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

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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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On the thermal stability of na-noscaled Cr/ta-C multilayers

U. Ratayski 1, Ch. Schimpf 1, T. Schucknecht 1, U. Mühle 1,2,C. Baehtz 3, M. Leonhardt 4, H.-J. Scheibe 4, D. Rafaja 1

1 Institute of Materials Science, TU Bergakademie Freiberg, Gustav-Zeuner Str. 5, 09599 Freiberg, Germany2 Fraunhofer Institute for Nondestructive Testing IZFP, Maria-Reiche-Straße 2, 01109 Dresden, Germany3 Helmholz-Zentrum Dresden Rossendorf, Bautzener Landstr. 400, 01328 Dresden, Germany4 Fraunhofer Institute for Materials and Beam Technology IWS, Winterbergstr. 28, 01277 Dresden, Germany

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AbstractThe influence of the initial microstructure in the as-deposited Cr/ta-C multilayer coatings on the thermal stability of the tetrahedrally bonded carbon (ta-C) layers and on the changes of the multilayer morphology with increasing annealing temperature was investigated. The multilayer structure was prepared by the combined DC arc/laser arc technology. The DC arc evaporation was employed for the deposition of the Cr layers, the laser arc technology for the deposition of the ta-C layers. The laser arc process was performed at three different energies of the carbon ions (Ec) in the range between ~25 eV and 500 eV in order to modify the degree of intermix-ing of species at the interfaces. In situ and ex situ small-angle and wide-angle X-ray scattering (SAXS, WAXS) and high resolution transmission electron microscopy (HRTEM) were used for the characterization of the microstructure evolution of the Cr/ta-C multilayer coatings. The in situ experiments were performed after each annealing step up to annealing temperature of 600 °C. SAXS revealed the changes in the density and thickness of the individual layers in the multilayer stack and in the interface morphology and correlation of the interface roughness. The density of the ta-C layers decreased with increasing annealing temperature, which can be explained by the graphitization of the ta-C layers. The grade of the graphitization depends on the sp³/sp² ratio in the as-deposited multilayers. In samples with originally higher sp³ fraction, ta-C graphitizes at higher temperatures than in samples with a lower sp³ fraction. The thermal treatment led to a significant smoothening of the Cr/ta-C interfaces in the multilayers with a wide intermixing zone and facilitated the formation of metastable fcc-CrC via interdiffusion of C and Cr.Keyword: DLC; multilayer; cathodic arc evaporation; Cr/ta-C; synchrotron; in situ

experiments; SAXS; TEM

Introduction

Diamond-like carbon (DLC) coatings are es-tablished as tribological and wear resistant protective coatings in abrasive environment, e.g., as coatings for piston rings or pins or for inserts in diesel injection devices, etc. [1-3]. The properties of DLC films varies from gra-phitic-like to diamond-like, thus correspond-ingly DLC’s show a wide range of properties [4]. The environmental constraints for the automotive industry are mainly responded by the reduction of friction and wear, which requires the development of more efficient, advanced wear protective coatings, for what DLC coatings are prominent candidates. However, a limiting factor for the application of DLC’s is the adhesion to the substrate [5, 6].

DLC’s are divided into three groups [6, 7]: (i) highly hydrogenated amorphous carbon (polymeric a-C:H), (ii) hydrogenated amor-phous carbon and (iii) non-hydrogenated amorphous carbon (a-C, ta-C). These three groups are further classified according to the ratio of sp³ (diamond) and sp² (graphit-

ic) bonds [6, 7]. However, the amount of sp³ bonds alone is not sufficient to describe the various properties of amorphous carbon coat-ings, e.g. tetrahedrally bonded diamond-like carbon (ta-C) films exhibit high hardness and low coefficient of friction, while polymeric amorphous carbon films (a-C:H) are ductile and exhibit a relatively low hardness. Both DLC’s containing a high amount of sp³ bonds. Therefore, the differentiation of DLC’s by the mechanical and physical properties is more practical.

