Functional structure design of new

high-performance materials via atomic

design and defect engineering (ADDE)

edited by

Prof. David Rafaja

Imprint

Copyright: Technische Universität Bergakademie Freiberg, Spitzentechnologiecluster ADDE

Publishing: SAXONIA Standortentwicklungs- und -verwaltungsgesellschaft mbH

Technical editing: Alexander Eisenblätter, Dr. Uta Rensch

Layout: Alexander Eisenblätter, Susann Müller

Print: SDV Direct World GmbH, Dresden

All rights reserved

No part of this book may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, microfilming, recpording or otherwise, without written permission from the Publisher.

The authors are responsible for the content of their publication as well as completeness and correctness of literature references cited. The publisher has performed only editorial changes to the original manuscripts.

ISBN 978-3-934409-68-2

Preface

Modern industry requires instantly new materials with tailored properties and energy efficient technologies for their production. Technical University Bergakademie Freiberg is one of the European universities, which are permanently active both in the materials science and engi-neering and in the development of modern technologies for the advanced materials produc-tion. Based on this long-time tradition of the Freiberg University, the development of modern high-performance materials with high functionality and efficiency for applications in the fields of communication, mobility, energy and environment was naturally the main goal of the Centre of Excellence “Functional structure design of new high-performance materials via atomic de-sign and defect engineering (ADDE)”, which was established in 2009 and funded until the end of 2014 by the European Regional Development Fund (ERDF) and by the Ministry of Science and Art of Saxony.

The main idea of the Cluster of Excellence ADDE was to control the crystal structure and mi-crostructure defects in order to tailor the properties of materials. This book presents an over-view of the results of 19 interdisciplinary projects, which dealt with the generation and manip-ulation of desired microstructure features in functionalised materials like metastable phases, controlled phase decomposition, precipitation and nano-sized structures, or with the elimina-tion of unwanted defects in large crystals intended for special electronic applications like for-eign atoms or dislocations. The main aim of the Cluster of Excellence ADDE was to understand the interactions between individual crystal defects and microstructure features as a first step towards targeted defect engineering. The choice of the materials was stimulated by the topics, which are established at the TU Bergakademie Freiberg and at the principal cooperation part-ners, which were the Helmholtz Centre Dresden Rossendorf and the Leibnitz Institute for Solid State and Materials Research Dresden. As the education and training of young professionals and academics for the Saxonian industry and research was one of the central tasks of the Centre of Excellence ADDE, a close cooperation with the local industry played a very important role.

The contributions in this book are divided into four groups. The first one is devoted to the technologies for production of thin silicon solar cells. It comprises the growth and processing of silicon ingots, the methods for their characterisation and the technologies for recycling of sawing slurries. In the second group of the contributions, selected materials for microelectron-ics, information storage and sensor technology are presented like the wide-gap semiconductor GaN, the dielectrics TiO2, SrTiO3, ZrO2 and TaZrOx, the azulene and anthraquinone structures for information storage and the conductive coordination polymers for transducers in sensor applications. In the third part of this book, protective coatings for wear reduction and corrosion protection are discussed with a special focus on the use of metastable phases. The last group of the contributions is dedicated to the mechanical properties and thermodynamics of light metal alloys.

David Rafaja November 2015Coordinator of the ADDE project

Contents

TECHNOLOGIES FOR PRODUCTION OF SILICON SOLAR CELLS

Ultra-short time processing of silicon solar cellsS. Prucnal, F. L. Bregolin, K. Krockert, H. J. Möller, W. Skorupa

Growth and characterization of multi-crystalline silicon ingotsE. Schmid, C. Funke, Th. Behm, S. Würzner, O. Pätzold, V. Galindo, M. Stelter, H. J. Möller

Experimental and numerical investigations of the formation of surface defects during machining of silicon wafersM. Budnitzki, T. Behm, M. Kuna, H. J. Möller

Development of a new process for recycling of used sawing slurries from solar industryI. Nitzbon, A. Obst, U. Šingliar, M. Bertau

MATERIALS FOR MICROELECTRONICS AND SENSOR TECHNOLOGY

Defect engineering in GaN layers grown by hydride vapor phase epitaxyG. Lukin, O. Pätzold, M. Stelter, M. Barchuk, D. Rafaja, C. Röder, J. Kortus

Strontium titanate – Breaking the symmetryH. Stöcker, J. Hanzig, F. Hanzig, M. Zschornak, E. Mehner, S. Jachalke, D. C. Meyer

Atomic layer deposition of dielectric thin films in the ternary system TiO2-SrTiO3B. Abendroth, S. Rentrop, W. Münchgesang, H. Stöcker, J. Rensberg, C. Ronning, S. Gemming, D. C. Meyer

Synthesis and characterization of Ge nanocrystals embedded in high-k materials for alternative non-volatile memory devicesD. Lehninger, P. Seidel, M. Geyer, F. Schneider, A. Schmid, V. Klemm, D. Rafaja, J. Heitmann

Novel molecular materials for information storage – Synthesis, electronic properties and electrode designM. Mazik, E. Weber, N. Seidel, S. Förster, E. Kroke, J. Wagler, A. Kämpfe, J. Kortus, T. Hahn, S. Liebing, Y. Joseph, R. Dittrich

Development of an electrically conductive coordination polymer based transducer for sensor applicationsM. Günthel, J. Hübscher, F. Katzsch, R. Dittrich, Y. Joseph, E. Weber, F. Mertens

6

26

42

64

78

100

118

134

146

166

PROTECTIVE COATINGS AND HARD MATERIALS

On the thermal stability of nanoscaled Cr/ta-C multilayersU. Ratayski, Ch. Schimpf, T. Schucknecht, U. Mühle, C. Baehtz, M. Leonhardt, H.-J. Scheibe, D. Rafaja

Defect engineering in Ti-Al-N based coatings via energetic particle bombardment during cathodic arc evaporationCh. Wüstefeld, M. Motylenko, D. Rafaja, C. Michotte, Ch. Czettl

Experimental and numerical assessment of protective coatings deposited by high velocity oxygen fuel flame spraying: Spraying process and thermo-mechanical behaviorS. Roth, M. Hoffmann, C. Skupsch, M. Kuna, H. Biermann, H. Chaves

Synthesis, properties and potential applications of rocksalt-type aluminium nitride (rs-AlN)K. Keller, M. R. Schwarz, S. Schmerler, E. Kroke, G. Heide, D. Rafaja, J. Kortus

MECHANICAL PROPERTIES AND THERMODYNAMICS OF LIGHT METAL ALLOYS

Influence of multi-pass roll-bonding on the mechanical properties of twin roll cast magnesium sheetsF. Schwarz, St. Reichelt, L. Krüger, R. Kawalla

Mg-Al composite wiresE. Knauer, J. Freudenberger, A. Kauffmann, L. Schultz

A unified approach to identify material properties from small punch test experimentsM. Abendroth

Atomistic modeling of defects in the framework of the modified embedded-atom methodS. Groh

Thermodynamic investigations in the ternary Al-Ti-Cr systemM. Kriegel, O. Fabrichnaya, D. Heger, D. Chmelik, D. Rafaja, H. J. Seifert

List of authors

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260

278

288

306

328

346

Defect engineering in Ti-Al-N based coatings via energetic particle bombardment during cathodic arc evaporation

Christina Wüstefeld 1, Mykhaylo Motylenko 1, David Rafaja 1, Claude Michotte 2, Christoph Czettl 3

1 Institute of Materials Science, Technische Universität Bergakademie Freiberg, D-09599 Freiberg, Germany

2 CERATIZIT Luxembourg S.à.r.l., L-8201 Mamer, Luxembourg3 CERATIZIT Austria GmbH, A-6600 Reutte, Austria

200

AbstractThis study describes the impact of the bias voltage during reactive cathodic arc evaporation (CAE) on the microstructure development in Ti-Al-N based coatings and their thermal stabil-ity. For this purpose, Ti-Al-N monolayers and Ti-Al-N / Al-Ti-(Ru)-N multilayers were depos-ited from powder metallurgical Ti100-xAlx (x = 40, 50, 60, 67) targets and Ti50Al50 / Ti33-yAl67Ruy (y = 0, 1, 5) targets, respectively. In order to vary the energy of the impinging ions, the CAE dep-osition was done at different bias voltages (UB) between -20 V and -120 V. The microstructure of the coatings was characterized with the aid of glancing angle X-ray diffraction (GAXRD) in the as-deposited and annealed state. GAXRD was done ex situ as well as in situ during anneal-ing up to 950 °C. The microstructure investigated by X-ray diffraction was described in terms of the phase composition, the stress-free lattice parameter, crystallite size and macroscopic lat-tice strain of the face centred cubic (fcc) (Ti,Al)N phase. Furthermore, analytical transmission electron microscopy was performed in conjunction with electron energy loss spectroscopy and X-ray energy dispersive spectroscopy. The microstructure properties were correlated with the hardness of the samples. It was found that the microstructure of the Ti-Al-N based coatings can be designed by the applied bias voltage during the CAE deposition. With increasing bias voltage, higher concentration fluctuations of titanium and aluminium within the fcc-(Ti,Al)N phase and a rising defect density were observed. The high defect density and the higher con-centration fluctuations were accompanied by the formation of minor segregations of Al-rich fcc- (Al,Ti)N at high bias voltages. At high UB the formation of wurtzitic AlN in the Al-rich coatings was apparently retarded in the as-deposited coatings but at elevated temperatures the decomposi-tion was accelerated.

1. Introduction

Ti-Al-N coatings were presented already in 1985/1986 as industrially promising protec-tive coatings [1, 2]. Especially their oxidation resistance [2, 3] and hardness at elevated tem-peratures [4, 5] make them attractive as pro-tective coatings for high speed drills [6] and cutting tools. Thus, already in 1989 Ti-Al-N coated cutting tools were produced for the commercial market [7]. In the following time, further improvements of the coating proper-ties and their thermal stability were aspired by the addition of doping elements like Si [8, 9], B [10, 11], Nb [10], Ta [11], Hf [10], Ru [12, 13] etc. to Ti-Al-N based coatings or by the deposition of a multilayer architecture [11, 14].

The Ti-Al-N based hard coatings are main-ly produced by physical vapour deposition methods (PVD), which facilitate the for-mation of the metastable face centred cubic Ti1-xAlxN phase. The highest experimentally observed solubility limit of Al in fcc-Ti1-xAlxN was x = 0.67 [15, 16], which agrees with ab

initio calculations [17, 18] that revealed the same value. If the metastable fcc-Ti1-xAlxN phase is exposed to high temperatures, it un-dergoes spinodal decomposition and forms Ti-rich and Al-rich fcc-(Ti,Al)N domains, as it was observed by Hörling et al. [16] and Mayrhofer et al. [4] in the temperature range of ~860 - 900 °C. This decomposition step is accompanied by an age hardening effect [4]. At higher temperatures, the authors of Refs. [4, 16] observed the transformation of fcc-AlN into its thermodynamically stable wurt- zite (w) modification. This decomposition step was frequently reported to take place above 1000 °C [4, 14, 16, 19-24]. Furthermore, the formation of w-AlN was considered for a long time to be detrimental for the coating properties like hardness [25-27] and adhesion [5, 28]. In the recent time, it could be shown that w-AlN can form during annealing al-ready in the temperature range of 850 - 900 °C [29-35] and thus coexists with fcc Al-rich (Al,Ti)N. This implies that also a direct trans-formation of fcc-(Ti,Al)N to w-(Al,Ti)N is

201

possible [36]. A hardness increase after an-nealing at 900 °C and the presence of minor amounts of w-AlN that were reported in Refs. [30, 35] support our idea [8, 37-39] that the presence of w-AlN is not generally harmful with regard to the hardness. In particular, the hardness is enhanced through the partial coherent interfaces between wurtzite and fcc phase [8]. The age hardening effect reported for the Ti-Al-N based coatings at ~ 900 °C e.g. in Refs. [4, 16, 30, 35] can be activated during the metal cutting. It could be shown that the temperature at the cutting edge of Ti0.6Al0.4N coated WC-Co cutting inserts was 850 - 900 °C during dry cutting of carbon steel [40].

