c-SiNx is explained by the minimization of surface area energy at the early stages of layer nucleation, i.e., instead of forming a high energy crystalline/amorphous interface a low-energy crystalline/crystalline interface is formed. However, when the interfacial strain energy, which depends on the thickness of the SiNx layer, is sufficiently large the epitaxial growth breaks down, and an amorphous Si3N4 layer forms. Figure 6.2 shows results from a high-resolution scanning transmission electron microscope study from the lSiNx = 13 Å TiN/SiNx multilayer sample from Paper II of

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29

several areas exhibit local epitaxy as in a superlattice. A HAADF detector was used with a camera length of 300 mm for the purpose of increased image intensity. Correspondingly, both diffraction and z-contrast are present in the images, which promote the TiN crystals to be more or less pronounced. The dark and bright contrast

Figure 6.2 High resolution scanning transmission electron micrographs of a TiN/SiNx multilayer with lSiNx=13 Å. The inset image shows local epitaxy of c-SiNx and TiN. A. Flink, H. Söderberg, P. O. Å. Persson, L. Hultman, M. Odén, unpublished.

correspond to SiNx and TiN, respectively. The higher-magnification inset shows that cubic SiNx is stabilized between the TiN layers. In Paper II also the hardness of the multilayers was determined by nanoindentation. Both multilayer series deposited on MgO(001) and SiO2 substrates had their maximum hardness (34±2 GPa) at lSiNx close to the epitaxial breakdown limit in the respective series. The SiNx/TiN superlattice films exhibit Koehler hardening1, in which dislocation generation and glide across interfaces is hindered for phases with a large difference in shear moduli. However, the maximum hardening observed in the present multilayer samples is significantly larger than expected from Koehler hardening. Thus, additional hardening mechanism, including strengthening due to the coherency strain between the

10 Å

TiN

SiNx

20 Å

TiN

SiN

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30

two epitaxial layers, (SiNx and TiN(001)) must be present. Effects of lattice strain can be seen in Fig. 6.2 with the meandering trace of {200} crystallographic planes in the successive TiN and SiNx layers. Finally, ab initio DFT calculations were used to calculate different surface formations. Three different coverages were tested; 0.4 ML with a Si to N ratio of 1, 0.9 ML with a Si to N ratio of 5/6, and 1 ML with a Si to N ratio of 1. The DFT calculations suggest that for increasing Si content a tetrahedral binding configuration as in hexagonal or amorphous Si3N4 is preferred over an octahedral as in NaCl-structure SiN. 1 Koehler, J. S., Attempt to Design a Strong Solid. Phys. Rev. B 2 (1970) 547

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PAPER I

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Influence of Si on the microstructure of arc evaporated (Ti,Si)N thin films;

evidence for cubic solid solutions and their thermal stability

A. Flink a,*, T. Larsson b, J. Sjolen c, L. Karlsson c, L. Hultman a

a Thin Film Physics Division, Department of Physics, IFM, Linkoping University, SE-581 83 Linkoping, Swedenb Ombenning by 14, SE-737 90 Angelsberg, Sweden

c SECO Tools AB, SE-737 82 Fagersta, Sweden

Available online 19 September 2005

Abstract

Ti1�xSixN (0�x�0.14) thin solid films were deposited onto cemented carbide (WC-Co) substrates by arc evaporation. X-ray diffraction

and transmission electron microscopy showed that all films were of NaCl-structure type phase. The as-deposited films exhibited a

competitive columnar growth mode where the structure transits to a feather-like nanostructure with increasing Si content. Films with

0�x�0.01 had a b111� crystallographic preferred orientation which changed to an exclusive b200� texture for 0.05�x�0.14. X-ray

photoelectron spectroscopy revealed the presence of Si–N bonding, but no amorphous Si3N4. Band structure calculations performed using a

full potential linear muffin tin orbital method showed that for a given NaCl-structure Ti1�xSixN solid solution, a phase separation into cubic

SiN and TiN is energetically favorable. The microstructure was maintained for the Ti0.86Si0.14N film annealed at 900 -C, while

recrystallization in the cubic state took place at 1100 -C annealing during 2 h. The Si content influenced the film hardness close to linearly, by

combination of solid-solution hardening in the cubic state and defect hardening. For x =0 and x =0.14, nanoindentation gave a hardness of

31.3T1.3 GPa and 44.7T1.9 GPa, respectively. The hardness was retained after annealing at 900 -C, while it decreased to below 30 GPa for

1100 -C following recrystallization and W and Co interdiffusion.

D 2005 Elsevier B.V. All rights reserved.

Keywords: Nitrides; Arc evaporation; Transmission electron microscopy (TEM); Thin films; Solid solution; Microstructure

1. Introduction

Advanced surface engineering of transition metal nitride

wear-resistant coatings by the introduction of alloying

elements is a growing field of research. TiN has been

widely used as hard coating on cutting tools, but the poor

oxidation resistance at temperatures above 500 -C [1,2] has

created an interest in ternary compounds, e.g., on Ti–Al–N

[3–5] and Cr–Al–N [6–8] and also the more complex

quaternaries, e.g., Ti–Al–Si–N [9,10]. These coating

materials show a much improved oxidation behavior at

high temperatures and are now used by market leaders

within metal cutting tools. Recently, Choi et al. [11] showed

for Ti–Si (11 at.%)–N compounds that Si forms SiO2 on

the surface which acts as an oxygen diffusion barrier up to

an annealing temperature of 800 -C for a duration of 150

min. Also, improvement on mechanical properties has been

realized with super hardness, H >45 GPa, for specific Si

contents [12,13] in TiN–Si3N4 nanocomposites.

The equilibrium phase diagram for Ti–Si–N does not

contain any stable ternary phases [14]. However, Vaz et al.

found phases originating from a possible (Ti,Si)N solid

solution [15]. The maximum Si content in metastable

supersaturated cubic solution Ti1�xSixN is 10–15 at.%

[14]. Furthermore, Prochazka et al. [16] conclude that a

totally segregated amorphous Si3N4 only can occur when

the nitrogen activity is larger than about 10�6. This suggests

that PVD is a preferred technique to suppress nitrogen

segregation by virtue of the lower substrate temperature

employed. Thus, during deposition by reactive magnetron

sputtering Soderberg et al. [17], first, and Hu et al. [18],

more recently, stabilized sub-nm thick layers of cubic SiN in

TiN/SiNx multilayers by epitaxial growth.

0257-8972/$ - see front matter D 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.surfcoat.2005.08.096

* Corresponding author.

E-mail address: axefl@ifm.liu.se (A. Flink).