The formation of ta-C coatings requires high compressive stress in order to stabilize sp³ bonds, which can be achieved by the ion bom-bardment [5]. Thus, the enhancement of the adhesion of the ta-C coatings is required to prevent the delamination [8]. A possible strat-egy for the improvement of the adhesion are the surface pretreatment of the substrate by metal ion bombardment and the deposition of a buffer layer between the substrate and the DLC film consisting of transition metals such

as Ti, Cr, Nb or V [2] as well as the deposition of multilayer coatings. Among the transition metals, chromium is a prominent candidate due to the formation of three stable carbides at moderate annealing temperatures. The chromium carbides with narrow homogene-ity ranges are assumed to act as efficient barri-ers for further diffusion [9]. Furthermore, the environmental temperature plays a crucial role for the usability of ta-C coatings, because their application temperature is limited to about 300 °C due to the graphitization of the sp³ bonds [6].

In our first contribution the effect of the car-bon ion energy (EC) on the microstructure of Cr/ta-C multilayers [10] were investigated ex situ. It was shown that the carbon ion energy used for the deposition of the ta-C layer chang-es the amount of sp³ bonds in the carbon lay-ers, which reached the maximum at medium Ec = 200 eV. A further increase of the carbon ion energy led to the decrease of the tetrahe-drally bonded carbon fraction. Furthermore, both the morphology and the quality of the Cr/ta-C interfaces were affected significantly by the energy of the carbon ions. In order to be able to describe the kinetics of the phase transitions and the related microstructure changes in the Cr/ta-C system in more details, three Cr/ta-C multilayer coatings were pre-pared in analogy to [10] by using a combined DC arc/pulsed arc technique and investigat-ed in situ during annealing up to 600 °C in vacuum. The microstructure evolution caused by the thermal treatment was investigated by glancing angle X-ray diffraction (GAXRD), small angle X-ray scattering (SAXS), X-ray reflectivity (XRR) and transmission electron microscopy (TEM) with high resolution.

Fig. 1: Scheme of the Sulzer Metaplas MZR324 PVD coater equipped with three DC arc sources (1), with a laser arc module with rotating graphite cathodes (2) and with a planetary sample holder in the vacuum chamber (3).

Film preparation

The deposition of the Cr/ta-C multilayer coat-ings was performed in the industrial PVD coater Sulzer Metaplas MZR324. The cham-ber was equipped with three DC arc cathodes made from Cr (diameter 63 mm) and a laser

Experimental Details

controlled pulsed arc module (LAM) with a cylindrical graphite cathode (length 400 mm and a diameter of 160 mm), as schematically illustrated in Fig. 1. No additional gases were required for the coating preparation, thus the base pressure was kept unchanged at 10-3 Pa during the whole deposition process. No ad-ditional substrate heating was applied. There-fore, the substrate temperature did not exceed 100 °C in order to ensure the formation of amorphous diamond-like carbon layers con-taining a high amount of sp³ bonds. The Cr/ta-C films were deposited onto polished (001) oriented Si wafers, which were placed on the outer circle of the planetary sample holder.

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In order to avoid the increase of the surface roughness of the Si substrates, no ion-assisted substrate cleaning was performed prior to the deposition. The periodic motif of Cr and ta-C layers was built up through the rotation of the sample holder and the subsequent switching of the DC arc and LAM sources. The nomi-nal thickness of individual layers was set to ~ 10 nm; it was controlled by the deposition time for the Cr layers and the number of im-pulses for the ta-C layers in synchronization to the rotation speed of the sample holder. The intended total layer thickness of the mul-tilayer stacks was below 120  nm in order to be able to resolve Kiessig oscillations in the XRR curves. Therefore, the periodic motif (Cr/ta-C) was repeated 6 times in the coat-ings. In order to investigate solely the effect of the carbon ion energy on the microstructure evolution during the thermal treatment, the process parameters for the Cr layers were kept constant for the three multilayer films and the substrate bias was set to -50  V. In contrast, different peak currents and pulse lengths were adjusted during the deposition of ta-C layers in subsequent deposition runs in order to ob-tain three different mean carbon ion energies (Ec) of about 25 eV (DC1), 200 eV (DC2) and 500 eV (DC3).