The properties of Ti-Al-N-based coatings are determined by the microstructure fea-tures like phase composition, crystallite size and lattice strain. The microstructure can be adjusted on the one hand by the Al concen-tration and on the other hand by the depo-sition parameters. The most used techniques for the deposition of the hard coatings based on Ti-Al-N are cathodic arc evaporation and magnetron sputtering (MS). Both deposition techniques are usually carried out in reactive mode, where aluminium and titanium are transferred from the solid phase of the tar-get to the vapour phase. Nitrogen is provided as reactive gas. The generated plasma differs substantially for both methods. In case of MS, energetic ions of an inert gas, usually argon, are produced in a glow discharge and eject at-oms of the target material due to the momen-tum transfer. The fraction of the ionized metal atoms is very low in MS. In contrast to MS, 50 to 100  % of the evaporated metallic spe-cies are ionized [41] in the CAE process as a result of the evaporation of the target material by a high energy arc [42]. The CAE process represents an energetic deposition. If a neg-ative bias voltage is applied to the substrate with respect to the plasma potential, the ions are accelerated towards the growing film. The energy of the ion impact can be modified by the bias voltage which influences the surface mobility of the deposited species.

This study addresses the microstructure for-mation of Ti-Al-N based coatings that were deposited by cathodic arc evaporation. On the one hand, the influence of the Al content and the bias voltage on the microstructure and thermal stability of Ti1-xAlxN monolay-ers were investigated. On the other hand, the effect of the Ru addition and the bias voltage on the microstructure and thermal stability of Ti-Al-N / Al-Ti-Ru-N multilayers were ana-lysed.

2. Experimental

The Ti-Al-N monolayers [39] and Ti-Al-N / Al-Ti-(Ru)-N multilayers [43] were produced in an industrial CAE facility of the Balzers RCS type [44]. The deposition chamber is equipped with six arc sources that are in-stalled at the circumference of the chamber. On two arc sources, pure titanium targets were mounted that were used for the deposi-tion of a ~200 nm thick TiN adhesion layer, which was deposited prior the actual coating. The deposition of the actual coating was done from mixed Ti-Al-(Ru) targets that were in-stalled on the other four arc sources. All tar-gets were produced by Plansee CM using a powder metallurgical route [45]. In case of the Ti-Al-N monolayer coatings, four composi-tions series were deposited via four different target compositions: (I) Ti60Al40, (II) Ti50Al50, (III) Ti40Al60 and (IV) Ti33Al67. In order to modify the ion impact during the deposition process, each composition series was made at four different bias voltages of -20 V, -40 V, -80 V and -120 V. In this way, a matrix of 4 × 4 samples was produced.

The Ti-Al-N / Al-Ti-(Ru)-N multilayers were deposited from two Ti50Al50 and two Ti33-yAl67Ruy targets. Three multilayer series with different Ru content were produced by the variation of the Ru concentration in the target, namely: (I) y = 0, (II) y = 1 and (III) y = 5. Each multilayer series (I to III) was de-posited at four different substrate bias voltages of -20 V, -40 V, -60 V and -80 V. The CAE dep-osition of all coatings was done in a nitrogen

202

(1)

atmosphere with a working pressure of 3.2 Pa and at a temperature of 450  °C. A two fold rotation was applied during deposition. As substrates polished cemented carbides con-taining 12 wt.% Co and mixed carbides (hex-WC, fcc-TiC) with SNUN 120412 geometry according to DIN ISO 1832 [46] were used.

Additionally, monolithic Ti-Al-N coatings that were deposited onto cemented carbide inserts were investigated [47]. These coatings were reactively deposited from a Ti40Al60 tar-get at a nitrogen pressure of 5 Pa using CAE. The deposition was done in the π300 appara-tus from PIVOT plc with one vertical rotat-ing target in the middle of the chamber. The deposition temperature was ~500 °C and a three-fold rotation was applied. Within this contribution, the vertically positioned sam-ples that were deposited at three different bias voltages of -20 V, -40 V and -80 V as well as the horizontally positioned samples deposited at UB = -80 V are considered.

The chemical composition of the as-deposit-ed coatings was analysed by wavelength-dis-persive electron probe microanalysis (EPMA / WDS) done on a JXA 8900 RL from Jeol. Using EPMA/WDS, the weight fractions of Ti, Al, O and Ru were determined. Since the nitrogen content could not be measured directly due to the overlap of the spectral line Ll of titanium with the Kα1 line from nitro-gen, the nitrogen content in the coating was calculated from the analytical total assuming that N is the complement to the analysed el-ements. The nitrogen content was checked by glow discharge optical emission spectroscopy (GDOES) done on SPETRU-MAT 750 from Leco. The GAXRD experiments were done on a D8 diffractometer (Bruker AXS) that was equipped with a sealed X-ray tube with copper anode; further details are given in Ref. [39]. The angle of incidence of the primary beam at the sample surface was 3°.

(2)

where the constants , , and arerelated to the single-crystal elasticcompliance constants of cubic materials, fordetails see Ref. [49], and

𝑎𝑎𝜓𝜓ℎ𝑘𝑘𝑘𝑘 = 𝐴𝐴2𝑎𝑎0𝜎𝜎 sin2 𝜓𝜓 + 𝐵𝐵2𝑎𝑎0𝜎𝜎 Γℎ𝑘𝑘𝑘𝑘sin2 𝜓𝜓 + 2𝐵𝐵1𝑎𝑎0𝜎𝜎Γℎ𝑘𝑘𝑘𝑘 + 2𝐴𝐴1𝑎𝑎0𝜎𝜎 + 𝑎𝑎0,

𝑎𝑎𝜓𝜓ℎ𝑘𝑘𝑘𝑘 = 𝐴𝐴2𝑎𝑎0𝜎𝜎 sin2 𝜓𝜓 + 𝐵𝐵2𝑎𝑎0𝜎𝜎 Γℎ𝑘𝑘𝑘𝑘sin2 𝜓𝜓 + 2𝐵𝐵1𝑎𝑎0𝜎𝜎Γℎ𝑘𝑘𝑘𝑘 + 2𝐴𝐴1𝑎𝑎0𝜎𝜎 + 𝑎𝑎0,

𝑎𝑎𝜓𝜓ℎ𝑘𝑘𝑘𝑘 = 𝐴𝐴2𝑎𝑎0𝜎𝜎 sin2 𝜓𝜓 + 𝐵𝐵2𝑎𝑎0𝜎𝜎 Γℎ𝑘𝑘𝑘𝑘sin2 𝜓𝜓 + 2𝐵𝐵1𝑎𝑎0𝜎𝜎Γℎ𝑘𝑘𝑘𝑘 + 2𝐴𝐴1𝑎𝑎0𝜎𝜎 + 𝑎𝑎0,

GAXRD was employed to characterize themicrostructure in terms of phasecomposition as well as stress-free latticeparameter ( ), microstrain ( ),macroscopic lattice strain ( ) and residualstress ( ) of the fcc-(Ti,Al)N phase [32]. Thestress-free lattice parameter of the fcc-(Ti,Al)N phase in the Ti-Al-N monolayerswas determined from the linear dependenceof the lattice parameters on [32],where is the inclination of the diffractionvector from the sample surface normaldirection. The Poisson’s ratio was adoptedfrom Ref. [48], which gives ν = 0.3 for fcc-TiN. In case of the Ti-Al-N / Al-Ti-Ru-Nmultilayers, the vs. plots revealed astrong anisotropy of the elastic latticedeformation of the fcc-(Ti,Al)N phase. Dueto the medium [ ratio of 0.53 ±0.01, the lattice strain was higher in the

direction than in the direction,, which corresponds to an

anisotropy factor .

Due to the strong observed anisotropy of theelastic deformation, first of all thedependence of on and on thediffraction indices was described using Eq.(1) [49]:

203

(3)

(4)

The symbol denotes the out-of-plane lat-tice parameter, which is obtained together with from the linear regression (Eq. (5)). The macroscopic lattice strain can be consid-ered as more reliable than the residual stress, since is determined from the XRD data us-ing just two parameters, the stress-free lattice parameter and the Poisson’s ratio [32].

(5)

The microstructure on the nanoscale was vis-ualized with transmission electron microsco-py (TEM) using a JEM 2010 FEF and a JEM 2200 FS from Jeol for the Ti-Al-N monolay-ers and the Ti-Al-N / Al-Ti-(Ru)-N multi-layers, respectively. The Ti-Al-N monolay-ers were prepared in plan-view orientation; details are given in Ref. [39]. The Ti-Al-N / Al-Ti-(Ru)-N multilayers were prepared as cross-sections using focused ion beam tech-nique. Fast Fourier transformation (FFT) was applied locally to the high resolution (HR) TEM images in order to perform the local phase identification, to determine the grain orientations and to describe the orientation relationship (OR) between fcc-(Ti,Al)N and wurtzite phase. The multilayer cross-sections were additionally investigated in scanning TEM modus in order to obtain information about the fluctuations of the metallic spe-cies across the layered structure using X-ray energy dispersive spectroscopy (EDS) and electron energy loss spectroscopy (EELS) of the Ti L2,3 edge. Additionally, the nitrogen K edge was recorded and analysed [50]. The generation of small electron probes and the operation in scanning modus over extended sample areas was achieved by the microscope JEM 2200 FS that is equipped with a spherical aberration (Cs=0.5) corrector. The EELS and EDS data were analysed using the software DigitalMicrographTM from Gatan.

Instrumented nanoindentation experiments were done on locally polished surfaces using the nanohardness tester from CSM Instru-ments equipped with a diamond Berkovich indenter in order to determine the indenta-tion hardness (HIT) of the coatings. The load of 30  mN was applied and removed with a loading rate of 60 mN/min. The recorded load displacement curves were analysed using the Oliver and Pharr method [51]. The coating thickness ranged between 4 and 5 µm and the penetration depth was less than 10 % of the thickness. The average indentation hardness was determined from at least 15 indentations.

were approximated with the linear depen-dence

is the cubic invariant. Afterwards, the Γ-dependent part of Eq. (1) was subtracted from the measured lattice parameters. Finally, the lattice parameters independent of diffrac-tion indices hkl,

with ν=0.3 [48]. For the calculation of the re-sidual stress a Young’s modulus of 500 GPa was assumed [32]. This assumption could be supported by nanoindentation experiments performed on the Ti-Al-N based coatings that contained fcc-(Ti,Al)N as major phase. These nanoindentation experiments yield-ed an indentation modulus in the range of 450-550 GPa. Additionally to the residual stress, the macroscopic lattice strain ε was determined by using a modification of Eq. (4)

204

In order to investigate thermally activated mi-crostructure changes in the coatings, anneal-ing experiments were done on selected coat-ings. The Ti-Al-N monolayer coatings were annealed in vacuum up to 850 °C, details of the annealing sequence are given in Ref. [33]. During annealing of the Ti-Al-N monolayer coatings, in situ GAXRD experiments were performed using synchrotron radiation, see Ref. [33]. For the thermal treatment of the Ti-Al-N / Al-Ti-(Ru)-N multilayers coatings, the samples were encased in silica glass tubes that were filled with argon gas at a pressure of 500 mbar. The glass tubes were annealed at the oven temperatures of 450 °C, 650 °C, 850 °C and 950 °C for 60 min. It is assumed that the temperature at the coating surface is slight-ly below the oven temperature and cannot be compared directly with the temperatures given for the in situ GAXRD experiments. The GAXRD experiments of the Ti-Al-N/Al-Ti-(Ru)-N multilayers coatings were done ex situ.

3. Results and Discussion

3.1 Chemical composition and phase composition of the samples

The metal ratios in the as-deposited Ti-Al-N monolayer [39] and Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings [43] as measured by EPMA / WDS are given in Table 1. The meas-urement revealed a lower [Al] / Σ[Me] ratio (where Σ[Me] = ([Ti] + [Al] + [Ru])) in the coatings as compared to the average Al to metal ratio of the utilized targets. This ef-fect is in accordance with Refs. [52-54] and is caused by an increased ionisation of titanium and by a higher re-sputtering yield of already deposited aluminium atoms. The reduction of the [Al] / Σ[Me] ratio is more pronounced in the Ti-Al-N / Al-Ti-(Ru)-N multilayers than in the monolayer coatings. In the multilay-ers, the Al to metal ratio is decreased approx. 10 % as compared to the targets. In case of the Ti-Al-N monolayer coatings, the reduction of the [Al] / Σ[Me] ratio increases from ~5 % to ~8 % for the coatings deposited from Ti60Al40 and Ti33Al67 targets, respectively. The applied UB had no effect on the chemical composition. GDOES confirmed these results and showed that the coatings contain 50 at.% of nitrogen.

Table 1: Metal ratios of the Ti-Al-N monolayer and Ti-Al-N / Al-Ti-(Ru)-N multilayers as determined by EPMA / WDS where Σ[Me] = ([Ti] + [Al] + [Ru]).