Surface & Coatings Technology 200 (2005) 1535 – 1542

www.elsevier.com/locate/surfcoat

33

In this work, we show that as-deposited Ti1�xSixN thin

films with a Si content up to x=0.14 can be prepared by arc

deposition. Kinetic limitations during film deposition com-

bined with ion bombardment induced collisional mixing are

proposed as conditions for the growth of metastable

Ti1�xSixN films instead of the thermodynamically stable–

and much more studied–TiN/Si3N4 system. Interestingly,

film hardness was found to increase close to linearly with the

Si content. The as-deposited solid solutions exhibited thermal

stability above 900 -C for 2 h. Residual stress recovery and

recrystallization, however, resulted in transformation to a

nanocrystalline structure after annealing at 1100 -C. Bandstructure calculations performed using a full potential linear

muffin tin orbital method showed that for a given NaCl-

structure, a phase separation into cubic SiN and TiN phases is

energetically more favorable than a Ti1�xSixN solid solution.

2. Experimental details

Cemented carbide WC-Co (6 wt.%) 12�12�4 mm3

plates were used as substrates. Before the deposition the

substrates were ground and polished to a mirror-like finish,

Ra�0.01 Am, and cleaned in an ultrasonic alkaline

degreasing agent.

The films were deposited by a commercial arc evapo-

ration system. Three cathodes of composition Ti, Ti90Si10,

and Ti80Si20, respectively, located on top of each other in the

chamber were used to produce Ti1�xSixN films of varying

composition from one batch. Substrates with a bias of �50

V were kept at a temperature of 500 -C in an Ar/N2

atmosphere with N2-flow of 300 sccm.

Isothermal annealing of samples were performed in a

Sintevac Furnace from GCA Vacuum Industries. The

samples were annealed for 2 h at 900 -C and 1100 -C,respectively. The annealing experiments were performed in

an Ar flow at atmospheric pressure to prevent oxidation of

the sample surfaces.

The microstructure of the coatings was studied with X-

ray diffraction (XRD), cross-sectional transmission electron

microscopy (XTEM), and scanning electron microscopy,

(SEM). X-ray diffractometry was performed using a Philips

PW 1820 powder diffractometer with a line-focused Cu Ka

X-ray source. h –2h scans were recorded in the 2h range of

5- to 90-. A Philips EM 400T microscope operating at 120

kV was used for the overview imaging and an FEI Technai

G2 UT FEG microscope equipped with an electron energy

loss spectrometer, EELS, and an energy-dispersive X-ray

analysis spectrometer, EDX, operating at 200 kV was used

for the high-resolution imaging and the EELS analysis.

Chemical analysis of the film compositions was per-

formed using an Oxford Link ISIS EDX equipment,

operating at 20 kV, in connection with a LEO 1550 SEM.

Elemental mapping by EDX was measured for 30 min on a

plan-view sample magnified by 2.5k. The chemical bonding

structure in the near-surface region was analyzed by X-ray

photoelectron spectroscopy, XPS, using a VG Microlab

310F system. The XPS was equipped with a non-mono-

chromated Al Ka at 1486.6 eV X-ray source and a

hemispherical electron energy analyzer. To compensate for

eventual sample charging, the peak position of the

adventitious carbon was recorded before Ar-etching,

thereby setting the peak position to 284.75 eV. The samples

were then Ar-etched and survey scans of the binding energy

0–1100 eV was recorded with a step size of 1 eV for each

sample. For accurate determination of the exact peak

positions of the Si2p and C1s peaks, local region scans

were recorded with step size of 0.1 eV. To suppress the

background noise, each scan was recorded 10 times.

After mechanical polishing of the surface, nanoindenta-

tion analysis of films was performed using a Nano Instru-

ments NanoIndenter II with a Berkovich diamond tip. The

maximum indent load was 25 mN. The indentation

procedure is described in more detail by Horling et al. [5].

Ten indents in each sample were made to obtain statistically

reliable results. Indents in a bulk fused silica reference

sample were made with an indent load of 8 mN, yielding a

similar penetration depth, <200 nm, as in the investigated

coatings. The average hardness and Young’s modulus with

standard deviations was determined [19]. Poisson’s ratio, t,for TiN was set to t =0.22, as used, e.g., by Sue [20]. The

average hardness and Young’s modulus for the reference

sample, SiO2, was measured to 9.65 T0.4 GPa and

72.31T1.32 GPa, respectively.

3. Computational details

Ab initio calculations on the TiN–SiN system were

carried out using the full potential linear muffin tin orbital

method (FP-LMTO) [21,22] within the local density

approximation (LDA) of density functional theory (DFT).

The exchange correlation function used is described by

Hedin and Lundqvist [23]. Starting with the TiN structure,

and by exchanging Si for Ti atoms in the NaCl and

zincblende, ZnS, structure, respectively, a total of five

different compositions, x =0, 0.25, 0.5, 0.75, and 1, were

investigated. The unit cell dimensions for all structures were

varied uniformly and the theoretical size of the unit cells

was obtained from the minimum in total energy. Energy

convergence was reached for all compositions with respect

to the number of k points used.

4. Results and discussion

EDX analysis of as-deposited films from the different

positions within the deposition chamber showed a continuous

composition range between 0�x�0.14 Si for samples

positioned between Ti and Ti0.9Si0.1 or Ti0.8Si0.2 targets.

The film closest to the Ti0.8Si0.2 target was expected, due to

setup geometry, to have higher Si content than the actual

A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–15421536

34

x =0.14. However, the apparent loss of Si during arc

evaporating can be explained by the stronger impact from

Ti atoms than Si on the sample surface. Since the Ti atoms

have a higher grade of ionization, they will retrieve higher

acceleration towards the negatively biased substrate. This

will distribute the Ti atoms deeper into the structure than for

Si. Such a mechanism was shown to operate in the (Ti,Al)N

system [24]. In our case, preferential sputtering of the lighter

Si component is also possible.

Fig. 1a shows X-ray diffractograms from the as-deposited

Ti1�xSixN films. The diffractograms for all compositions

revealed a NaCl-structure compound with a lattice parameter

very close to TiN, 0.424 nm. The preferential film growth

orientation changed from mixed b111� to exclusive b200�

with increasing Si content. For x�0.07, however, the (200)

peak broadening increased substantially.