Microstructure characterization and thermal treatment

The microstructure evolution during the an-nealing of the Cr/ta-C multilayer coatings was analyzed by in situ synchrotron experiments at the Rossendorf beamline (ROBL) BM20 at the European Synchrotron Radiation Facili-ty (ESRF) in Grenoble. The diffractometer at ROBL was equipped with a vacuum chamber (~6x10-5 mbar) with a heating stage installed that enabled in situ experiments within a tem-perature range between room temperature and 600 °C. The annealing of the Cr/ta-C layers was performed by a gradual increase of the annealing temperature from room tem-perature to 600 °C in steps of 100 K. The sam-ples were annealed for 1 h at each annealing temperature. After each temperature step, the Cr/ta-C multilayers were cooled down to

100 °C for the synchrotron measurements. The photon energy for the synchrotron exper-iments was set to 11.5 keV (λ = 0.107812 nm). The diffraction patterns were recorded using a 0D Mythen detector. By means of the in situ synchrotron X-ray reflectivity (XRR) exper-iments, the microstructure was described in terms of the changes of the mean mass den-sity ρ and the layer thickness t with increas-ing temperature. Furthermore, the crystalline phases were revealed by in situ glancing angle X-ray diffraction (GAXRD) experiments. The interface correlation in the multilayers was obtained from the measurement of the res-onant diffuse scattered intensity in the small angle region by laboratory ex situ SAXS ex-periments and from the HRTEM images.

The SAXS measurements were performed on the laboratory diffractometer D8 Advanced (BRUKER AXS), which was equipped with a sealed X-ray tube with a copper anode (λ = 0.15418 nm) and a Goebel mirror in the primary beam in order to assure a parallel pri-mary beam with a divergence of approximate-ly 150 arc sec. The primary beam was reduces in size by using a slit with a width of 0.1 mm in order to prevent the irradiation of the sam-ple holder. In the secondary beam, two slits with a width of 0.1  mm were used. A small divergence of the primary beam is necessary to achieve a high resolution in the reciprocal space. The SAXS measurements were per-formed as a series of ω – scans composed to reciprocal space maps (RSM).

Transmission electron microscopy (TEM) was performed in the high resolution (HR-TEM) mode of the analytical transmission electron microscope JEM 2020 FS from JEOL, which was equipped with an in column filter and a spherical aberration corrector (Cs) lo-cated in the primary beam. The TEM operat-ed at the acceleration voltage of 200 keV. TEM investigations yielded additional information of the morphology of the multilayers and the interface quality. The TEM investigations were performed on the Cr/ta-C multilayers in the as-deposited state as well as ex situ after the synchrotron annealing experiments at 600 °C for 1h.

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Microstructure of the as-deposited Cr/ta-C multilayer coatings

The periodicity of the Cr/ta-C motif in the multilayers under study was confirmed by XRR measurements (see Fig. 2) and TEM images (see Fig. 3). The XRR curves contain the multilayer features such as Bragg maxima from the periodic motif and the Kiessig oscil-lations related to the total multilayer thick-ness. In order to compare the SAXS measure-ments at different wavelength the diffraction angle θ was converted into the diffraction vector

Due to a variation of the individual layer thickness in the Cr/ta-C multilayers Bragg maxima weaken with increasing diffraction vector qZ. The Bragg maxima of DC1 and DC2 are broader than for sample DC3, which is caused by the variation of the thickness of the periodic motif. The thickness variations of the periodic motif seen by XRR measure-

Results ments were caused either by a chemical gra-dient at the interfaces or by the sample holder rotation during the deposition process. The position of the sample holder could not be controlled to guarantee the same initial posi-tion when switching from Cr cathode to the graphite cathode. Furthermore, the decay of the Bragg maxima was more pronounced with increasing carbon ion energy (see sam-ples DC1 and DC3 in Fig. 2) indicating the loss of the well pronounced periodicity of the multilayer stacks. As seen in Fig. 2 the mea-sured intensities of the XRR curves decreased faster with increasing diffraction vector qZ for the Cr/ta-C multilayer DC3. The decline of the XRR curves is proportional to the sur-face roughness [11, 12]. The high carbon ion energy used for the preparation of the mul-tilayer sample DC3 caused a significant in-crease of the surface roughness.