Coating Series Targetsystem combination

Ti-Al-N monolayers

Ti-Al-N / Al-Ti-(Ru)-N multilayers

4 × Ti60Al40

4 × Ti50Al50

4 × Ti40Al60

4 × Ti33Al67

2× Ti50Al50 & 2× Ti33Al67

2× Ti50Al50 & 2× Ti32Al67Ru1

2× Ti50Al50 & 2× Ti28Al67Ru5

0.62 ± 0.01

0.53 ± 0.01

0.44 ± 0.01

0.38 ± 0.01

0.48 ± 0.01

0.47 ± 0.01

0.45 ± 0.01

0.38 ± 0.01

0.47 ± 0.01

0.56 ± 0.01

0.62 ± 0.01

0.52 ± 0.01

0.53 ± 0.01

0.53 ± 0.01

0

0

0

0

0I

II

III

0.005 ± ≤ 0.001

0.02 ± ≤ 0.001

TiΣ[Me]

AlΣ[Me]

RuΣ[Me]

205

The phase composition of all coatings was concluded from the intensity of the diffrac-tion lines corresponding to the expected crystalline phases w-AlN and fcc-(Ti,Al)N. Each diffraction line in the range of 30° to 150° was fitted by a symmetrical Pearson VII function as shown exemplarily in Fig.  1 for the low-angle part of the GAXRD patterns of the Ti-Al-N monolayer coatings deposited at UB = -40 V and UB = -80 V. Further infor-mation about the phase composition was de-duced from the comparison of the stress-free lattice parameter of the fcc-(Ti,Al)N phase with the expected lattice parameter accord-ing to the Vegard-like dependence given in Eq. (6) [32, 38, 55].

(6)

Fig. 1: Low angle parts of the GAXRD patterns measured for the Ti1-xAlxN monolayer coatings that were deposited onto cemented carbide substrates at a bias voltage of -40 V (a) and -80 V (b). The measured data is shown by black points. The grey lines show the individual diffraction peaks as fitted by Pearson VII functions. The positions of the fcc-TiN, fcc-AlN and w-AlN diffraction lines are labelled at the bottom of the figure. The positions of fcc-TiC and hex-WC coming from the substrate are shown at the top of the figure.

The phase analysis revealed a different phase composition of the coatings that had the same chemical composition but were deposited at low bias voltages (UB = -20 V and UB = -40 V) and at high bias voltages (UB = -60  V, UB = -80 V and UB = -120 V).

The phase composition of the Ti-Al-N mon-olayer coatings is summarized in Fig. 2a. The Ti0.62Al0.38N and Ti0.53Al0.47N coatings depos-ited at UB = -20 V and UB = -40 V contained fcc-(Ti,Al)N as a single phase, since no dif-fraction lines of wurtzitic phase could be observed in the diffraction patterns. The for-mation of traces of wurtzite phase beside the major fcc-(Ti,Al)N phase could be recognized for an Al content of x = 0.56 (see Fig. 1a) at low bias voltages. If the Al content was fur-ther increased to x = 0.62, the wurtzite phase became the major phase at low bias voltages. The shift of the diffraction lines of the wurt-zite phase in the coatings from the expected

206

2θ positions of w-AlN with the lattice param-eter a = 0.311 nm and c = 0.498 nm could be attributed on the one hand to the expansion of the elementary cell of the wurtzite phase indi-cating the incorporation of Ti into the wurt-zite phase. The expansion of the elementary cell of the wurtzite phase was also observed in the monolithic Ti0.36Al0.64N and Ti0.39Al0.61N coatings deposited onto cutting inserts by a different CAE process [47]. On the other hand, the shift of the diffraction lines of the wurtzite phase could also be affected by resid-ual stress, which is expected to be lower in the coatings deposited at low UB than in the coat-ings deposited at high UB.

The application of a high UB during the dep-osition of Ti-Al-N coatings suppressed the formation of the wurtzite phase, because fcc-(Ti,Al)N was the major phase in all Ti-Al-N monolayer coatings that were depos-ited at UB = -80  V and UB = -120  V. At high bias voltages, traces of the wurtzite phase could be recognized in the diffraction pat-terns only for an Al content of x = 0.62 (see Fig.  1b). A similar behaviour could be ob-served in Ti1-xAlxN (0.61 ≤ x ≤ 0.64) coatings

deposited onto cutting inserts by another CAE process done at UB = -20  V, -40  V and -80 V [47]. In that case, the estimated phase fraction of the wurtzite phase decreased from (77 ± 6) mol% for UB = -20 V and UB = -40 V to only (7 ± 3) mol% for UB = -80 V. The re-duction of the volume fraction of w-AlN at high bias voltages was also observed by Pfeiler et al. for Ti-Al-N based coatings con-taining tantalum [56] and vanadium [57]. However in our samples [39], local fluctua-tions of the Al content led to the formation of an Al-rich fcc-(Al,Ti)N phase beside the major fcc-(Ti,Al)N phase in the Ti-Al-N monolayer coatings that were deposited at high bias voltages (-80  V and -120  V). This could be concluded from a slight asymmetry of the fcc-(Ti,Al)N diffraction lines to high diffraction angles. This asymmetry was fitted by another Pearson VII function (see Fig. 1b). The asymmetry is more visible for XRD lines with even diffraction indices than for odd indices [36]. This is attributed to the differ-ent magnitude of the structure factors (and thus different diffracted intensities) for fcc- Ti1-xAlxN with low and high Al contents. For diffraction lines with even and odd indices,

Fig. 2: Influence of the bias voltage and the [Ti]/([Ti]+[Al]) ratio on the phase composition (a) and on the stress-free lattice parameter (b) of the fcc-Ti1-xAlxN phase in the Ti1-xAlxN monolayer coatings. The coloured horizontal lines in (b) indicate the stress-free lattice parameter of the fcc-Ti1-xAlxN phase as calculated from the Vegard-like dependence (Eq.(6)) for the mean composition of the coatings determined by using EPMA / WDS (x = 0.38, 0.47, 0.56 and 0.62). The stress-free lattice parameters were obtained by using the modified sin²ψ method as described in [32].

207

(7)

the structure factors of the fcc-Ti1-xAlxN phase correspond to Eq.(7) and (8), respectively:

(8)

Furthermore, the stress-free lattice parameter of the fcc-(Ti,Al)N phase of the coatings de-posited at high UB was larger than the expected lattice parameter according to the Vegard-like dependence (Eq. (6)) for the respective coat-ing composition (Fig.  2b). This implies that the whole amount of Al present in the coating is not incorporated in the fcc-(Ti,Al)N phase, but some Al is present in an Al-rich phase [32].

In the case of the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings, all three series contained fcc-(Ti,Al)N as major phase. No distinct dif-fraction lines from w-AlN appeared in the GAXRD patterns (Fig.  3). But the increased intensity between 32° and 35.2° 2θ in the GAXRD pattern measured in a laboratory ex-periment with Cu radiation suggests the for-mation of traces of wurtzitic phase as a second phase in all coatings, as shown exemplarily for the Ti-Al-N  /  Al-Ti-Ru-N multilayers of series III in Fig. 3a. The presence of traces of wurtzite phase in the coatings is better visible in the GAXRD patterns that were obtained with synchrotron radiation ( =0.10781  nm) at Rossendorf Beamline BM 20 at the Euro-pean Synchrotron Radiation Facility (ESRF) in Grenoble and which are shown in Fig. 3b.

Fig. 3: Low angle parts of the Ti-Al-N / Al-Ti-Ru-N multilayer coatings from series III deposited at different bias voltages that were obtained in a laboratory experiment using Cu radiation (a) and in a synchrotron experiment with a wavelength of λ=0.10781 nm (b). The positions of fcc-TiN, fcc-AlN and w-AlN are shown at the bottom of the figure, whereas the positions of hex-WC and fcc-TiC are given at the top of the figure. The green arrows indicate the increased intensity in the diffraction patterns caused by traces of wurtzite.

where , and are the atomicscattering factors of nitrogen, aluminium andtitanium, respectively. As ,the increasing Al content leads always to adecrease of the diffracted intensity, which ishowever stronger for diffraction lines withodd indices than for diffraction lines witheven indices. For the diffraction lines 200 and220, the ratio isnearly twice as high as compared to thediffraction line 111. For the atomic scatteringfactors calculated after the routine given byWaasmaier et al. [58], the above calculationyielded

, and

. Thus, the asymmetry is more visible for the200 and 220 diffraction lines than for the 111diffraction line.

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The increased intensity between ~22° and ~25° 2θ as well as between ~38° and ~40.5° 2θ in the GAXRD patterns (Fig. 3b) is attributed to the wurtzite phase. The coatings deposited at high UB (-60  V and -80  V) contain apparently Al-rich fcc- (Al,Ti)N as a third phase, since an asymmetry of the fcc-(Ti,Al)N diffraction lines to high diffraction angles (cf. Fig.  3) was observed like in the case of the Ti-Al-N monolayer coatings. Furthermore, the stress-free lattice parameter of the fcc-(Ti,Al)N phase, which was calculated according to the routine giv-en in Ref. [49], was substantially increased as compared to the coatings deposited at low UB (see Fig. 4).

Fig. 4: Dependence of the stress-free lattice parameter of the fcc-(Ti,Al)N phase on the bias voltages for all three Ti-Al-N / Al-Ti-(Ru)-N multilayer series.

3.2 Influence of the bias voltage and Al content on the microstructure parameters in Ti-Al-N coatings

The phase composition, which was adjused by the bias voltage and by the aluminium con-tent, influences the crystallite size and the macroscopic lattice strain. In order to deter-mine the crystallite size, the dependence of the line broadening on the diffraction vector was analysed, which is shown exemplarily for the Ti0.53Al0.47N coatings deposited at different UB in Fig.  5. The dependence found for the

fcc-(Ti,Al)N phase was similar to that one shown in Ref. [32] and suggested the presence of partially coherent nanocrystallites with small mutual misorientations [8, 59].

The size of the fcc-(Ti,Al)N nanocrystal-lites was determined according to the rou-tine proposed in Ref. [59] and is shown in Fig.  6a. At the bottom of Fig.  6a, the phase composition is displayed, which indicates the correlation of the nanocrystallites with the phase composition of the coatings. The single-phase Ti0.62Al0.38N and Ti0.53Al0.47N coatings that were deposited at low bias volt-ages (-20  V and -40  V) contained the larg-est crystallites having a size of ~ 9 - 11  nm. Substantial smaller crystallites were found in the dual phase coatings. In Ti0.44Al0.56N coatings deposited at UB = -20  V and -40  V, which contained a minor amount of w-AlN as second phase, the size of the nanocrystal-lites was ~ 6  nm. When w-AlN became the major phase in the Ti0.38Al0.62N coatings with the highest Al content that were deposited at low bias voltages, the size of the fcc-(Ti,Al)N nanocrystallites could not be determined, be-cause the diffraction lines from fcc-(Ti,Al)N were too broad and too weak. The dual-phase Ti1-xAlxN coatings with 0.38 ≤ x ≤ 0.62 and deposited at high UB (-80 V and -120 V), which contained Al-rich fcc-(Al,Ti)N as second

209

Fig. 5: Dependence of the integral breadth of XRD lines of the fcc-(Ti,Al)N phase on the modulus of the diffraction vector for the Ti0.53Al0.47N coatings deposited at different UB.

crystalline phase, were characterized by crys-tallites having a size of ~ 3.5 - 5 nm. If w-AlN was formed as third phase in these coatings the size of the crystallites became ~ 3  nm.

The macroscopic lattice strain in the fcc-(Ti,Al)N phase, which was determined ac-cording to the procedure described in Ref. [32], is given in Fig. 6b for all Ti1-xAlxN coat-ings that contained fcc-(Ti,Al)N as major phase. Assuming a Poisson’s ratio of 0.3 and a Young’s modulus of 500 GPa, the macroscopic lattice strain in Fig. 6b was transformed into the residual stress (z axis on the right-hand side) [32].

It can be seen that the bias voltage strong-ly influences the stress state of the coatings. Low tensile stress (< 0.5  GPa) and almost zero residual stress were observed in the Ti1-xAlxN coatings deposited at UB = -20  V. When the bias voltage was increased to -40 V, -80 V and -120 V the coatings were character-ized by a compressive stress state. The stress rose by a factor of ~ 4 for x ≤ 0.56 when the bias voltage was increased from -40 V to -80 V. A further increase of the bias voltage from -80  V to -120  V resulted in a minor stress increase for x ≤ 0.47 and in a saturation of the stress for x ≥ 0.56, respectively. The in-crease of the residual stress with increasing UB to -80 V is partially caused by an increasing

kinetic energy of the bombarding ions giving rise to a higher density of lattice defects, e.g., displaced atoms due to direct and recoil im-plantation of film atoms [60]. At UB = -120 V, partial defect annihilation obviously takes place as indicated by the minor increase or rather saturation of the stress as compared to UB = -80 V. According to Ref. [60], this effect is apparently facilitated by vibration and gen-eration of phonons in the vicinity of the colli-sion cascade as a result of the delivery of en-ergy of the impinging ions to the surface. This leads to short range movements of atoms and relaxation of displaced atoms. Additionally local heating of the deposited coating takes place at a high ion flux leading to an increased annealing rate of the created defects [60]. However, the high residual stress in the coat-ings deposited at high UB (-80 V and -120 V) stabilizes Al-rich fcc-(Al,Ti)N which is in ac-cordance with ab initio calculations [61].