Cross-sectional TEM micrographs from the as-deposited

samples are presented in Fig. 2a (TiN), Fig. 2b

(Ti0.92Si0.08N), and Fig. 3 (Ti0.86Si0.14N), respectively. The

films exhibited a dense columnar structure where the

column width ranged from 100 to 400 nm. Interestingly,

for all films, the top surface correlated directly with the

substrate topography. This implies that 3D-island growth

and eventual faceting was effectively suppressed during the

deposition. From Figs. 2 and 3, it is evident that the Si

content also affected the structure and increased the defect

density. The as-deposited Ti0.86Si0.14N, see Fig. 3, exhibited

within each column a feather-like structure. Higher magni-

fication imaging revealed nm-structure of feathers (elon-

gated crystalline grains) with large strain contrast and moire

fringes from overlapping features. The high-resolution

electron micrograph (HREM) in Fig. 3b shows a typical

appearance of three feather features of the cubic (Ti,Si)N

phase with high defect density of dislocations. The

observations in Fig. 3 show an interesting growth mode

with a rotating lattice by branching into subgrains over each

column. Branching begins at the column boundaries and the

subgrains merge at the apparent stem of the columns. This

takes place to maintain the (002) growth surface. The

selected area electron diffraction pattern (SAED) in Fig. 3a

confirms the texture seen in XRD. No volumes of any

amorphous phase were found by the TEM analysis.

Fractured cross-sections from as-deposited Ti1�xSixN,

x =0, x =0.05, x =0.08, x =0.14, presented in Fig. 4, were

investigated by scanning electron microscopy. The micro-

graphs showed dense columnar microstructure with macro

particles incorporated to a similar density as for (Ti,Al)N

coatings. The thickness of the (Ti,Si)N coating ranged

between 1.6 and 2.0 Am. As-deposited Ti0.86Si0.14N

a)

b)

c)

34 36 38 40 42 44 46 48 50

WC (101)TiN (200)TiN (111)WC (100)

Inte

nsi

ty (

arb

. un

its)

2 Theta (degrees)

Co

34 36 38 40 42 44 46 48 50

Inte

nsi

ty (

arb

. un

its)

2 Theta (degrees)

WC (100) TiN (111) TiN (200) WC (101)

0.14

0.13

0.11

0.08

0.07

0.06

0.05

0.03

0.01

0

Co x:

0.14

0.13

0.11

0.08

0.07

0.06

0.05

0.03

0.01

0

34 36 38 40 42 44 46 48 50

Inte

nsi

ty (

arb

. un

its)

2 Theta (degrees)

WC (100) TiN (111) TiN (200) WC (101)

Cox:

0.14

0.13

0.11

0.08

0.07

0.06

0.05

0.03

0.01

0

x:

Fig. 1. X-ray diffractograms from Ti1�xSixN films in (a) as-deposited, (b)

annealed at 900 -C, and (c) annealed at 1100 -C states. The Si content of

each film is indicated.

a) b)

Fig. 2. Cross-sectional transmission electron micrographs of (a) TiN and (b)

Ti0.92Si0.08N films on WC-Co substrates.

A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–1542 1537

35

behaved as a fine structure when fractured in agreement

with the nanostructure seen by XTEM.

In Fig. 5a, the total energy of a relaxed NaCl type unit cell

of Ti1�xSixN is plotted as a function of Si content. As a

comparison the total energy of Ti1�xAlxN, Ti1�zZrzN, and

TiC1�xNx are included in the figure. The calculations show

that NaCl-structure (Ti,Si)N is metastable with respect to

phase separation into NaCl-structure SiN and TiN. Compar-

ison of the energy for ZnS and NaCl structure Ti1�xSixN, Fig.

5b, however, suggests that a NaCl-structure (Ti,Si)N solid

solution is energetically more favorable than a cubic ZnS-

structure only up to x¨0.67. The calculations of lattice

parameters for NaCl-structure Ti1�xSixN gave aTi0.75Si0.25N=

0.985a0, aTiN=0.986a0, and aSiN=0.99a0, where a0 is the

lattice parameter for TiN. Underestimation of the unit cell size

is normal for the local density approximation (LDA) within

density functional theory (DFT) as seen for TiN here. The

relatively constant a values are consistent with our exper-

imental results, c.f. Figs. 1a, 3, and 6). Also, Soderberg et al.

[17] was not able to distinguish the y-SiN and y-TiN phases

by XRD due to the small difference in lattice parameters. The

present findings, further, implies that the peak with a lattice

parameter of 0.429 nm observed in an alleged (Ti,Si)N solid

solution by Vaz et al. [15] is not from a NaCl-structure phase.

Results from XRD of Ti1�xSixN films isothermally

annealed at 900 -C and 1100 -C are presented in Fig. 1b

and c. The texture of the as-deposited samples was

maintained. However, the peak width decreases as annealing

temperature increases.

Fig. 6 shows a cross-sectional TEM micrograph from the

Ti0.86Si0.14N sample annealed at 1100 -C. The film now

exhibits a cellular structure of elongated ¨10-nm-wide

grains with a texture that is reminiscent from the columnar

d)

a)

b)

c)

Fig. 4. Fractured cross-sectional SEM micrographs from Ti1�xSixN, (a)

x =0, (b) x =0.05, (c) x =0.08, and (d) x =0.14 on WC-Co substrates.

a)

b)

Column boundariesFeather features

8°

I II III

(002)

(002)

Fig. 3. Cross-sectional TEM micrographs from an as-deposited

Ti0.86Si0.14N thin film on WC-Co substrate in, (a) overview with selected

area electron diffraction pattern and (b) HREM image. The columnar

microstructure with internal branching of subgrains of (002) crystallo-

graphic orientation is indicated in (a). The trace of (200) and (002) planes in

neighboring subgrains I– III with zone axes [010], [hk0], and [110],

respectively, after mutual rotation around [001] is shown in (b).

A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–15421538

36

structure with feather-like features of the as-deposited state.

This nanocrystalline structure can be seen from the higher

magnification insert image, Fig. 5b and c, to consist of

recrystallized grains, but no amorphous phase. We find that

the boundaries were formed to relax the as-deposited

structure that contained a high density of line and point

defects. It is a further characteristic observation that the

structure exhibited cell-walls between the grains, see Fig. 6b

and c. The cell-walls that have dark atomic number contrast

were rich in Co and W as explained below.