TEM imaging as shown in Fig. 3 confirmed the gain of the surface and interface rough-ness of sample DC3 (Fig. 3b). The trend of the effect of the carbon ion energy on the inter-

Fig. 2: X-ray reflectivity curves measured for Cr/ta-C multilayercoatings deposited at variouscarbon ion energies of approxmately 25 eV (DC1), 200 eV (DC2)and 500 eV (DC3).

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Fig. 3: Bright field TEM images of the Cr/ta-C multilayer coatings deposited at EC ~25 eV (a) and at EC ~500 eV (b).

face roughness was already shown in [10]. The Cr layers showed a columnar growth, where-as the ta-C layers were amorphous as seen by the extinction of the diffraction contrasts in the bright field TEM images. The amorphous nature of the ta-C layers was confirmed by the Fast Fourier Transformation (FFT) of the carbon layers of the marked area in the ta-C layer of sample DC1 (see Fig. 4b) and sample DC3 (Fig. 5b). The FFT only showed a diffuse halo without any intensity maxima. The phase analysis performed by GAXRD revealed the body centered cubic (bcc) Cr as the only crys-talline phase. These results are in agreement with our previous published results [10]. The FFT of the Cr layers deposited at a low carbon ion energy (Fig. 4c) confirmed the crystallin-ity of Cr and revealed a <011> orientation of bcc-Cr in the growth direction. On contrary, the FFT of the Cr layer deposited at high car-bon ion energy of ~500 eV (Fig. 5c) showed an intensity ring. The absence of the intensity maxima indicated the loss of the crystalline nature of the Cr layers due to the high car-bon ion energy DC1. These results are in good agreement with our previous contribution on the Cr/ta-C multilayer coatings [10]. Further-more, TEM imaging shows well a different quality of the interfaces between the Cr and the ta-C layers for sample DC1 (see Fig. 4a) and DC3 (see Fig. 5a).

The refinement of the XRR curves yielded information about the thickness of the indi-vidual layers t, the mean thickness of the Cr and ta-C layers, the density of the individual layers ρ and on the interface roughness . The XRR curves were calculated by using the recursive optical Parratt formalism [13] and the computational routine from [14, 15] based on the Fullerton approach [12]. The XRR analysis revealed that the mean den-sity of the ta-C layers (ρC) increases from Ec~25  eV to ~200  eV from 2.8 to 3.0  g/cm³ and declines if the carbon ion energy is fur-ther increased to about 500 eV (see Fig. 6a). The lowest carbon density was obtained for the Cr/ta-C coating deposited at the highest carbon ion energy (Ec~500  eV). A similar trend was found for the mean density of the Cr layers (ρCr) in the multilayer stacks. The mean density of the Cr layers was in all three Cr/ta-C multilayers higher than the theoreti-cal density of Cr with 7.2 g/cm³. The higher mean density of the Cr layers can be caused either by the incorporation of C in the Cr lay-ers, by a high interface roughness between the Cr and the ta-C layer and/or the generation of a diffuse interface. The large error margin of the determined densities of the Cr and the ta-C layers references the correlation of the interface roughness and of the inhomogene-ity of the layers on the scattering potential.