The microstructure formation designed by the bias voltage and by the Al content influ-ences the indentation hardness of the coatings as shown in Fig. 6c. The highest hardness was found in the Ti1-xAlxN coatings with x ≤ 0.56 that were deposited at high UB and contained fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N. The high hardness in the range of 33 to 37  GPa is achieved by small nanocrystallites in the range of 3.5 to 5 nm and by a high compressive

210

Fig. 6: Influence of the bias voltage and the [Ti]/([Ti]+[Al]) ratio on the crystallite size of the fcc-(Ti,Al)N phase (a), the macroscopic lattice strain and residual stress of the fcc-(Ti,Al)N phase (b) and the indentation hardness (c). At the bottom of the figures the phase composition is illustrated, the symbols have the same meaning like in Fig. 2a.

stress in the range of 6 to 8 GPa. In the coat-ings with the highest Al content (Ti0.38Al0.62N), the hardness was reduced by 5 and 3 GPa for UB = -80  V and UB = -120  V, respectively, as compared to the Ti1-xAlxN coatings with lower Al content and the same bias voltage in each case. In the coatings with the reduced hard-ness, the smallest crystallites were found, which could facilitate increased grain bound-ary sliding [62] and thus the decrease of the hardness. Another possible explanation for the hardness reduction could be the pres-ence of the wurtzite phase exceeding a critical amount, because the volume fraction of the wurtzite phase in the Ti0.38Al0.62N coating was estimated to be (25 ± 10) vol%. A similar hard-ness evolution with different amount of molar fraction of w-AlN was observed in monolithic Ti1-xAlxN (0.61 ≤ x ≤ 0.64) coatings deposit-ed onto cutting inserts at UB = -80 V [47]. The Ti1-xAlxN coatings that contained a low amount of wurtzite phase namely (7 ± 3) mol% had a higher hardness of (31.1 ± 1) GPa than the coatings containing a high amount of wurtzite phase of (38 ± 4)  mol% which had a hard-ness of (27.7 ± 0.9) GPa.

Low compressive stress (< 2  GPa) and large crystallites (~ 9 - 11  nm) in single-phase fcc-(Ti,Al)N coatings, which were deposited at low bias voltages, resulted in low hardness of 25  GPa to 30  GPa. A hardness increase by ~ 3  GPa was observed in the Ti0.44Al0.56N coatings deposited at low bias voltages, which contained a low amount of wurtzite phase (approx. (14 ± 5) vol%) as second crystalline phase. This hardness increase can be attribut-ed to a slight increase of the compressive stress (Fig. 6b) and to the reduction of the crystallite size (Fig. 6a) as compared to the single phase coatings deposited at the same bias voltage. This observed hardness increase emphasiz-es the positive influence of low amounts of wurtzite phase in the Ti-Al-N based coat-ings with regard to their hardness. The posi-tive effect of small amounts of w-AlN on the hardness evolution was also shown for anoth-er Ti-Al-N coating series deposited by CAE [37, 55]. This effect is facilitated by partially coherent interfaces between fcc-(Ti,Al)N and

wurtzite phase, which will be discussed in the next Section (3.3). The hardness dropped to the lowest hardness values ranging between 21 and 23 GPa in the Ti0.38Al0.62N coatings de-posited at low bias voltages, which contained the wurtzite phase as major phase. This effect is caused by the predominance of the softer w-AlN phase.

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Fig. 7: (a) HRTEM of the Ti0.44Al0.56N monolayer coating deposited at UB = -40 V and local FFTs revealing the orienta-tion of fcc-(Ti,Al)N (yellow marked) and wurtzite phase (orange marked) that follows from the indices of the diffrac-tion spots in the simulated electron diffraction patterns (b).

3.3 Orientation relationships between fcc-(Ti,Al)N and wurtzite phase

Several experimentally observed orientationrelationships between partially coherent fcc-(Ti,Al)N and w-AlN are published inliterature [36, 37, 63, 64]. HRTEM done onthe Ti0.44Al0.56N coating deposited at UB = -40 V and local FFTs (Fig. 7) revealed afurther orientation relationship betweenfcc-(Ti,Al)N and w-AlN. Local FFTsperformed in the designated square areas inFig. 7a revealed fcc- (Ti,Al)N with the

direction being parallel to theincident electron beam and the wurtzitephase with the direction beingparallel to the incident electron beam. TheFFTs can be indexed with the simulatedelectron diffraction patterns given inFig. 7b. The OR between both phases ischaracterized by the planes beingparallel to the planes and the

direction being parallel to thedirection (see Fig. 8).

parallel to the planes and thedirection being parallel to the

direction (see Fig. 8).

Obviously this OR represents a preferredone, because it was found in two furthersamples namely in the Ti-Al-N / Al-Ti-Ru-Nmultilayer deposited at UB = -80 V that wasthermally treated at 950 °C and in the virginTi-Al-N / Al-Ti-N multilayer coating depo-sited at UB = -40 V [65]. This OR is charac-terized by a moderate misfit of ~ 3.6 % alongthe and direction (seeFig. 8) regarding the misfit between the Tiand Al atoms. But the misfit along the

and direction is ~ 16 %and huge. It is probably compen-sated bymicrostructure defects.

Obviously this OR represents a preferredone, because it was found in two furthersamples namely in the Ti-Al-N / Al-Ti-Ru-Nmultilayer deposited at UB = -80 V that wasthermally treated at 950 °C and in the virginTi-Al-N / Al-Ti-N multilayer coating depo-sited at UB = -40 V [65]. This OR ischaracterized by a moderate misfit of ~ 3.6 %along the and direction(see Fig. 8) regarding the misfit between theTi and Al atoms. But the misfit along the

and direction is ~ 16 % andhuge. It is probably compensated by micro-structure defects.

212

Obviously this OR represents a preferredone, because it was found in two furthersamples namely in the Ti-Al-N / Al-Ti-Ru-N multilayer deposited at UB = -80 V thatwas thermally treated at 950 °C and in thevirgin Ti-Al-N / Al-Ti-N multilayer coatingdeposited at UB = -40 V [65]. This OR ischaracterized by a moderate misfit of~ 3.6 % along the anddirection (see Fig. 8) regarding the misfitbetween the Ti and Al atoms. But the misfitalong the and direction is~ 16 % and huge. It is probablycompensated by microstructure defects.We sought to study the interface along the

and direction in moredetails by the deposition of a w-AlN filmwith a thickness of ~110 nm on top of a

orientated TiN single crystallinelayer using magnetron sputtering [65]. Thisfcc-TiN single crystalline layer was obtainedby heteroepitaxial growth on top of a

orientated MgO substrate. Furtherdetails and illustrations of the experimentaiming the heteroepitaxial growth of w-AlNon top of fcc-TiN using magnetronsputtering as deposition technology aregiven in Ref. [65]. After deposition a TEMcross-section sample was prepared with the

direction being parallel to thenormal of the TEM foil. The analysis of theORs using selected area electron diffraction(SAED) revealed two major types of AlNcolumns with a further OR [65], which wasnot published before. The observed SAEDpattern of the interface fcc-TiN / w-AlN wassimilar to the simulated one in Fig. 9. Theresults of the analysis can be summarizedbriefly as follows. Two w-AlN column typeswere found, which alternate every < 10 nmalong the flat TiN interface being parallel to

. The OR of the first w-AlN columnis characterized by and

(see red spots in Fig. 9).The second w-AlN column is created by therotation of the first w-AlN crystal (w1) by180° around the normal of the interfaceso that the direction is parallel tothe direction and theplane is parallel to the plane (seeblue spots in Fig. 9). The above mentionedOR that was observed in the Ti-Al-N basedcoatings can be suspected as a thirdorientation which appears only in a smallvolume fraction in this experiment [65].This illustrates that the material is veryflexible by the creation of partially coherentinterfaces offering a variety of orientationrelationships which could be beneficial forthe hardness development.

Obviously this OR represents a preferredone, because it was found in two furthersamples namely in the Ti-Al-N / Al-Ti-Ru-N multilayer deposited at UB = -80 V thatwas thermally treated at 950 °C and in thevirgin Ti-Al-N / Al-Ti-N multilayer coatingdeposited at UB = -40 V [65]. This OR ischaracterized by a moderate misfit of~ 3.6 % along the anddirection (see Fig. 8) regarding the misfitbetween the Ti and Al atoms. But the misfitalong the and direction is~ 16 % and huge. It is probablycompensated by microstructure defects.We sought to study the interface along the

and direction in moredetails by the deposition of a w-AlN filmwith a thickness of ~110 nm on top of a

orientated TiN single crystallinelayer using magnetron sputtering [65]. Thisfcc-TiN single crystalline layer was obtainedby heteroepitaxial growth on top of a

orientated MgO substrate. Furtherdetails and illustrations of the experimentaiming the heteroepitaxial growth of w-AlNon top of fcc-TiN using magnetronsputtering as deposition technology aregiven in Ref. [65]. After deposition a TEMcross-section sample was prepared with the

direction being parallel to thenormal of the TEM foil. The analysis of theORs using selected area electron diffraction(SAED) revealed two major types of AlNcolumns with a further OR [65], which wasnot published before. The observed SAEDpattern of the interface fcc-TiN / w-AlN wassimilar to the simulated one in Fig. 9. Theresults of the analysis can be summarizedbriefly as follows. Two w-AlN column typeswere found, which alternate every < 10 nmalong the flat TiN interface being parallel to

. The OR of the first w-AlN columnis characterized by and

(see red spots in Fig. 9).The second w-AlN column is created by therotation of the first w-AlN crystal (w1) by180° around the normal of the interfaceso that the direction is parallel tothe direction and theplane is parallel to the plane (seeblue spots in Fig. 9). The above mentionedOR that was observed in the Ti-Al-N basedcoatings can be suspected as a thirdorientation which appears only in a smallvolume fraction in this experiment [65].This illustrates that the material is veryflexible by the creation of partially coherentinterfaces offering a variety of orientationrelationships which could be beneficial forthe hardness development.

Fig. 9: Simulated electron diffractionpatterns of an OR found at theinterface of fcc-TiN (white spots)and w-AlN with two differentorientations of w-AlN (w1 = red,w2 = blue). The ,and directions areperpendicular to the plane of thepaper. The black broken line is || tothe interface and the arrowcorresponds to the normal of theinterface. The open squarescorrespond to forbidden reflectionsin the wurtzite crystallites.

Fig. 8: Schematic presentation of fcc-TiN andw-AlN in the mutual ORand . The directionsand are facing downwards. Theinteratomic distances within certain latticeplanes in fcc-TiN and w-AlN are marked bygrey dashed lines. The elementary cells arehighlighted by the broken blue lines. The Alatoms are plotted by orange, the Ti atoms byyellow and the N atoms by green spheres.Spheres that are hollow hatched representatoms lying below the plane of the paper.

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3.4 Influence of the bias voltage and the Ru content on the microstructure parameters in Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings

coatings can be suspected as a thirdorientation which appears only in a smallvolume fraction in this experiment [65].This illustrates that the material is veryflexible by the creation of partially coherentinterfaces offering a variety of orientationrelationships which could be beneficial forthe hardness development.

Another OR being important in Ti-Al-Nbased coatings is characterized by

and that isexplained in detail in Refs. [36, 55]. Thisorientation offers a good match of bothcrystal structures with a low misfit along the

and direction and alongthe and direction. Themisfit between the w-AlN phase(a = 0.311 nm, c = 0.498 nm) and the fcc-(Ti,Al)N phase depends on the latticeparameter of the fcc-(Ti,Al)N phase.Considering the interatomic distance of themetal atoms in the anddirections, it ranges between 3.6 % for fcc-TiN with a = 0.424 nm and 7 % for fcc-AlNwith a = 0.410 nm. In case of theand directions, the misfit rangesbetween 1.6 % and 4.9 %. With regard to theabove mentioned OR and to in situdiffraction experiments during annealing ofTi1-xAlxN coatings, it is suggested in Ref.[36] that a phase transformation of fcc-(Ti,Al)N to the wurtzite phase is facilitatedby the formation of stacking faults.