Elemental mapping by EDX in combination with SEM

from the Ti0.86Si0.14N sample annealed at 1100 -C, Fig. 7revealed that Co and W is leaking through the film. The

SEM/EDX and TEM/EELS/EDX analyses shows that such

high-temperature annealing results in W and Co diffusion

a)

b)

c)

Fig. 6. Cross-sectional TEM micrographs of a Ti0.86Si0.14N film on WC-Co

substrate annealed at 1100 -C for 2 h showing (a) an overview, (b) higher-

magnification view with a selected are diffraction pattern and (c) HREM

image of a recrystallized cubic-phase structure without amorphous material.

a)

b)

c)

0,00

0,04

0,08

0,12

0,16

0,20

0,24

0,28

0

100

200

300

400

500NaCl (δ)ZnS-zb

Composition, x

Ti1-xSixN

0,00 0,25 0,50 0,75 1,00

0,00 0,25 0,50 0,75 1,00

0,00 0,25 0,50 0,75 1,00

-0,02

0,00

0,02

0,04

0,06

0,08

0,10

0,12

-40

0

40

80

120

160

200

NaCl

E (

Ry/

cell)

E (

Ry/

cell)

Composition, x

Ti1-xAlxN

Ti1-xZrxN

Ti1-xSixN

TiC1-xNx

E (

meV

/ato

m)

E (

meV

/ato

m)

0

500

1000

1500

2000

2500

3000

3500

4000

0

500

1000

1500

2000

2500

3000

3500

Tem

per

atu

re (

K)

Ti1-xSixN

Tem

per

atu

re (

oC

)

-Ti1-xSixN

+ ,,,

Fig. 5. Ab-initio calculations showing (a) the total energy of one unit cell as

a function of composition for the Ti1�xAlxN, Ti1�zZrzN, Ti1�xSixN, and

TiC1�xNx systems in NaCl-structure, (b) comparison of which structure of

NaCl and ZnS that is most energetically favorable for different Si contents

and (c) phase diagram for TiN–SiN as calculated from the graph in (a),

together with the chemical spinodal, calculated down to 1727 -C. (Liquid

state not considered.)

A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–1542 1539

37

towards the film top surface via the grain boundaries formed

during the recrystallization. For comparison, Horling et al.

[3] reported Co interdiffusion in (Ti,Al)N films on similar

substrates annealed at 1250 -C for 2 h.

Similar to arc evaporated TiN [25], Ti(C,N) [25],

(Ti,Al)N [3], and (Ti,Zr)N [26], Ti1�xSixN solid solution

films undergo recovery from a growth-induced compres-

sive stress state during annealing above the deposition

temperature. From Fig. 1, the recovery is evident from a

decreasing peak broadening. The underlying processes can

be understood from the onset of N-related defect

annihilation with activation energy around 2 eV [27]

and defect-assisted diffusion on the Ti side with activation

energy comparable to that for surface diffusion of 3.5 eV

[28]. Based on its smaller atomic radius, however, Si

diffusion should not be the limiting factor for the

segregations.

Next we consider the tendency for phase separation of

the as-deposited Ti1�xSixN films. As was found by the ab

initio calculations, Fig. 5, the system exhibits a large

miscibility gap with a possible initial transformation path

to y-TiN and y-SiN or ZnS-structured SiN. The calculated

spinodal in Fig. 5 indicates that spinodal decomposition

can be preferred instead of nucleation and growth as the

phase separation mechanism for a wide range of compo-

sitions including 10–12 at.% Si (the percolation threshold

in nc-TiN/a-SiN) for most processing temperatures. That

state is of course in turn apt to transform to the equilibrium

phases y-TiN and Si3N4. From the present results, we find

that annealing temperatures in excess of 1100 -C would be

required to effectively reach closer to equilibrium. Consi-

dering the larger molar volume of the Si3N4 phases

compared to its cubic polytypes, the metastable cubic

states can be sustained. In the case of CVD deposition

[13,30] and other processes of relatively high surface

mobility of the elements, the strong segregation tendency

of the Ti and Si on the substrate surface during deposition

results in the direct formation of TiN and Si3N4 as in the

nc-TiN/a-Si3N4 nanocomposite thin films.

Fig. 8 shows how the Si content influences the hardness

and Young_s modulus of the as-deposited films and those

annealed at 900 -C and 1100 -C. The addition of Si

increased the hardness throughout the whole gradient series

for the as-deposited films from 31.3T1.3 GPa up to

44.7T1.9 GPa for Ti0.86Si0.14N. From XRD and TEM, it

is evident that the Ti0.86Si0.14N coating contains the highest

defect density. Thus, both solid solution and defect harden-

ing may be active. The 900 -C annealed samples retained

their hardness with a similar trend of Si content coupled to

hardness. At 1100 -C, however, the hardness has decreasedto below 30 GPa for all compositions. Veprek et al. [31]

suggested that an O content of as little as 1–1.5 at.% causes

a decrease of the hardness in nc-TiN/a-Si3N4 nanocompo-

sites if segregated to the grain boundaries. For the present

case of metastable cubic phase (Ti,Si)N, our analysis

showed that Co and W had diffused from the substrate into

a)

b)

520

540

560

580

600

620

640

660

680

700

720

740 as-dep. 900 1100

Yo

un

g's

mo

du

lus

(GP

a)

Composition, x

0,00 0,02 0,04 0,06 0,08 0,10 0,12 0,14 0,16

0,00 0,02 0,04 0,06 0,08 0,10 0,12 0,14 0,16

22

24

26

28

30

32

34

36

38

40

42

44

46

48

Har

dn

ess

(GP

a)

Composition, x

as-dep. 900 1100

Fig. 8. (a) Hardness and (b) Young’s modulus of as-deposited and annealed

Ti1�xSixN films at 900 -C and 1100 -C for 2 h.

SEMa)

N Kαb) Ti Kαc)

Co Kαd) W Lαe)

Fig. 7. (a) Scanning electron micrograph and (b–e) elemental maps of

different elements by EDX from an area containing the film, Ti0.86Si0.14N

annealed at 1100 -C, right part, and substrate WC-Co, left part, surface

(after local flaking). Bright contrast correlates to high concentration.

A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–15421540

38

the grain boundaries. The observed softening can then be

explained by grain boundary weakening by tissue phases,

e.g., CoSi, TiO2 or SiO2, the latter two assuming that

oxygen is segregated from within grains to the boundaries.

These are equilibrium phases expected from phase diagram

considerations all of which are soft in comparison to TiN

and SiNx compounds. In the form of a tissue phase, they can

be expected to weaken the matrix material.

The hardness of the as-deposited TiN film with 31.2 GPa

is significantly higher than of TiN (002) single-crystal films

of 20 GPa [32]. Defect hardening effects in arc evaporated

TiN, Ti(C,N) and TiC have, however, been reported earlier

[33]. It is due to small grain size and lattice point defects,

where the latter yields compressive residual stress state.