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The used computational model treats the in-terface roughness statistically and these effects yield similar effects on the scattering potential of the multilayer stack. Therefore, the used re-finement method cannot distinguish between surface roughness and the variation of the electron density of the layers caused by in-corporations or diffuse interfaces. Therefore, the large error margin can be understood as artifact of the scattering potential. The refined density and layer thickness of the Cr/ta-C multilayer are graphically summarized in Fig. 6c) and Fig. 6d), where it is shown that the mean layer thickness of the Cr and the ta-C layers decreased with increasing carbon ion energy. The individual layer thickness of the ta-C and the Cr layers decreased from about 10 nm to 6.5 nm (see Fig. 6c) and 12 nm

to 8  nm (Fig. 6d) for sample DC1 to DC3, respectively. The increase of the carbon ion energy rise the probability of resputtering of the carbon at higher energies, which explains more pronounced decrease of the ta-C layer thickness. If the kinetic energy of the car-bon ions is sufficient high the carbon can be ejected from the growing film by breaking the carbon bonds of the surface layers. Further-more, at higher kinetic energies of the carbon can be incorporated in the Cr interface by ion implantation.

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Fig. 4: HRTEM image of the Cr/ta-C multilayer coatings deposited with a low carbon ion energy of ~25 eV (a) and the FFT of the marked area of the ta-C (b) and the Cr (c) layer. The arrow marks the growth direction.

Fig. 5: HRTEM image of the Cr/ta-C multilayer coatings deposited with a high carbon ion energy of ~500 eV (a) and the FFT of the marked area of the ta-C (b) and the Cr (c) layer. The arrow marks the growth direction

The influence of the carbon ion energy on the interface quality was investigated by re-ciprocal space maps (RSM) in the small an-gle region (Fig. 7) and was confirmed by TEM investigation (see Fig. 3). The Cr/ta-C multilayers deposited at a carbon ion ener-gy of about 25  eV exhibited a well-defined interface between the individual layers (see Fig. 3a, Fig. 4a), whereas the interfaces of the Cr/ta-C multilayer coating deposited at a carbon ion energy of ~500 eV were strong-ly disturbed (Fig. 3b, Fig. 5a) indicating an enhanced defect density at the interfaces due to the ion bombardment. The influence of the carbon ion energy was further revealed by RSM measurements exemplarily for sam-ple DC1 (Fig. 7a), DC2 (Fig. 7b) and DC3 (Fig. 7c). The RSM of sample DC1 (Fig. 7b), which was deposited at Ec ~25 eV for the car-bon ions, exhibited the typical features of the resonant diffuse scattered intensity of a mul-tilayer coating such as the specular reflect-

ed intensity along qx = 0, well pronounced Yoneda wings and resonant diffuse scattering bananas parallel to the qx axis. The increase of the intensity of the resonant diffuse scatter-ing (so called Holý bananas) is caused by the partial correlation of the interface roughness profiles in the multilayer stack [16] and the position of the Holý bananas intersect the Bragg peaks of the specular reflected intensity. Furthermore, dynamical scattering effects are visible for sample DC1 such as the Bragg-like resonant lines, which intersect the Bragg peaks of the specular reflected intensity and process parallel to the Yoneda wings. The Bragg-like lines occur if the diffraction con-dition of the incident and outgoing waves are fulfilled. Furthermore, the intersection of the Bragg-like lines yields maxima of the diffuse scattered intensity, which are called Bragg-like resonant peaks. The appearance, the form and the intensity of these dynamical effects de-pends on the interface correlation, whereas

Fig. 6: Summary of the mean mass density of the C (a) and the Cr layers (b) in the as deposited Cr/ta-C multilayercoatings.

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the dynamical effects occur independently of the actual interface correlation function. The RSM of the Cr/ta-C multilayer coating DC3 deposited at Ec ~500 eV (Fig. 7c) showed the absence of the Holý bananas and a continu-ous decrease of the diffuse resonant scattered intensity for increasing qz values. The absence of the Holý bananas indicates the loss of the vertical interface roughness correlation in the multilayer stack, which was also seen by TEM investigations (see Fig. 3b). However, the Yoneda wings can be observed in the sample DC3. The width of the Holý bananas in the qz

direction can be understood as a measure of the replication of the interface profile in the multilayer stack. Therefore, the degradation of the diffuse resonant scattered intensity of the Holý bananas and their broadening are caused by the loss of the vertical interface cor-relation in the Cr/ta-C multilayers caused by the high defect density at the interfaces.