The stacking faults are located on theplanes in order to transform the

fcc phase to the wurtzite phase. The faultedfcc crystal is seen by XRD as a faultedwurtzite structure with the stacking faultson the planes. As a result of thefaulting, the diffraction lines 100w, 110w and200w are the most pronounced ones in thediffraction pattern, whereas the diffractionlines 101w, 102w, 103w and 202w are verybroadened due to microstructure defectslike stacking faults and the correspondingpartial dislocations [36]. The increasedlattice parameters of the wurtzite phaseindicated the incorporation of Ti into w-(Al,Ti)N, which in turn facilitated theformation of the stacking faults. This kindof phase transition would preserve thecoherence of the adjacent wurtzite and fccphase..

coatings can be suspected as a thirdorientation which appears only in a smallvolume fraction in this experiment [65].This illustrates that the material is veryflexible by the creation of partially coherentinterfaces offering a variety of orientationrelationships which could be beneficial forthe hardness development.

Another OR being important in Ti-Al-Nbased coatings is characterized by

and that isexplained in detail in Refs. [36, 55]. Thisorientation offers a good match of bothcrystal structures with a low misfit along the

and direction and alongthe and direction. Themisfit between the w-AlN phase(a = 0.311 nm, c = 0.498 nm) and the fcc-(Ti,Al)N phase depends on the latticeparameter of the fcc-(Ti,Al)N phase.Considering the interatomic distance of themetal atoms in the anddirections, it ranges between 3.6 % for fcc-TiN with a = 0.424 nm and 7 % for fcc-AlNwith a = 0.410 nm. In case of theand directions, the misfit rangesbetween 1.6 % and 4.9 %. With regard to theabove mentioned OR and to in situdiffraction experiments during annealing ofTi1-xAlxN coatings, it is suggested in Ref.[36] that a phase transformation of fcc-(Ti,Al)N to the wurtzite phase is facilitatedby the formation of stacking faults.

The stacking faults are located on theplanes in order to transform the

fcc phase to the wurtzite phase. The faultedfcc crystal is seen by XRD as a faultedwurtzite structure with the stacking faultson the planes. As a result of thefaulting, the diffraction lines 100w, 110w and200w are the most pronounced ones in thediffraction pattern, whereas the diffractionlines 101w, 102w, 103w and 202w are verybroadened due to microstructure defectslike stacking faults and the correspondingpartial dislocations [36]. The increasedlattice parameters of the wurtzite phaseindicated the incorporation of Ti into w-(Al,Ti)N, which in turn facilitated theformation of the stacking faults. This kindof phase transition would preserve thecoherence of the adjacent wurtzite and fccphase..

Information about the multilayerarchitecture of the Ti-Al-N / Al-Ti-(Ru)-Nmultilayers were obtained by analyticalscanning (S-) TEM. The layeredarchitecture of the coatings could bevisualized by dark field (DF) STEM imagesas shown in Fig. 10a and Fig. 10b for theRu-free Ti-Al-N / Al-Ti-N multilayer ofseries I and the Ti-Al-N / Al-Ti-Ru-Nmultilayer of series III, respectively. Thelength of the periodic motif estimated fromthe DF STEM images is ~ 21 nm for the Ti-Al-N / Al-Ti-N multilayer of series Ideposited at UB = -40 V and ~ 29 nm for theTi-Al-N / Al-Ti-Ru-N multilayer of seriesIII deposited at UB = -40 V.

The contrasts in the DF STEM micrographsreflect the changes in the mean atomicweight across the layers. In case of aconstant TEM foil thickness, the regionswith a high mean atomic weight appearbrighter than the areas with a lower meanatomic weight. The contrasts differ slightlyfor the coatings with Ru doping and withoutRu addition. Thin bright layers labelled as“A” are visible in the Ti-Al-N / Al-Ti-Ru-Nmultilayer (Fig. 10b) and indicate a highmean atomic weight. These layers were notvisible in the Ru-free multilayer (Fig. 10a).The X-ray spectroscopy (EDS) on the Ti Kα,Al Kα and Ru Kα lines, done in STEMacross the layer stacks with a step size of1 nm, gave information about the variationof the metal ratios (see Fig. 10c).Additionally, the measured intensity of theRu Kα line is given as blue dotted line inFig. 10c. The concentration profile given inFig. 10c revealed the gradual change in theconcentration of the metal species, which isattributed to the 2-fold rotation during CAEdeposition. The EDS analysis also showedthat the layers “A” are attributed to anenrichment of Ru. In these layers also thehighest Al ratio of ~ 0.65 was found,because Ru was deposited from the Al-richtargets. It is assumed that Ru is present in itsmetallic form because the depositiontemperature of 450 °C was much higherthan the decomposition temperature ofRuN that is approx. 100 °C [66]. The highest[Ti] / Σ[Me] ratio is located in the layerslabelled as “B” and approaches 0.56. EELSanalysis of the Ti L2,3 edge (see Fig. 10d)confirmed the fluctuation of the Ticoncentration observed by EDS.Furthermore, EELS analysis of the N K edgeverified the constant amount of nitrogen(Fig. 10d) across the individual layers.Further information about the EELSanalysis in Ti-Al-N based coatings is givenin Ref. [50].

In the Ti-Al-N / Al-Ti-Ru-N multilayer ofthe series III deposited at UB = -40 V, themedium [Ti] / Σ[Me], [Al] / Σ[Me] and[Ru] / Σ[Me] ratios averaged over threeperiodic layer stacks were (0.44 ± 0.01),(0.55 ± 0.01) and (0.01 ± 0.005),respectively. These ratios are in goodagreement with the metal ratios determinedby EPMA / WDS (see Tab. 1).

Since GAXRD indicated traces of wurtzitephase in all coatings (see Section 3.1), theformation of w-AlN in the Al-rich layerswould be possible. The highest Al ratio inthese layers is 0.65, which is above thecritical Al ratio for the formation of the fccsingle phase coatings that was found insimilarly deposited Ti1-xAlxN monolayercoatings (cf. Fig. 2a). Furthermore, this Alconcentration is close to the highestexperimentally observed solubility limit ofAl (x = 0.67) in fcc-Ti1-xAlxN [15, 16].However, HRTEM in combination withFFT across the multilayer stack revealedthat the fcc crystal structure of (Ti,Al)N isbasically maintained over the Ti-rich andAl-rich layers. Although phases with smallvolume fractions of w-AlN are difficult tobe identified by TEM, wurtzite phase in theTi-Al-N / Al-Ti-Ru-N multilayers wasfound at the column boundaries.Furthermore, the wurtzite phase waspartially coherent to the fcc-phase with thesame OR that is shown in Fig. 8, namely:

|| and || .

The macroscopic residual stress of the fcc-(Ti,Al)N phase was determined accordingto the routine proposed in Ref. [49] and isshown in Fig. 11a. For the calculation of theresidual stress, a Young’s modulus of500 GPa and Poisson’s ratio of 0.3 wereassumed. Additionally, the macroscopiclattice strain was calculated. The stressanalysis revealed that all coatings arecharacterized by a compressive stress state.The addition of Ru had no influence on thestress evolution, because the stress of allthree coating series, which were depositedat the same bias voltage, differed onlywithin the expected experimental error. Themacroscopic residual stress as well as thelattice strain is primarily influenced by thebias voltage applied during the CAEdeposition. At low bias voltage of UB = -20 V, the lowest compressive stress of~ 1 GPa was observed. The compressivestress increased nearly linearly to ~ 6.5 GPawith increasing bias voltage from -20 V to -60 V, which can be attributed to anincreasing kinetic energy of the bombardingions giving rise to a higher density of latticedefects [60]. The compressive stressincreased only slightly and saturated at~ 7 GPa, when the bias voltage wasincreased from -60 V to -80 V. The sameeffect was observed in the Ti1-xAlxNmonolayers, when the bias voltage wasincreased from -80 V to -120 V. Thesaturation of the residual stress at thehighest bias voltages was attributed to localheating of the deposited coating as a resultof intensive ion bombardment leading topartial defect annihilation e.g. relaxation ofdisplaced atoms (see Section 3.2)..

coatings can be suspected as a thirdorientation which appears only in a smallvolume fraction in this experiment [65].This illustrates that the material is veryflexible by the creation of partially coherentinterfaces offering a variety of orientationrelationships which could be beneficial forthe hardness development.

Another OR being important in Ti-Al-Nbased coatings is characterized by

and that isexplained in detail in Refs. [36, 55]. Thisorientation offers a good match of bothcrystal structures with a low misfit along the

and direction and alongthe and direction. Themisfit between the w-AlN phase(a = 0.311 nm, c = 0.498 nm) and the fcc-(Ti,Al)N phase depends on the latticeparameter of the fcc-(Ti,Al)N phase.Considering the interatomic distance of themetal atoms in the anddirections, it ranges between 3.6 % for fcc-TiN with a = 0.424 nm and 7 % for fcc-AlNwith a = 0.410 nm. In case of theand directions, the misfit rangesbetween 1.6 % and 4.9 %. With regard to theabove mentioned OR and to in situdiffraction experiments during annealing ofTi1-xAlxN coatings, it is suggested in Ref.[36] that a phase transformation of fcc-(Ti,Al)N to the wurtzite phase is facilitatedby the formation of stacking faults.

The stacking faults are located on theplanes in order to transform the

fcc phase to the wurtzite phase. The faultedfcc crystal is seen by XRD as a faultedwurtzite structure with the stacking faultson the planes. As a result of thefaulting, the diffraction lines 100w, 110w and200w are the most pronounced ones in thediffraction pattern, whereas the diffractionlines 101w, 102w, 103w and 202w are verybroadened due to microstructure defectslike stacking faults and the correspondingpartial dislocations [36]. The increasedlattice parameters of the wurtzite phaseindicated the incorporation of Ti into w-(Al,Ti)N, which in turn facilitated theformation of the stacking faults. This kindof phase transition would preserve thecoherence of the adjacent wurtzite and fccphase..

Information about the multilayer architec-ture of the Ti-Al-N / Al-Ti-(Ru)-N multi-layers were obtained by analytical scanning(S-) TEM. The layered architecture of thecoatings could be visualized by dark field(DF) STEM images as shown in Fig. 10aand Fig. 10b for the Ru-free Ti-Al-N / Al-Ti-N multilayer of series I and theTi-Al-N / Al-Ti-Ru-N multilayer of seriesIII, respectively. The length of the periodicmotif estimated from the DF STEM imagesis ~ 21 nm for the Ti-Al-N / Al-Ti-Nmultilayer of series I deposited at UB = -40 V and ~ 29 nm for the Ti-Al-N / Al-Ti-Ru-N multilayer of series III depo-sited at UB = -40 V.

214

Fig. 10: DF STEM images of the Ti-Al-N / Al-Ti-(Ru)-N multilayers of the coating series I (a) and III (b) deposited at UB = -40 V. The white line in (b) shows the position of the analytical TEM line scan. The results of EDS and EELS analysis are given in (c) and (d), respectively. The thin blue dotted line in (c) shows the intensity of the Ru Kα line.

Information about the multilayerarchitecture of the Ti-Al-N / Al-Ti-(Ru)-Nmultilayers were obtained by analyticalscanning (S-) TEM. The layeredarchitecture of the coatings could bevisualized by dark field (DF) STEM imagesas shown in Fig. 10a and Fig. 10b for theRu-free Ti-Al-N / Al-Ti-N multilayer ofseries I and the Ti-Al-N / Al-Ti-Ru-Nmultilayer of series III, respectively. Thelength of the periodic motif estimated fromthe DF STEM images is ~ 21 nm for the Ti-Al-N / Al-Ti-N multilayer of series Ideposited at UB = -40 V and ~ 29 nm for theTi-Al-N / Al-Ti-Ru-N multilayer of seriesIII deposited at UB = -40 V.