The Young’s modulus of the films, see Fig. 8, increased

slightly with increasing Si content for the as-deposited films

from 670 GPa, for x =0, to 700 GPa, for x =0.047. For

higher amount of Si, however, the Young’s modulus

decreased to ¨600 GPa. It should be noted that nano-

indentation of nitrides under compressive stress systemati-

cally gives high values for the moduli. Here, we are

concerned with the relative changes and conclude that the

hardening by Si substitution in Ti1�xSixN films is not

correlated to the strength of SiN per se.

X-ray photoelectron spectroscopy of the as-deposited

Ti0.86Si0.14N showed the presence of the elements Ti, Si, N,

as well as a very small amount of O after Ar-etching. In

addition, a pronounced sample charging was detected,

suggesting the presence of an insulating phase on the film

surface. This is in contrast to the as-deposited TiN sample,

which did not exhibit sample charging before or after Ar-

etching. Analysis of the Si2p peak, shown in Fig. 9, showed

a binding energy 100.9 eV for the Ti0.86Si0.14N film, which

suggests Si–N bonding, close to the reported value of a-

Si3N4 at 100.6 eV [34]. Amorphous silicon nitride, a-Si3N4,

is usually positioned at 101.8 eV [29]. No Ti–Si bonding

was seen in the present XPS operating with a ¨1.0 eV

resolution.

5. Conclusions

Ti1�xSixN solid solution films of NaCl-structure con-

taining at least x =0.14 Si can be synthesized by arc-

evaporation at a substrate temperature of 500 -C. For thehighest Si contents, however, XRD and TEM reveal a

defect-rich lattice whereas XPS profiles do not support the

presence of amorphous Si3N4. The findings are supported

by ab initio calculations which show that the pseudo-

binary TiN–SiN system with metastable solid solutions

exhibits a wide miscibility gap with respect to y-TiN and

NaCl-structured SiN or ZnS-structured SiN depending on

composition.

A new characteristic nanostructured feather-like growth

mode of films was observed for the Ti0.86Si0.14N samples as

the columnar microstructure exhibited internal branching of

subgrains with more or less continuous rotation to align

growth to a common (002) direction.

The as-deposited solid solutions exhibited thermal

stability up to 900 -C for 2 h where the hardness was

retained. Film hardness was close to a linear function of Si

content (from 31 GPa in TiN to 45 GPa in Ti0.86Si0.14N).

Residual stress recovery and recrystallization, however,

was induced by 1100 -C annealing. It resulted in

transformation to a nanocrystalline cellular structure that

was void of amorphous phases, but with concomitant

diffusion of W and Co from the substrate through the film

via the grain boundaries to form a tissue phase. This

resulted in grain boundary weakening which decreased the

film hardness.

Finally, our results have impact for the interpretation of

nanocomposite or multilayer (superlattice) TiN/Si3N4 films

as the presence of cubic (Ti,Si)N compounds must now be

considered.

Acknowledgements

The present work was performed under the Swedish

Foundation for Strategic Research (SSF) Low-Temperature

Thin Film Synthesis program with additional support from

SECO Tools AB. Dr. Per Persson and Dr. Hans Hogberg are

acknowledged for assistance with electron microscopy and

XPS analyses, respectively.

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A. Flink et al. / Surface & Coatings Technology 200 (2005) 1535–15421542

40

Erratum Paper I p. 1541: reference [4] reads: “A Hörling PhD Thesis...” should read: “W. –D. Münz, J. Göbel, Proc. 7th Int Conf. Vacuum Metallurgy, Linz, Austria 1985”

41

42

PAPER II

43

44

1

Toward understanding interface structure in

superhard TiN-SiN nanolaminates and nanocomposites

Lars Hultman 1,*, Javier Bareño2, Axel Flink1, Hans Söderberg3, Karin Larsson4, Vania Petrova2, Magnus Odén3, J. E. Greene2, and Ivan Petrov2

1 Department of Physics, Chemistry, and Biology (IFM), Linköping University, SE-581 83 Linköping, Sweden

2 Frederick Seitz Materials Research Laboratory and the Materials Science Department, University of Illinois at Urbana-Champaign, Urbana, Illinois 61801, USA

3 Engineering Materials, Luleå University of Technology, SE 971 87 Luleå, Sweden 4 Department of Materials Chemistry, The Ångström Laboratory, Uppsala University, P.O. Box 538, SE-751 21

Uppsala, Sweden * Communicating Author: LH; e-mail: larhu@ifm.liu.se

Nanostructured materials – the subject of much of contemporary materials research – are defined by internal interfaces, the nature of which are largely unknown. Yet, the interfaces determine the properties of nanocomposites and nanolaminates. An example is nanocomposites with extreme hardness �70-90 GPa [1,2], which is of the order of, or higher than, diamond. The Ti-Si-N system, in particular, is attracting attention for the synthesis of such superhard materials. In this case, the nanocomposite structure consists of TiN nanocrystallites encapsulated in a fully percolated SiNx “tissue phase” (1-2 monolayers thick) that is assumed to be amorphous [1-3]. Here, we show that the interfacial tissue phase can be crystalline, and even epitaxial with complex surface reconstructions. Using in situ structural analyses combined with ab initio calculations, we find that SiNx layers grow epitaxially, giving rise to strong interfacial bonding, on both TiN(001) and TiN(111) surfaces. In addition, TiN overlayers grow epitaxially on SiNx/TiN(001) bilayers in nanolaminate structures. These results provide insight into the development of design rules for novel nanostructured materials.

An intense research area is the synthesis of superhard (hardness H � 40 GPa [1]) nano- composites for use as wear-resistant coatings on tools and mechanical components as well as scratch-resistant coatings on optics. The nanocomposites are composed of nano- crystallites (�10 nm) of transition metal nitrides, carbides or borides encapsulated by 1-2 monolayers (ML) of a covalent nitride (e.g., Si3N4, BN, CNx, or diamond-like C) interfacial layer. Due to the small dimensions across the nanograins, nucleation and glide of dislocations is impeded, while the high cohesive strength of the thin intergranular tissue phase inhibits grain-

boundary sliding [4]. Together, these effects provide qualitative explanation for the observed superhardness of the nanocomposites. The pseudobinary TiN-SiNx system, which presently serves as an archetype in the quest for superhard nanocomposite materials [1-3, 5-8], exhibits strong surface segregation (TiN and Si3N4 have essentially no solid solubility), a prerequisite for self-organized nanocomposite formation during vapor phase deposition. Recently, the growth of superhard SiNx/TiN nanolaminates has also been reported [5-7]. Despite the critical dependence of nanocomposite properties on the nature of the