Fig. 7: Reciprocal space maps (RSM) ofthe Cr/ta-C multilayer coatings in theas-deposited state. The RSM are sortedby increasing carbon ion energy fromthe top to the bottom with the order of (a)25 eV (DC1), (b) 200 eV (DC2) and (c)500 eV (DC3).

Microstructure evolution of Cr/ta-C multi-layers during the thermal treatment

The Cr/ta-C multilayers were annealed up to an annealing temperature of 600 °C for 1h per temperature step. The thermal treatment yielded information about the microstructure evolution of the multilayer stacks, especially on the stability of the diamond-like ta-C lay-

ers and on the interdiffusion processes at the Cr/ta-C interfaces. From the XRR analysis, the evolution of the mean densities of the Cr and the ta-C layers in dependence on the an-nealing temperatures are presented in Fig. 8a. The XRR refinement revealed the decrease of

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the ta-C density in all three multilayer stacks above an annealing temperature of 300 °C (Fig. 8a). Above annealing temperatures of 400 °C, the decrease of the ta-C density was decelerated in all three layers and the density loss achieved saturation above the theoretical density of graphite. It is assumed that the de-crease of the mean density of the ta-C layers is caused by the transformation of the sp³ bonds to the sp² bonds, which can be understood as graphitization of the ta-C layers. However, the refined ta-C layer densities after the annealing at 600 °C varied in the three multilayers be-tween 2.39±0.2 g/cm³ (DC1), 2.49±0.1 g/cm³ (DC2) and 2.3±0.1  g/cm³ (DC3). The high-est density of carbon after the annealing was observed in sample DC2, which was depos-ited at the carbon ion energy of about 200 eV. The Cr/ta-C multilayer coating DC2 also ex-hibited the highest density of carbon in the as-deposited state. The lowest C density after the thermal treatment was observed for sam-ple DC3, which was deposited at the highest carbon ion energy during the coating pro-cess. These results indicate that the expected highest amount of the sp³ bonds in the as-de-posited ta-C layers correlates with the higher density obtained from the XRR analysis. In our first report on the Cr/ta-C layers [10], the amount of sp³ and sp² bonds in DLC coatings was estimated from the intensity ratio of the π* and σ* peaks [17] in the electron energy

loss spectra, which were measured locally in a transmission electron microscope. The com-parison of the densities of the carbon layers with the amount of the sp³ bonds revealed al-most linear relationship between these quan-tities. In this work, the above relationship was used to predict higher amount of sp² bonds in the ta-C layers from their lower density. It was concluded that the amount of sp³ bonds the after thermal treatment depends on the amount of sp³ bonds in the as-deposited ta-C layers. More sp³ bonds in the as-deposited layers lead to more sp³ bonds in the annealed layers.

Also for the Cr layers, a decrease of the mean density was observed (Fig. 8b). In the tem-perature range between 200 to 300 °C, the Cr densities were observed to decrease only slightly in the Cr/ta-C multilayers. Howev-er, in contrary to the ta-C layer a significant decrease of the mean Cr layer densities were observed at annealing temperatures above 300 °C, where the Cr densities decreased to approximately 6.5  ±  0.3  g/cm³ in all three Cr/ta-C multilayer stacks. After the thermal treatment, the determined Cr layer density is lower than the theoretical densities of pure Cr of 7.2 g/cm³ and the three thermodynam-ically described chromium carbides Cr23C6, Cr6C3 and Cr3C2 having the tabulated densi-ties of 6.95, 6.88 and 6.6 g/cm³, respectively.