The contrasts in the DF STEM micrographsreflect the changes in the mean atomicweight across the layers. In case of aconstant TEM foil thickness, the regionswith a high mean atomic weight appearbrighter than the areas with a lower meanatomic weight. The contrasts differ slightlyfor the coatings with Ru doping and withoutRu addition. Thin bright layers labelled as“A” are visible in the Ti-Al-N / Al-Ti-Ru-Nmultilayer (Fig. 10b) and indicate a highmean atomic weight. These layers were notvisible in the Ru-free multilayer (Fig. 10a).The X-ray spectroscopy (EDS) on the Ti Kα,Al Kα and Ru Kα lines, done in STEMacross the layer stacks with a step size of1 nm, gave information about the variationof the metal ratios (see Fig. 10c).Additionally, the measured intensity of theRu Kα line is given as blue dotted line inFig. 10c. The concentration profile given inFig. 10c revealed the gradual change in theconcentration of the metal species, which isattributed to the 2-fold rotation during CAEdeposition. The EDS analysis also showedthat the layers “A” are attributed to anenrichment of Ru. In these layers also thehighest Al ratio of ~ 0.65 was found,because Ru was deposited from the Al-richtargets. It is assumed that Ru is present in itsmetallic form because the depositiontemperature of 450 °C was much higherthan the decomposition temperature ofRuN that is approx. 100 °C [66]. The highest[Ti] / Σ[Me] ratio is located in the layerslabelled as “B” and approaches 0.56. EELSanalysis of the Ti L2,3 edge (see Fig. 10d)confirmed the fluctuation of the Ticoncentration observed by EDS.Furthermore, EELS analysis of the N K edgeverified the constant amount of nitrogen(Fig. 10d) across the individual layers.Further information about the EELSanalysis in Ti-Al-N based coatings is givenin Ref. [50].

In the Ti-Al-N / Al-Ti-Ru-N multilayer ofthe series III deposited at UB = -40 V, themedium [Ti] / Σ[Me], [Al] / Σ[Me] and[Ru] / Σ[Me] ratios averaged over threeperiodic layer stacks were (0.44 ± 0.01),(0.55 ± 0.01) and (0.01 ± 0.005),respectively. These ratios are in goodagreement with the metal ratios determinedby EPMA / WDS (see Tab. 1).

Since GAXRD indicated traces of wurtzitephase in all coatings (see Section 3.1), theformation of w-AlN in the Al-rich layerswould be possible. The highest Al ratio inthese layers is 0.65, which is above thecritical Al ratio for the formation of the fccsingle phase coatings that was found insimilarly deposited Ti1-xAlxN monolayercoatings (cf. Fig. 2a). Furthermore, this Alconcentration is close to the highestexperimentally observed solubility limit ofAl (x = 0.67) in fcc-Ti1-xAlxN [15, 16].However, HRTEM in combination withFFT across the multilayer stack revealedthat the fcc crystal structure of (Ti,Al)N isbasically maintained over the Ti-rich andAl-rich layers. Although phases with smallvolume fractions of w-AlN are difficult tobe identified by TEM, wurtzite phase in theTi-Al-N / Al-Ti-Ru-N multilayers wasfound at the column boundaries.Furthermore, the wurtzite phase waspartially coherent to the fcc-phase with thesame OR that is shown in Fig. 8, namely:

|| and || .

The macroscopic residual stress of the fcc-(Ti,Al)N phase was determined accordingto the routine proposed in Ref. [49] and isshown in Fig. 11a. For the calculation of theresidual stress, a Young’s modulus of500 GPa and Poisson’s ratio of 0.3 wereassumed. Additionally, the macroscopiclattice strain was calculated. The stressanalysis revealed that all coatings arecharacterized by a compressive stress state.The addition of Ru had no influence on thestress evolution, because the stress of allthree coating series, which were depositedat the same bias voltage, differed onlywithin the expected experimental error. Themacroscopic residual stress as well as thelattice strain is primarily influenced by thebias voltage applied during the CAEdeposition. At low bias voltage of UB = -20 V, the lowest compressive stress of~ 1 GPa was observed. The compressivestress increased nearly linearly to ~ 6.5 GPawith increasing bias voltage from -20 V to -60 V, which can be attributed to anincreasing kinetic energy of the bombardingions giving rise to a higher density of latticedefects [60]. The compressive stressincreased only slightly and saturated at~ 7 GPa, when the bias voltage wasincreased from -60 V to -80 V. The sameeffect was observed in the Ti1-xAlxNmonolayers, when the bias voltage wasincreased from -80 V to -120 V. Thesaturation of the residual stress at thehighest bias voltages was attributed to localheating of the deposited coating as a resultof intensive ion bombardment leading topartial defect annihilation e.g. relaxation ofdisplaced atoms (see Section 3.2)..

FFT across the multilayer stack revealed thatthe fcc crystal structure of (Ti,Al)N isbasically maintained over the Ti-rich andAl-rich layers. Although phases with smallvolume fractions of w-AlN are difficult to beidentified by TEM, wurtzite phase in theTi-Al-N / Al-Ti-Ru-N multilayers was foundat the column boundaries. Furthermore, thewurtzite phase was partially coherent to thefcc-phase with the same OR that is shown inFig. 8, namely: || and

|| .

215

Fig. 11: Influence of the bias voltage and the Ru addition on the residual stress as well as macroscopic lattice strain (a) and indentation hardness (b).

The macroscopic residual stress of the fcc-(Ti,Al)N phase was determined according tothe routine proposed in Ref. [49] and isshown in Fig. 11a. For the calculation of theresidual stress, a Young’s modulus of500 GPa and Poisson’s ratio of 0.3 wereassumed. Additionally, the macroscopiclattice strain was calculated. The stressanalysis revealed that all coatings arecharacterized by a compressive stress state.The addition of Ru had no influence on thestress evolution, because the stress of allthree coating series, which were deposited atthe same bias voltage, differed only withinthe expected experimental error. Themacroscopic residual stress as well as thelattice strain is primarily influenced by thebias voltage applied during the CAEdeposition. At low bias voltage of UB = -20 V,the lowest compressive stress of ~ 1 GPa wasobserved. The compressive stress increasednearly linearly to ~ 6.5 GPa with increasingbias voltage from -20 V to -60 V, which canbe attributed to an increasing kinetic energyof the bombarding ions giving rise to ahigher density of lattice defects [60].

The compressive stress increased onlyslightly and saturated at ~ 7 GPa, when thebias voltage was increased from -60 V to-80 V. The same effect was observed in theTi1-xAlxN monolayers, when the bias voltagewas increased from -80 V to -120 V. Thesaturation of the residual stress at the highestbias voltages was attributed to local heatingof the deposited coating as a result ofintensive ion bombardment leading topartial defect annihilation e.g. relaxation ofdisplaced atoms (see Section 3.2).

The indentation hardness of the Ti-Al-N/ Al-Ti-(Ru)-N multilayers is given in Fig. 11b.The hardness evolution vs. UB can beclassified into two groups. The first groupcomprises the coatings deposited at low biasvoltages (-20 V and -40 V). Their hardnesslies in the range between 30 and 32 GPa. Thesecond group contains the coatingsdeposited at high bias voltage (-60 V and-80 V). This group is characterized by thehighest hardness values lying between 33 and35 GPa. The increasing hardness with risingbias voltage indicates the contribution of thehigher compressive stress to the hardnessenhancement in the coatings. It seems thatthe Ru addition had no remarkable effect onthe hardness. The indentation hardnessdiffers for different Ru additions within therange of errors.

The macroscopic residual stress of the fcc-(Ti,Al)N phase was determined according tothe routine proposed in Ref. [49] and isshown in Fig. 11a. For the calculation of theresidual stress, a Young’s modulus of500 GPa and Poisson’s ratio of 0.3 wereassumed. Additionally, the macroscopiclattice strain was calculated. The stressanalysis revealed that all coatings arecharacterized by a compressive stress state.The addition of Ru had no influence on thestress evolution, because the stress of allthree coating series, which were deposited atthe same bias voltage, differed only withinthe expected experimental error. Themacroscopic residual stress as well as thelattice strain is primarily influenced by thebias voltage applied during the CAEdeposition. At low bias voltage of UB = -20 V,the lowest compressive stress of ~ 1 GPa wasobserved. The compressive stress increasednearly linearly to ~ 6.5 GPa with increasingbias voltage from -20 V to -60 V, which canbe attributed to an increasing kinetic energyof the bombarding ions giving rise to ahigher density of lattice defects [60].

The compressive stress increased onlyslightly and saturated at ~ 7 GPa, when thebias voltage was increased from -60 V to-80 V. The same effect was observed in theTi1-xAlxN monolayers, when the bias voltagewas increased from -80 V to -120 V. Thesaturation of the residual stress at the highestbias voltages was attributed to local heatingof the deposited coating as a result ofintensive ion bombardment leading topartial defect annihilation e.g. relaxation ofdisplaced atoms (see Section 3.2).

The indentation hardness of the Ti-Al-N/ Al-Ti-(Ru)-N multilayers is given in Fig. 11b.The hardness evolution vs. UB can beclassified into two groups. The first groupcomprises the coatings deposited at low biasvoltages (-20 V and -40 V). Their hardnesslies in the range between 30 and 32 GPa. Thesecond group contains the coatingsdeposited at high bias voltage (-60 V and-80 V). This group is characterized by thehighest hardness values lying between 33 and35 GPa. The increasing hardness with risingbias voltage indicates the contribution of thehigher compressive stress to the hardnessenhancement in the coatings. It seems thatthe Ru addition had no remarkable effect onthe hardness. The indentation hardnessdiffers for different Ru additions within therange of errors.

216

Fig. 12: Summary of the phase composition in the as-deposited coatings (a) and after annealing at 850 °C as estimated from Rietveld analysis (b). The lines are guide for the eyes.

3.5 Thermal stability of Ti-Al-N based coatings

The thermal induced microstructure chang-es were studied by in situ high temperature GAXRD experiments at the Rossendorf Beamline BM 20 at the ESRF and are de-scribed in detail in terms of the stress-free lattice parameter, microstrain, macroscopic lattice strain and phase composition dur-ing annealing of Ti1-xAlxN coatings in Refs. [32, 33, 36]. The results of these analyses re-vealed that the bias voltage applied during the deposition influences also the thermal stability of the coatings. During the annealing at 850 °C, the Ti1-xAlxN coatings with 0.38 ≤ x ≤ 0.56 decomposed into Al deplet-ed fcc-(Ti,Al)N, fcc-AlN and wurtzite phase. It was found that the bias voltage controls the relative amount of fcc-AlN and wurtzite phase, whereas the Al content determines the amount of fcc-(Ti,Al)N. The effect of the bias voltage and the chemical composition on the development of the phase composition of the Ti1-xAlxN coatings in course of the thermal treatment up to 850 °C is presented in a com-pact form in Fig. 12. The changes are illus-

trated by the phase composition in the as-de-posited state (Fig. 12a) and after annealing at 850 °C (Fig. 12b) for three different coating compositions (x = 0.38, 0.47, 0.56) and three bias voltages (UB = -40 V, -80 V and -120 V).

As discussed in Section 3.1, the bias volt-age influences the phase composition in the as-deposited (ad) state. At high UB (-80 V and -120 V), the samples contained fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N, whereas the sam-ples deposited at low UB (-40 V) contained fcc-(Ti,Al)N as single-phase for x ≤ 0.47 (Fig. 12a). During annealing at 850 °C, the formation of Al-rich fcc-(Al,Ti)N and wurt-zite phase were observed simultaneously in the coatings. However, the relative amount of wurtzite phase increased with increasing bias voltage (see Fig. 12b). The Ti0.44Al0.56N coating deposited at UB = -40 V deviated from this trend, since wurtzite phase was present al-ready in the as-deposited sample. Hence, only traces of fcc-AlN formed during annealing. The increase of the wurtzite phase fraction and the decrease of fcc-AlN with increasing UB after annealing at 850 °C imply an accel-erated decomposition process in the coatings

217

Fig. 13: Microstructure evolution of the Ti-Al-N / Al-Ti-(Ru)-N multilayers of the three coating series (I = circle, II = triangle, III = square) deposited at UB = -40 V (left column) and UB = -80 V (right column) in course of the thermal treatment showing the stress-free lattice parameter (a,b), macroscopic lattice strain (c,d) and indentation hardness (e,f). The filled and open symbols in (a-d) represent the microstructure parameters of the fcc-(Ti,Al)N and the nearly Al-free fcc-(Ti,Al)N, respectively.

deposited at high bias voltages. It was sug-gested that a retarded reduction of the mac-roscopic lattice strain during annealing of the coatings deposited at UB = -120 V accelerates the formation of w-AlN [33].