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interfaces between constituent phases, there is a complete lack of basic knowledge regarding the three-dimensional structure of the interfacial phase. The nanocurvature of such interfaces presents an extreme challenge to the use of transmission electron microscopy (TEM) and other standard analytical methods. We approach the problem of isolating and probing SiNx/TiN interface chemistry and structure by preparing planar interfaces in the form of bilayers (heterostructures), trilayers, and multilayers (superlattices or nanolaminates) starting with well-defined TiN(001) and TiN(111) surfaces. In order to minimize contamination effects, film growth experiments are performed in ultra-high-vacuum (UHV). In situ UHV variable-temperature scanning tunneling microscopy (VT-STM), low-energy electron diffraction (LEED), high-resolution cross sectional electron microscopy (HR-XTEM), nanoindentation measurements, and ab initio density functional theory (DFT) calculations are used to provide atomistic information regarding bonding and crystallographic order at SiNx/TiN interfaces. As a first step, we determine the growth mode of SiNx layers in SiNx/TiN nanolaminates deposited on MgO(001) at 500 °C by reactive dual magnetron sputtering of high-purity Si (99.99%) and Ti (99.97%) targets in mixed Ar/N2 atmospheres (pAr = 4.0 mTorr, pN2 = 0.5 mTorr; 99.9997%). The thickness of the TiN layers is maintained constant at 40 Å (~19 ML) while varying the SiNx layer thickness (lSiNx) from 3 Å (~1.5 ML) to 25 Å (~12 ML). Figure 1(a) is a typical HR-XTEM image from a nanolaminate with lSiNx = 5 Å (~2.5 ML). Both the SiNx(001) and TiN(001) layers grow epitaxially to form a superlattice. HR-XTEM images reveal continuous lattice fringes across successive layers with no indication of the presence of an amorphous or polycrystalline SiNx phase. In fact, all nanolaminate samples with lSiNx � 5 Å are epitaxial SiNx/TiN(001) superlattices. X-ray reflectivity (XRR) scans of films with lSiNx = 5 Å and lSiNx = 13 Å exhibit

FIG. 1. Cross-sectional high-resolution electron microscopy images from TiN/SiNx multilayers grown at 500 °C on MgO(001), with a substrate bias of -50 V, by reactive dual magnetron sputtering of Si and Ti targets in an Ar/N2 discharge. The TiN layers are 40 Å thick, while the SiNx layer thickness is (a) 5 Å and (b) 13 Å. (c) A higher-magnification image of the sample in (b). (d) Image from a region in sample (b) exhibiting local epitaxial growth across the entire the 13 Å thick SiN layer. superstructure reflections which have a sharpness and periodicity that denote well-defined continuous layers. Elastic recoil detection analyses (ERDA) of our SiNx/TiN(001) multilayers reveal, in all cases, that the N/(Ti+Si) ratio is ~1.00. Thus, x=1 for the SiNx layers since the TiN(001) layers are grown under conditions known to provide stoichiometry [9]. In addition, we have established that the N/Ti ratio of thick TiN layers is 1.00±0.03 using a combination of x-ray diffraction (XRD) and Rutherford backscattering spectroscopy. θ-2θ XRD scans of SiN/TiN(001) superlattices exhibit only single peaks with no

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splitting at 00l positions, consistent with our DFT calculations showing that the lattice parameter misfit between NaCl-structure TiN and isostructural metastable SiN is ~0.5%. The only impurities detected by ERDA are O (�0.1 at.%), C (~0.01 at.%), and H (~0.01 at.%). Increasing lSiNx to 13 Å, however, leads to a transition in SiNx layer growth from epitaxial to amorphous (see Figs. 1(b) and 1(c)). The combination of XTEM, HR-XTEM, XRR, and XRD analyses shows that the film consists of polycrystalline TiN layers, with strong 002 preferred orientation, alternating with continuous SiNx interlayers which are initially crystalline (and exhibit local epitaxy with underlying TiN grains) to a thickness of 5-13 Å before becoming amorphous. We observe a similar transition from epitaxial to amorphous SiNx growth in multilayers deposited on MgO(001) at 300 °C and 700 °C. We assume that the amorphous SiNx phase has the composition Si3N4 since growth was carried out in a highly N2-rich ambient. The HREM images in Figs. 1(c) and 1(d) show that local epitaxy of cubic-SiN occurs on some TiN(001) grains over the complete 13 Å layer thickness. Our results show that there is an epitaxial breakdown thickness (see Bratland et al. [10] for discussion of breakdown mechanisms during low-temperature strained layer growth [10]), l*SiNx ~2.5-6.5 ML (5-13 Å), beyond which

growing SiNx layers become amorphous. We attribute the epitaxial stabilization of metastable cubic SiN on TiN(001) to pseudomorphic forces. However, the corresponding interfacial strain increases linearly with lSiNx [11] and as lSiNx > l*SiNx, the strain energy becomes sufficiently large that the epitaxial growth front breaks down. We infer that strain energy minimization is responsible for the reduced N content in metastable cubic-SiN epitaxial layers compared to the equilibrium phase, hexagonal-structure β-Si3N4, whose lattice constant (co = 2.90 Å; ao = 7.59 Å) mismatch with that of TiN (ao = 4.24 Å)

is >10%. Parallel SiNx/TiN growth experiments using SiO2 substrates (oxidized Si(001); Ts = 500 °C) with lSiNx = 3 Å (~1.5 ML) result in polycrystalline nanolaminates with local epitaxial growth of SiN on individual TiN grains. In order to probe the dependence of the hardness H of our SiNx/TiN multilayers on lSiNx, we carried out nanoindentation measurements on two sets of 0.5-µm-thick samples grown at Ts = 500 °C with lTiN maintained constant at 40 Å. A Nanoindenter II system with a Berkovich diamond indenter was used with loads of 2-5 mN. The first sample set (Series 1) was grown on MgO(001) substrates (the initial layer is epitaxial TiN(001)), while the second set (Series 2) was deposited on SiO2 (the initial TiN layer is polycrystalline). H(lSiNx) data are plotted in Fig. 2. For both sample series, the maximum hardness (34±2 GPa) is obtained at lSiN ~ l*SiNx. The lower l*SiNx value observed for the SiNx/TiN nanolaminates compared to the SiN/TiN(001) superlattices is due to the higher crystalline quality and lateral homogeneity of the epitaxial TiN interlayer. A similar hardness dependence on lSiNx has

been reported for TiN-SiNx nanocomposite films [1,2]. The nanocomposites reach their maximum hardness over a narrow SiNx thickness range near the percolation limit at ~1-2 ML. Figure 2 shows that this thickness is well within the pseudomorphic SiN regime for SiNx/TiN nanolaminates containing polycrystalline TiN layers grown on SiO2. Our SiN/TiN(001) superlattice films exhibit Koehler hardening [12,13] in which dislocation glide is hindered across interfaces between phases with large differences in shear moduli (G). In this case, GSiN = 42 GPa (calculated using DFT), while GTiN = 217-278 GPa [14,15]. For glide along maximum shear directions <110> oriented 45˚ to the [001] loading direction with a Taylor factor of 0.3, we estimate,