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Fig. 8: Densities evolution in the ta-C(a) and the Cr layers (b) depending on the annealing temperatures and the energy of the carbon ions 25 eV (DC1 - black), 200 eV (DC2 - blue) and 500 eV (DC3 - red). The gray planes represent the theoretical density of graphite with 2.26 g/cm³ (a), Cr and the metastable CrC (b). The theoretical density of diamond is 3.51 g/cm³, which is not shown in (a)

However, no significant changes of the Cr and the ta-C layers were observed by XRR refine-ment due to the thermal treatment of the Cr/ta-C multilayer coatings. The formation of the metastable CrC by interdiffusion would led to the change of the thickness ratio of the Cr and ta-C layers therefore it can be assumed that the formation of the metastable CrC was caused by incorporated C in the Cr layers due to ion implantation during the deposition process [19]. The TEM investigation showed the reduction of the ta-C layer thickness. The thickness reduction could not be observed by XRR analysis because of the chemical gradi-ent in the as-deposited samples.

As it can be seen in Fig. 9, the interface morphology of the Cr/ta-C multilayer DC3 changed significantly after the heat treatment. In the as-deposited Cr/ta-C coating (Fig. 3b and Fig. 5a), the interfaces were blurred, whereas the interfaces in the annealed mul-tilayer stack appeared sharp and smoothed (Fig. 9). The smoothening of the interfaces was also observed by the RSM measurements (Fig. 10). The measured resonant diffuse scat-tered intensity of the Cr/ta-C multilayers were enhanced, whereas the specular reflected in-tensity for sample DC1 and DC2 (Fig. 10a,b)

Fig. 9: HRTEM image of the Cr/ta-Cmultilayer coating deposited with a highcarbon ion energy of ~500 eV after thethermal treatment at 600 °C for 1h. Theinset represents the FFT of the markedarea of the Cr layer. The arrow marks thegrowth direction.

However, the lower Cr density might be ex-plained by the formation of the metastable CrC phase in the Cr layers. The formation of the metastable face centered cubic (fcc) CrC phase was reported by Bewilogua et al. [18] as well as Wang et al. in 1993 [19]. However, the existence of the metastable CrC phase is discussed due to the discrepancy to the em-pirical Hägg’s rule, which limits the formation of the solid solutions of metalloid – metal by the radius ratio rc/rMe < 0.59. The stabilization of the metastable fcc-CrC phase was reported to be caused by the high ion dose during the ion implantation of C ions [19]. The theoret-ical density of the fcc-CrC phase is tabulated with 6.49  g/cm³ as calculated for the NaCl structure type and the cubic lattice parame-ter of a = 4.03 Å, which is in good agreement with the determined Cr densities for all three Cr/ta-C multilayer coatings. The existence of the metastable fcc-CrC phase was supported by the TEM investigations of sample DC3 af-ter the thermal treatment at 600 °C (see Fig. 9). The TEM investigation showed that the ta-C layers are amorphous, whereas the ther-mal treatment of the Cr/ta-C multilayers im-proved the crystalline nature of the Cr layers (Fig. 9) and the FFT of the marked area (inset Fig. 9) could be indexed as the fcc-CrC phase.

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disappeared in the RSM measurements after the thermal annealing at 600 °C (Fig. 7a vs. Fig. 10a). The GAXRD analysis showed ad-ditional peaks of chromium silicides above annealing temperature of 400 °C and TEM investigations of the annealed multilayers revealed the formation of the chromium sil-icides at the Si/Cr interfaces. The formed reaction zone exhibited a rough, corrugated interface between the chromium silicides and the Si substrate, which was responsible for the disappearance of the specular reflected inten-sity in the RSM. However, the Holý bananas (maxima of the resonant diffuse scattering) were still present in the RSM measurements, which revealed that the multilayer structure of the remaining Cr/ta-C multilayer with cor-related interface roughness was still present after the thermal annealing. For sample DC3 (Fig. 10c), local intensity maxima of the dif-fuse scattered intensity were observed at the related Bragg peaks. In the measured RMS of the as-deposited multilayers (Fig. 7c), the

intensity of the resonant diffuse scattering decreased continuously with increasing dif-fraction vector qz with no local maxima of the diffuse scattered intensity, which was caused by the blurred and corrugated Cr/ta-C inter-faces with no vertical correlation in the as-de-posited state. The intensity and the width of the Holý bananas are a measure of the vertical correlation of the interfaces [16]. For samples DC1 and DC2, the increase of the width of the Holý bananas indicated the decrease of the vertical correlation of the interface rough-ness profile. In the RSM of sample DC3, the Holý bananas are weak but still visible indi-cating the sharpening of the Cr/ta-C interfac-es. The sharpening of the interfaces and the improvement of the vertical correlation of the interface roughness profile were confirmed by TEM and HRTEM investigations as seen in Fig. 9.