A similar influence of the bias voltage on the development of the phase composition was observed in the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings. Each multilayer coating was annealed consequently at 450 °C, 650 °C, 850 °C and 950 °C for 60 min in argon at-mosphere. After each annealing step, ex situ GAXRD was done and the diffraction pat-terns were analysed with respect to the stress-free lattice parameter and macroscopic lattice strain ε of the fcc phases. Additionally, the indentation hardness was measured after each annealing step. The changes in course of the thermal treatment are shown exemplarily for the coatings with three different Ru addi-tions that were deposited at UB = -40 V and

UB = -80 V in Fig. 13. Comparing the trends of the evolution of (Figs. 13a and 13b), ε (Figs. 13c and 13d) and indentation hardness (Figs. 13e and 13f) in the coatings that were deposited at the same bias voltage, no distinct influence of the Ru addition on their develop-ment during annealing in argon atmosphere could be revealed. Nevertheless, a beneficial effect of the Ru concentration on the thermal stability or hardness is possible during the ap-plication of the Ti-Al-N / Al-Ti-(Ru)-N mul-tilayer coatings in cutting processes as report-ed in [13]. The reason is that the conditions of the annealing experiment differed from the conditions prevailing during cutting. During cutting, the coatings are exposed to high tem-peratures for less than 30 min [13], but the de-composition of metastable phases takes place simultaneously with the oxidation of the sam-ples since cutting is usually done in air.

218

However, the influence of the bias voltage on the thermal stability of the Ti-Al-N / Al-Ti-(Ru)-N multilayer during annealing in argon atmosphere was evident and is shown for the coatings deposited at bias voltages of -40 V and -80 V because both bias voltages represent the general changes taking place in the coatings deposited at low and at high UB, respectively. The first obvious microstruc-ture changes became apparent after annealing at 650 °C. The stress-free lattice parameter and the macroscopic lattice strain of the fcc-(Ti,Al)N decreased. This can be attributed to defect annihilation and the incorporation of Al into fcc-(Ti,Al)N. In the coatings depos-ited at UB = -80 V, the relative reduction of ε and was more pronounced than in the coating deposited at UB = -40 V. For exam-ple after annealing at 650 °C, of the mul-tilayer coating of series III deposited at UB = -80 V decreased to 0.4178 nm (see Fig. 13b). According to the Vegard-like dependence from Eq. (6), this corresponds to an Al con-tent of x = 0.45. However, after annealing at 650 °C, and ε are still higher in the coatings deposited at UB = -80 V than in the coatings deposited at UB = -40 V (see Figs. 13a to 13d).

In all coatings, the onset of the decomposi-tion was observed after annealing at 850 °C. The Ti-rich and Al-rich fcc-(Ti,Al)N formed and the fraction of wurtzite phase increased slightly. These three phases existed beside the major fcc-(Ti,Al)N phase. This coexistence of the phases is exemplarily shown for the Ti-Al-N / Al-Ti-Ru-N multilayer coatings of series III deposited at UB = -40 V and UB = -80 V in Fig. 14. For the coating de-posited at UB = -40 V, the stress-free lattice parameter of the Ti-rich fcc-(Ti,Al)N phase (see open symbols in Fig. 13a) indicated the incorporation of x = 0.16 aluminium in the Ti-rich fcc-(Ti,Al)N. The Ti-rich (Ti,Al)N was under compressive stress (see open sym-bols in Fig. 13c). In the coatings deposited at high bias voltages, the wurtzite phase became more apparent as compared to the coatings deposited at low bias voltages after annealing at 850 °C (cf. Fig. 14).

After annealing at 950 °C, the fcc-(Ti,Al)N phase decomposed nearly completely into wurtzite phase, Ti-rich fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N (see Fig.14). The micro-structure analysis revealed that the decom-position process after annealing at 950 °C proceeded further in the coatings, which were deposited at high UB, than in the coat-ings deposited at low UB. This was concluded from (i) of the Ti-rich fcc-(Ti,Al)N phase, (ii) of the Al-rich fcc-(Al,Ti)N phase, (iii) the intensity of the diffraction lines of the wurtzite phase in the GAXRD patterns (Fig. 14) and (iv) the hardness reduction. The indi-vidual items can be explained with the aid of Figs. 13 and 14 as follows.

(i) of the Ti-rich fcc-(Ti,Al)N phase is higher for the coatings deposited at UB = -80 V than for UB = -40 V (see open symbols in Figs. 13a and13b), which indicates a lower Al incorporation in fcc-(Ti,Al)N at high UB than at low UB. For instance, the Al incorpo-ration in the Ti-rich fcc-(Ti,Al)N phase of the multilayer coatings of series III is x = 0.15 for UB = -40 V and x = 0.09 for UB = -80 V, as cal-culated by using Eq. (6).

(ii) of the Al-rich fcc-(Al,Ti)N phase is lo-wer for coatings deposited at UB = -80 V than for UB = -40 V (see filled symbols in Figs. 13a and 13b), which signifies a lower Ti concen-tration in the Al-rich fcc phase at high UB. Both observations denote that the separation of titanium nitride and aluminium nitride progressed to a further extent in the coatings, which were deposited at high bias voltages.

After annealing at 950 °C, the fcc-(Ti,Al)Nphase decomposed nearly completely intowurtzite phase, Ti-rich fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N (see Fig.14). Themicrostructure analysis revealed that thedecomposition process after annealing at950 °C proceeded further in the coatings,which were deposited at high UB, than inthe coatings deposited at low UB. This wasconcluded from (i) of the Ti-rich fcc-(Ti,Al)N phase, (ii) of the Al-rich fcc-(Al,Ti)N phase, (iii) the intensity of thediffraction lines of the wurtzite phase in theGAXRD patterns (Fig. 14) and (iv) thehardness reduction. The individual itemscan be explained with the aid of Figs. 13 and14 as follows.

(i) of the Ti-rich fcc-(Ti,Al)N phase ishigher for the coatings deposited at UB = -80 V than for UB = -40 V (see open symbolsin Figs. 13a and13b), which indicates alower Al incorporation in fcc-(Ti,Al)N athigh UB than at low UB. For instance, the Alincorporation in the Ti-rich fcc-(Ti,Al)Nphase of the multilayer coatings of series IIIis x = 0.15 for UB = -40 V and x = 0.09 forUB = -80 V, as calculated by using Eq. (6).

(ii) of the Al-rich fcc-(Al,Ti)N phase islower for coatings deposited at UB = -80 Vthan for UB = -40 V (see filled symbols inFigs. 13a and 13b), which signifies a lowerTi concentration in the Al-rich fcc phase athigh UB. Both observations denote that theseparation of titanium nitride andaluminium nitride progressed to a furtherextent in the coatings, which were depositedat high bias voltages.(iii) The diffraction lines of the wurtzitephase after annealing of the sampledeposited at UB = -80 V possess higherintensity than the diffraction lines observedin the GAXRD pattern for the thermallytreated sample deposited at UB = -40 V (cf.Fig. 14).

(iv) The indentation hardness of themultilayer coatings deposited at UB = -80 Vdropped from approx. 33 GPa to approx.27 GPa (Fig. 13f), whereas after annealing at950 °C the hardness of the coatingsdeposited at UB = -40 V decreased onlyslightly below the hardness level of the as-deposited state, namely to approx. 30 GPa(Fig. 13e). The obvious hardness reductionafter 950 °C in the coatings deposited atUB = -80 V can be attributed to a higherfraction of the wurtzite phase than in thecoatings deposited at UB = -40 V, which wasindicated also by the higher intensity of thediffraction lines of the wurtzite phase. Thehigher fraction of the wurtzite phase inannealed Ti-Al-N / Al-Ti-(Ru)-N multilayercoatings deposited at high UB coincides withthe results found for the Ti1-xAlxN coatings(see first part of Section 3.5) and confirmsthe effect of the bias voltage on the thermalstability of Ti-Al-N based coatings.

The microstructure analysis of the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings produced by CAEelucidated the role of the energetic particlebombardment during CAE deposition. Itcould be shown that the bias voltage is avaluable deposition parameter in order todesign the microstructure, properties andthermal stability of Ti-Al-N based coatings.

For the Ti-Al-N monolayer coatings, it wasshown that the phase composition, the sizeof the fcc-(Ti,Al)N nanocrystallites as wellas the macroscopic lattice strain andconsequently the hardness of the coatingscan be adjusted by the Al content and thebias voltage. Low bias voltages resulted in auniform distribution of Ti and Al, whereashigh bias voltages caused local fluctuationsof Al and Ti in (Ti,Al)N and a rising defectdensity that facilitated the formation of anAl-rich fcc-(Al,Ti)N phase in monolayerand multilayer coatings at high UB (-60 V, -80 V and -120 V). Furthermore, a high UBlimited the volume fraction of wurtzitephase in the as-deposited Ti1-xAlxN coatingsfor x ≥ 0.56 as compared to low UB. It couldbe shown that the presence of low amountsof wurtzite phase (≤ 15 vol%) is notgenerally detrimental for the hardness. Onthe contrary, partially coherent interfacesbetween wurtzite and fcc phase arebeneficial for the coatings’ hardness. Severalorientation relationships between wurtziteand fcc phase offering partial coherentinterfaces could be identified by local FFTsof HRTEM images.

219

Fig. 14: Low-angle parts of GAXRD patterns of the Ti-Al-N / Al-Ti-Ru-N multilayer coatings of series III deposited at UB = -40 V (a) and UB = -80 V (b) after annealing at 850 °C and 950 °C. The measured data are shown as blue points and the whole fit as red lines. The grey lines show the individual diffraction peaks as fitted by Pearson VII functions. The positions of the fcc-TiN, fcc-AlN and w-AlN diffraction lines are indicated at the bottom of the figures. The positions of fcc-TiC and hex-WC coming from the substrate are shown at the top.

4. Conclusions

The microstructure analysis of the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-N multi-layer coatings produced by CAE elucidatedthe role of the energetic particle bombar-dment during CAE deposition. It could beshown that the bias voltage is a valuabledeposition parameter in order to design themicrostructure, properties and thermalstability of Ti-Al-N based coatings.

For the Ti-Al-N monolayer coatings, it wasshown that the phase composition, the sizeof the fcc-(Ti,Al)N nanocrystallites as wellas the macroscopic lattice strain andconsequently the hardness of the coatingscan be adjusted by the Al content and thebias voltage. Low bias voltages resulted in auniform distribution of Ti and Al, whereashigh bias voltages caused local fluctuationsof Al and Ti in (Ti,Al)N and a rising defectdensity that facilitated the formation of anAl-rich fcc-(Al,Ti)N phase in monolayerand multilayer coatings at high UB (-60 V,-80 V and -120 V). Furthermore, a high UBlimited the volume fraction of wurtzitephase in the as-deposited Ti1-xAlxN coatingsfor x ≥ 0.56 as compared to low UB. It couldbe shown that the presence of low amounts

After annealing at 950 °C, the fcc-(Ti,Al)Nphase decomposed nearly completely intowurtzite phase, Ti-rich fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N (see Fig.14). Themicrostructure analysis revealed that thedecomposition process after annealing at950 °C proceeded further in the coatings,which were deposited at high UB, than inthe coatings deposited at low UB. This wasconcluded from (i) of the Ti-rich fcc-(Ti,Al)N phase, (ii) of the Al-rich fcc-(Al,Ti)N phase, (iii) the intensity of thediffraction lines of the wurtzite phase in theGAXRD patterns (Fig. 14) and (iv) thehardness reduction. The individual itemscan be explained with the aid of Figs. 13 and14 as follows.

(i) of the Ti-rich fcc-(Ti,Al)N phase ishigher for the coatings deposited at UB = -80 V than for UB = -40 V (see open symbolsin Figs. 13a and13b), which indicates alower Al incorporation in fcc-(Ti,Al)N athigh UB than at low UB. For instance, the Alincorporation in the Ti-rich fcc-(Ti,Al)Nphase of the multilayer coatings of series IIIis x = 0.15 for UB = -40 V and x = 0.09 forUB = -80 V, as calculated by using Eq. (6).

(ii) of the Al-rich fcc-(Al,Ti)N phase islower for coatings deposited at UB = -80 Vthan for UB = -40 V (see filled symbols inFigs. 13a and 13b), which signifies a lowerTi concentration in the Al-rich fcc phase athigh UB. Both observations denote that theseparation of titanium nitride andaluminium nitride progressed to a furtherextent in the coatings, which were depositedat high bias voltages.(iii) The diffraction lines of the wurtzitephase after annealing of the sampledeposited at UB = -80 V possess higherintensity than the diffraction lines observedin the GAXRD pattern for the thermallytreated sample deposited at UB = -40 V (cf.Fig. 14).