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FIG. 2. Nanoindentation hardness of 0.5-µm-thick SiN/TiN(001) superlattices grown on MgO(001) at 500 °C (Series 1) and SiNx/TiN nanolaminates deposited on Si(001) at 500 °C (Series 2) as a function of the SiNx layer thickness lSiNx. The TiN layer thickness is 40 Å in both cases. For comparison, the hardnesses of TiN(001) and amorphous Si3N4 films are 20 and 24 GPa, respectively. following the analysis in [16], a hardness increase of 8 to 8.7 GPa for SiN/TiN(001) superlattices compared to TiN(001) layers (20 GPa [17]). The maximum hardening observed in our samples is, however, significantly larger than that predicted by the Koehler effect. Thus, additional hardening mechanisms, including strengthening due to the coherency strains between the two epitaxial layers, SiN and TiN(001), must be present. For probing SiNx/TiN interface formation at the atomic level, we use in situ UHV STM and LEED to follow SiNx structural evolution during the early stages of growth on TiN(001). In these experiments, high-quality epitaxial TiN(001) template layers exhibiting 1×1 LEED patterns are grown on MgO(001) following the procedures developed in [5,18]. Next, Si is deposited by thermal evaporation on TiN(001) at room temperature and subsequently nitrided by exposure to N2 at pN2 ~2 × 10-8 Torr for up to 12 h at temperatures T ranging from 600 to 800 ºC. LEED analysis of layers with SiNx coverages �SiN < 0.30 ML reveals the coexistence of c-3×3 and 1×5 surface phases, oriented along <110> directions. As �SiN approaches 0.30 ML, the

intensity of the c-3×3 reconstruction spots decrease and the 1×5 pattern dominates. At �SiN = 0.5 ML, the 1×5 spots become streaks along (110) reflections. Figure 3(a) shows STM and LEED results for �SiN = 0.32±0.05 ML. The LEED pattern exhibits a 90º-rotated two-domain 1×5 reconstruction along orthogonal <110> directions. The corresponding STM image is composed of two domains, each with a characteristic “corn-cob” atomic-scale morphology exhibiting a height modulation of ~1 Å peak-to-peak. The width of the five-row repeat structure, four “corn rows” and a missing row, along both [110] and [1-10] is 14.4±1.1 Å, which is in agreement with the TiN interplanar spacing of ~3 Å. SiN deposition experiments on TiN(111) with �SiN values ranging from 0.1 to 0.3 ML also lead to epitaxial growth of NaCl-structure SiN. In this case, LEED patterns exhibit a 2×2 reconstruction, with weakly spotted streaks along (100), and STM images also show SiN row formation.

FIG. 3. In situ STM images and corresponding 100 eV LEED patterns from samples consisting of (a) 0.32 ML SiN on TiN(001) exhibiting a two-domain 1×5 reconstruction and (b) 3 ML TiN on 0.84 ML SiN on TiN(001) exhibiting a 1×1 surface. STM images were obtained using an Omicron VT/STM operated in constant current mode (1.5 V, 0.12 nA).

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To further investigate the interfacial self-assembly of TiN-SiNx nanocomposites, we deposited 3 ML of TiN, following the same procedure as with the initial TiN(001) layer, onto SiN/TiN(001) bilayers with �SiN = 0.85 ML. Figure 3(b) is an STM image from such a sample. The TiN overlayer grows epitaxially on SiN(001) and the surface exhibits the same 1×1 LEED pattern observed for the initial TiN(001) template layer. Moreover, the TiN(001) islands have their equilibrium shape; nearly square with <110> edge facets and <100> corner facets [21]. Furthermore, the measured <110>/<100> island step-edge anisotropy, � = 0.83, is in good agreement with previous results for TiN/TiN(001) homoepitaxy [19]. To provide additional insight into the atomic structure of the SiN/TiN(001) interface and the row-like SiN surface reconstruction, we employ ab-initio DFT to determine minimum energy structures as a function of �SiN. The calculations, with periodic boundary conditions, were carried out using the generalized gradient approximation of Perdew, Burke, and Ernzerhof [20] for electronic exchange and correlation. Figure 4 shows the atomic-scale geometry of 0.9 ML SiN on TiN(001), following relaxation of the interfacial region, in (a) cross-section and (b) plan view. The results were obtained by placing Si and N on TiN lattice positions (Si on top of NTiN and N on top of TiTiN) and with every 5th Si [-110] row empty. The relaxed structure geometry is comparable to the STM results (see Fig. 3). Moreover, the SiN overlayer assumes a cubic SiN structure. Si binds to N supplied during nitridation as well as to the surface, forming a surface reconstruction with a repeat distance of five TiN unit cells along <110>, in agreement with the experimental results. There is an outward relaxation of the N atoms, in both the SiN overlayer and the upper TiN planes, similar to that observed in vacuum-terminated TiN(001) surfaces [21]. N atoms in the SiN layer move closer to Si-Si bridge positions, with characteristic Si-N-Si bond angles of�111°, 127°,

[110]

(a)

[110]

[001

](b)

[-110

]

[110]

(a)

[110]

[001

](b)

[-110

]

FIG. 4. DFT (a) cross-sectional and (b) plan views showing relaxed interfacial atomic positions for 0.9 ML SiN on TiN(001). Ti: grey; N: blue; Si: yellow. Bond lengths and angles in the bottom TiN layer were fixed. A vacuum slab of 3 ML in the <001> direction was used. Atoms in the SiN overlayer are enlarged for clarity. and 143°, and in-plane bond lengths between 1.66 and 1.87 Å. Out-of-plane Si-N bond lengths are 1.84 ± 0.01 Å, while Ti-N bonds are somewhat longer than in bulk TiN (2.12 Å); between 2.40 and 2.43 Å. The original TiN(001) surface also reconstructs (see Fig. 4(a)). Ab initio DFT calculations were used to determine the minimum-energy configuration with �SiN = 1 ML. Two-domain, 90°-rotated, SiN row structures along <110> directions were observed. Eighty percent of the Si atoms retain cubic lattice positions almost perfectly above the surface Ti atoms, similar to the �SiN = 0.9 ML case, but with additional pronounced