(iv) The indentation hardness of themultilayer coatings deposited at UB = -80 Vdropped from approx. 33 GPa to approx.27 GPa (Fig. 13f), whereas after annealing at950 °C the hardness of the coatingsdeposited at UB = -40 V decreased onlyslightly below the hardness level of the as-deposited state, namely to approx. 30 GPa(Fig. 13e). The obvious hardness reductionafter 950 °C in the coatings deposited atUB = -80 V can be attributed to a higherfraction of the wurtzite phase than in thecoatings deposited at UB = -40 V, which wasindicated also by the higher intensity of thediffraction lines of the wurtzite phase. Thehigher fraction of the wurtzite phase inannealed Ti-Al-N / Al-Ti-(Ru)-N multilayercoatings deposited at high UB coincides withthe results found for the Ti1-xAlxN coatings(see first part of Section 3.5) and confirmsthe effect of the bias voltage on the thermalstability of Ti-Al-N based coatings.

The microstructure analysis of the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings produced by CAEelucidated the role of the energetic particlebombardment during CAE deposition. Itcould be shown that the bias voltage is avaluable deposition parameter in order todesign the microstructure, properties andthermal stability of Ti-Al-N based coatings.

For the Ti-Al-N monolayer coatings, it wasshown that the phase composition, the sizeof the fcc-(Ti,Al)N nanocrystallites as wellas the macroscopic lattice strain andconsequently the hardness of the coatingscan be adjusted by the Al content and thebias voltage. Low bias voltages resulted in auniform distribution of Ti and Al, whereashigh bias voltages caused local fluctuationsof Al and Ti in (Ti,Al)N and a rising defectdensity that facilitated the formation of anAl-rich fcc-(Al,Ti)N phase in monolayerand multilayer coatings at high UB (-60 V, -80 V and -120 V). Furthermore, a high UBlimited the volume fraction of wurtzitephase in the as-deposited Ti1-xAlxN coatingsfor x ≥ 0.56 as compared to low UB. It couldbe shown that the presence of low amountsof wurtzite phase (≤ 15 vol%) is notgenerally detrimental for the hardness. Onthe contrary, partially coherent interfacesbetween wurtzite and fcc phase arebeneficial for the coatings’ hardness. Severalorientation relationships between wurtziteand fcc phase offering partial coherentinterfaces could be identified by local FFTsof HRTEM images.

220

After annealing at 950 °C, the fcc-(Ti,Al)Nphase decomposed nearly completely intowurtzite phase, Ti-rich fcc-(Ti,Al)N and Al-rich fcc-(Al,Ti)N (see Fig.14). Themicrostructure analysis revealed that thedecomposition process after annealing at950 °C proceeded further in the coatings,which were deposited at high UB, than inthe coatings deposited at low UB. This wasconcluded from (i) of the Ti-rich fcc-(Ti,Al)N phase, (ii) of the Al-rich fcc-(Al,Ti)N phase, (iii) the intensity of thediffraction lines of the wurtzite phase in theGAXRD patterns (Fig. 14) and (iv) thehardness reduction. The individual itemscan be explained with the aid of Figs. 13 and14 as follows.

(i) of the Ti-rich fcc-(Ti,Al)N phase ishigher for the coatings deposited at UB = -80 V than for UB = -40 V (see open symbolsin Figs. 13a and13b), which indicates alower Al incorporation in fcc-(Ti,Al)N athigh UB than at low UB. For instance, the Alincorporation in the Ti-rich fcc-(Ti,Al)Nphase of the multilayer coatings of series IIIis x = 0.15 for UB = -40 V and x = 0.09 forUB = -80 V, as calculated by using Eq. (6).

(ii) of the Al-rich fcc-(Al,Ti)N phase islower for coatings deposited at UB = -80 Vthan for UB = -40 V (see filled symbols inFigs. 13a and 13b), which signifies a lowerTi concentration in the Al-rich fcc phase athigh UB. Both observations denote that theseparation of titanium nitride andaluminium nitride progressed to a furtherextent in the coatings, which were depositedat high bias voltages.(iii) The diffraction lines of the wurtzitephase after annealing of the sampledeposited at UB = -80 V possess higherintensity than the diffraction lines observedin the GAXRD pattern for the thermallytreated sample deposited at UB = -40 V (cf.Fig. 14).

(iv) The indentation hardness of themultilayer coatings deposited at UB = -80 Vdropped from approx. 33 GPa to approx.27 GPa (Fig. 13f), whereas after annealing at950 °C the hardness of the coatingsdeposited at UB = -40 V decreased onlyslightly below the hardness level of the as-deposited state, namely to approx. 30 GPa(Fig. 13e). The obvious hardness reductionafter 950 °C in the coatings deposited atUB = -80 V can be attributed to a higherfraction of the wurtzite phase than in thecoatings deposited at UB = -40 V, which wasindicated also by the higher intensity of thediffraction lines of the wurtzite phase. Thehigher fraction of the wurtzite phase inannealed Ti-Al-N / Al-Ti-(Ru)-N multilayercoatings deposited at high UB coincides withthe results found for the Ti1-xAlxN coatings(see first part of Section 3.5) and confirmsthe effect of the bias voltage on the thermalstability of Ti-Al-N based coatings.

The microstructure analysis of the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings produced by CAEelucidated the role of the energetic particlebombardment during CAE deposition. Itcould be shown that the bias voltage is avaluable deposition parameter in order todesign the microstructure, properties andthermal stability of Ti-Al-N based coatings.

For the Ti-Al-N monolayer coatings, it wasshown that the phase composition, the sizeof the fcc-(Ti,Al)N nanocrystallites as wellas the macroscopic lattice strain andconsequently the hardness of the coatingscan be adjusted by the Al content and thebias voltage. Low bias voltages resulted in auniform distribution of Ti and Al, whereashigh bias voltages caused local fluctuationsof Al and Ti in (Ti,Al)N and a rising defectdensity that facilitated the formation of anAl-rich fcc-(Al,Ti)N phase in monolayerand multilayer coatings at high UB (-60 V, -80 V and -120 V). Furthermore, a high UBlimited the volume fraction of wurtzitephase in the as-deposited Ti1-xAlxN coatingsfor x ≥ 0.56 as compared to low UB. It couldbe shown that the presence of low amountsof wurtzite phase (≤ 15 vol%) is notgenerally detrimental for the hardness. Onthe contrary, partially coherent interfacesbetween wurtzite and fcc phase arebeneficial for the coatings’ hardness. Severalorientation relationships between wurtziteand fcc phase offering partial coherentinterfaces could be identified by local FFTsof HRTEM images.

The bias voltage applied during thedeposition of the coatings influenced alsotheir thermal stability during annealing,since the decomposition in the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings deposited at low UB wasretarded as compared to the coatingsdeposited at high UB. Furthermore, theannealing of the Ti-Al-N based coatings invacuum and argon atmosphere led to theformation of Ti-rich fcc-(Ti,Al)N as well aswurtzite phase and its coexistence with fcc-AlN in the temperature range of 850 - 950 °C. This observation supports the ideathat a direct transformation of fcc-(Ti,Al)Nto wurtzite phase is possible.

The addition of Ru to the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings had noremarkable effect on the microstructure inthe as-deposited state and during annealingin argon atmosphere.

According to the results shown above, Ti-Al-N / Al-Ti-(Ru)-N multilayer coatingswith a medium [ ratio of ~ 0.53and Ti-Al-N coatings with an average[ ratio of ~ 0.47 that weredeposited at low UB (e.g. UB = -40 V) can berecommended for the application atelevated temperatures up to 950°C, since theformation of the large volume fractions ofwurtzite phase is retarded and the hardnessis maintained after annealing. The coatingswith similar [ ratios that weredeposited at high UB (e.g. UB = -80 V) couldbe favoured for the application attemperatures below <850°C, since thehardness is high in the as deposited state butthe generation of wurtzite phase exceeding acritical volume fraction is acceleratedduring annealing leading to a hardnessreduction.

This work was performed within the Clusterof Excellence “Structure Design of NovelHigh-Performance Materials via AtomicDesign and Defect Engineering (ADDE)”that is financially supported by theEuropean Union (European Fund ofRegional Development) and by the Ministryof Science and Art of Saxony (SMWK). Weacknowledge Helmholtz-Zentrum Dresden-Rossendorf for providing us with the in-house beam time at ROBL and thankCarsten Baehtz for assistance during themeasurements on the beamline BM20.

The bias voltage applied during thedeposition of the coatings influenced alsotheir thermal stability during annealing,since the decomposition in the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings deposited at low UB wasretarded as compared to the coatingsdeposited at high UB. Furthermore, theannealing of the Ti-Al-N based coatings invacuum and argon atmosphere led to theformation of Ti-rich fcc-(Ti,Al)N as well aswurtzite phase and its coexistence with fcc-AlN in the temperature range of 850 - 950 °C. This observation supports the ideathat a direct transformation of fcc-(Ti,Al)Nto wurtzite phase is possible.

The addition of Ru to the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings had noremarkable effect on the microstructure inthe as-deposited state and during annealingin argon atmosphere.

According to the results shown above, Ti-Al-N / Al-Ti-(Ru)-N multilayer coatingswith a medium [ ratio of ~ 0.53and Ti-Al-N coatings with an average[ ratio of ~ 0.47 that weredeposited at low UB (e.g. UB = -40 V) can berecommended for the application atelevated temperatures up to 950°C, since theformation of the large volume fractions ofwurtzite phase is retarded and the hardnessis maintained after annealing. The coatingswith similar [ ratios that weredeposited at high UB (e.g. UB = -80 V) couldbe favoured for the application attemperatures below <850°C, since thehardness is high in the as deposited state butthe generation of wurtzite phase exceeding acritical volume fraction is acceleratedduring annealing leading to a hardnessreduction.

This work was performed within the Clusterof Excellence “Structure Design of NovelHigh-Performance Materials via AtomicDesign and Defect Engineering (ADDE)”that is financially supported by theEuropean Union (European Fund ofRegional Development) and by the Ministryof Science and Art of Saxony (SMWK). Weacknowledge Helmholtz-Zentrum Dresden-Rossendorf for providing us with the in-house beam time at ROBL and thankCarsten Baehtz for assistance during themeasurements on the beamline BM20.

Acknowledgement

The bias voltage applied during thedeposition of the coatings influenced alsotheir thermal stability during annealing,since the decomposition in the Ti-Al-Nmonolayer and Ti-Al-N / Al-Ti-(Ru)-Nmultilayer coatings deposited at low UB wasretarded as compared to the coatingsdeposited at high UB. Furthermore, theannealing of the Ti-Al-N based coatings invacuum and argon atmosphere led to theformation of Ti-rich fcc-(Ti,Al)N as well aswurtzite phase and its coexistence with fcc-AlN in the temperature range of 850 - 950 °C. This observation supports the ideathat a direct transformation of fcc-(Ti,Al)Nto wurtzite phase is possible.

The addition of Ru to the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings had noremarkable effect on the microstructure inthe as-deposited state and during annealingin argon atmosphere.

According to the results shown above, Ti-Al-N / Al-Ti-(Ru)-N multilayer coatingswith a medium [ ratio of ~ 0.53and Ti-Al-N coatings with an average[ ratio of ~ 0.47 that weredeposited at low UB (e.g. UB = -40 V) can berecommended for the application atelevated temperatures up to 950°C, since theformation of the large volume fractions ofwurtzite phase is retarded and the hardnessis maintained after annealing. The coatingswith similar [ ratios that weredeposited at high UB (e.g. UB = -80 V) couldbe favoured for the application attemperatures below <850°C, since thehardness is high in the as deposited state butthe generation of wurtzite phase exceeding acritical volume fraction is acceleratedduring annealing leading to a hardnessreduction.

This work was performed within the Clusterof Excellence “Structure Design of NovelHigh-Performance Materials via AtomicDesign and Defect Engineering (ADDE)”that is financially supported by theEuropean Union (European Fund ofRegional Development) and by the Ministryof Science and Art of Saxony (SMWK). Weacknowledge Helmholtz-Zentrum Dresden-Rossendorf for providing us with the in-house beam time at ROBL and thankCarsten Baehtz for assistance during themeasurements on the beamline BM20.

According to the results shown above,Ti-Al-N / Al-Ti-(Ru)-N multilayer coatingswith a medium [ ratio of ~ 0.53and Ti-Al-N coatings with an average[ ratio of ~ 0.47 that weredeposited at low UB (e.g. UB = -40 V) can berecommended for the application at eleva-ted temperatures up to 950 °C, since theformation of the large volume fractions ofwurtzite phase is retarded and the hardnessis maintained after annealing. The coatingswith similar [ ratios that weredeposited at high UB (e.g. UB = -80 V) couldbe favoured for the application at tempe-ratures below <850 °C, since the hardnessis high in the as deposited state butthe generation of wurtzite phase exceeding a

The addition of Ru to the Ti-Al-N / Al-Ti-(Ru)-N multilayer coatings had no remarkable effect on the microstructure in the as-deposited state and during annealing in argon atmosphere.

221

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