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reconstruction due to Si atom pairs forming a dimer row above the SiN(001) surface. Si atoms in the upper (non-dimer row) are covalently bonded to their four nearest N neighbors (3 NSiN

+ 1 NTiN) in a distorted tetrahedral sp3 configuration. The calculated structures indicate that SiN surface coverages � 1 ML result in tetrahedral environments, while lower �SiN values favors octahedral environments. In addition, for �SiN = 1 ML, the range of Si and N bond lengths with �SiN = 1ML is 1.60-1.87 Å, close to the reported value for Si-N compounds, 1.74 Å [22]. Si-NTiN bonds have a strong covalent character and a high local electron density (between 0.45 and 1.01 electrons/bond), comparable to diamond (1 electron/bond). Epitaxial stabilization of non-equilibrium phases in thin films and multilayers has been well documented for a wide variety of materials systems [23] including cubic AlN (whose equilibrium structure is hexagonal wurtzite) on NaCl-structure TiN(001) [24]. Another example is the growth of 9-12 ML of cubic AlN in AlN/TiN(001) superlattices (lattice mismatch = -3.84%) before transformation to the equilibrium structure [25]. The stabilization is due to a reduction in interfacial energy during nucleation of the metastable heterolayer and a decrease in the overall strain energy during subsequent island coalescence and layer growth. In the present case of c-SiN/TiN(001), the misfit is much smaller (0.5%) than for AlN/TiN(001), yet the epitaxial breakdown thickness is only 2-3 ML. There are two primary differences between c-AlN/TiN(001) and c-SiN/TiN(001). (1) For both the equilibrium and metastable AlN phases, the stoichiometric ratio is 1, while for SiNx, N/Si = 1 for the epitaxial cubic phase and 1.33 for the amorphous phase. As a consequence, our DFT calculations show that the Si bonding coordination changes from octahedral (6-fold with N and/or Ti) for c-SiN to tetrahedral (4-fold with N) for Si3N4. (2) In the AlN case, both the

cubic and hexagonal phases exhibit a lattice mismatch (i.e., misfit strain) with TiN(001), while for SiNx there is no strain associated with amorphous layer growth. Thus, the growth transition from epitaxial cubic SiN to amorphous Si3N4 phase takes place at lSiNx � l*SiNx in response to a change in both composition and bonding coordination during elastic strain relaxation. The results presented here suggest that the interfacial structure of superhard TiN-SiNx nanocomposites and nanolaminates are much more complex and play a more important role in controlling film properties than previously thought. Figure 5(a) is a schematic illustration of the originally proposed phase structure, adapted from Ref. [1-3], of superhard TiN-Si3N4 nanocomposites. It consists of nm-sized TiN crystallites surrounded by a 1-1.5 ML thick amorphous Si3N4 tissue phase at the percolation threshold (~10 vol.%) [1,2]. Figure 5(b) includes our results with epitaxial cubic SiN layers growing on low-index facets of TiN crystallites to a maximum thickness of 2.5-6.5 ML (5-13 Å) before the transition to amorphous Si3N4. The epitaxial breakdown thickness may be extended if the SiN tissue phase is between two TiN crystallites with which it forms coherent interfaces. This situation corresponds to the TiN/SiN(001) multilayers discussed above, for which maximum hardness is reached at a SiN thickness of 2-5 ML, well above the percolation limit. In summary, we report the first results on the structure and chemistry of the interfaces which control the properties of TiN-SiNx nano- composites and nanolaminates. Using direct observations supported by ab initio calculations, we show that the SiNx/TiN interfaces are far more complex than previously considered, including the epitaxial stabilization of a metastable cubic SiN phase on low index TiN crystal surfaces. A variety of surface reconstructions are also observed as a function of SiN coverage. For SiN growth on TiN(001), a

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Fig. 5. Cross-sectional schematic diagrams of superhard TiN-SiNx nanocomposites consisting of TiN nanocrystallites encapsulated by a SiNx tissue phase with volume fraction just above the percolation limit. (a) Two-phase model adapted from Refs. [1-3] in which SiNx is amorphous Si3N4. (b) Three-phase model in which the SiNx layer adjacent to low-index TiN planes is either epitaxial cubic SiN (lSiNx � 2-3 ML) or a bilayer consisting of cubic SiN and amorphous Si3N4 (lSiNx > 3 ML). Ti: grey; N: blue; Si: yellow. c-3×3 surface reconstruction is obtained with 0 < �SiN < 0.3 ML, while 0.3 � �SiN � 0.5 ML corresponds to a 90°-rotated two-domain 1×5 reconstruction along <110>. On TiN(111), SiN forms a 2×2 reconstruction at coverages between 0.1 and 0.3 ML. We also show that TiN grows epitaxially on SiN/TiN(001) bilayers, recovering

the 1×1 LEED pattern observed for the initial TiN(001) template. Continued heteroepitaxial growth can be maintained in SiN/TiN(001) superlattices wih SiN layer thicknesses of up to a few ML giving rise to greatly enhanced hardness. A transition to amorphous Si3N4 layer growth follows with increasing lSiNx > l*SiNx due to bonding strain in the SiN layer as Si atoms prefer to be tetrahedrally coordinated with N. These results suggest new design strategies for nanoscale materials in which nanocomposite systems and processing conditions are purposefully selected in order to obtain tissue phases which exhibit local epitaxy with the encapsulated crystallites. This provides higher interfacial bond strength, which further reduces the probability of grain boundary sliding and thus yields enhanced materials’ strength. One can envision a new class of super-to-ultra hard all-crystalline ceramic nanocomposites formed during primary (film deposition) or secondary (age processing) phase transformations in which both the interfacial structure and the overall preferential crystallographic orientation is controlled. ACKNOWLEDGEMENTS LH, KL, and MO acknowledge support from the Swedish Research Council and the Swedish Foundation for Strategic Research. JB, VP, JEG, and IP gratefully acknowledge the financial support of the Department of Energy, Division of Materials Science, and the use of the facilities of the Center for Microanalysis of Materials, University of Illinois, which is partially supported by the U.S. Department of Energy under grant DEFG02/91/ER45439. Prof. Stan Veprek and Dr. Andreas Bergmaier are acknowledged for kind assistance with ERDA measurements. Dr. Jens Birch is acknowledged for useful discussions. The DFT calculations were performed using the CASTEP package [26], developed by Accelrys Inc., San Diego, CA.

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