Fig. 10: Reciprocal space maps (RSM) of the Cr/ta-C multilayer coatings after a thermal treatment up to 600 °C. The RMS are sorted according to increasing carbon ion energy Ec applied during the deposition process from the top to the bottom with the order of (a) 25 eV (DC1), (b) 200 eV (DC2) to (c) 500 eV (DC3).

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Discussion

The carbon ion energy affected the density of the ta-C layers, which is directly proportion-al to the sp³/sp² ratio in the ta-C layers. The sp³/sp² ratio reached a maximum at the me-dium carbon ion energy. A further increase of the carbon ion energy led to the graphiti-zation of the ta-C layers. This trend was ob-served already in our first report on the Cr/ta-C multilayers [10]. The presented results of the XRR refinements showed the decrease of the mean carbon density with increasing annealing temperature. In all three Cr/ta-C multilayers, the density reached saturation above the annealing temperature of 500 °C. The mean carbon density of the ta-C layers remained higher in the multilayer coatings with a higher sp³/sp² ratio in the as-deposited state. These results indicated that the amount of sp³ bonds in the as-deposited multilayers plays a crucial role for the graphitization of the ta-C layers. Furthermore, the saturation of the transformation from sp³ to sp² bonds was shifted to higher annealing temperatures (see Fig. 8a) with increasing mean carbon density. Sample DC3 reached the saturation of the density loss at the annealing temperature of 300 °C, whereas the decline of the carbon density in DC2 was shifted to the annealing temperatures above 500 °C. Therefore, we as-sume that the initial amount of the sp³ bonds in the as-deposited state is the limiting factor the sp³ to sp² conversion in the Cr/ta-C mul-tilayer coatings. The formation of sp³ bonds in DLC coatings is caused by the subplanta-tion of the carbon ions but the sp³/sp² ratio decreased at high carbon ion energies [4]. The decrease of the fraction of the sp³ bonds with increasing carbon ion energy can be explained by the thermal spike model [4, 20], which de-scribes the thermal activated relaxation of the sp³ bonds and the formation of the stable sp² bonds during the deposition. Furthermore, the amount of sp³ bonds in ta-C coatings is directly correlated to the intrinsic compres-sive stress in these coatings [4]. The increase of the sp³/sp² ratio increases the compressive stress in DLC coatings [21]. The chemical bonding of the ta-C coating is shifted above the stability criterion for sp³ and sp² known

as Berman-Simon line [21]. We assume that due to the increase of the sp³/sp² ratio in the investigated Cr/ta-C multilayers, the thermal activation barrier for the transformation from sp³ to sp² bonds is shifted to higher values, be-cause the higher compressive stresses in the ta-C layers stabilized the sp³ bonds. There-fore, the thermal activation barrier for the graphitization of the sp3 bond is enhanced. Furthermore, the transformation of sp³ to sp² bonds is assumed to be decelerated due to the high compressive stress in the ta-C layers.

As already discussed in our first contribution [10], the interface quality in the Cr/ta-C mul-tilayers is strongly affected by the carbon ion energy, which was applied for the deposition of the ta-C layers. The increase of the carbon ion energy led to an increasing impact of the carbon ions in the Cr layers, which caused an increasing intermixing zone at the interfaces between Cr and C and the incorporation of C in the Cr layer. The intermixing zone can be explained by the subplantation model [4]. The penetration depth of the carbon ions can be explained as a function of the carbon ion en-ergy with a certain penetration threshold (Ep). The penetration threshold is proportional to the difference of the surface binding energy and the displacement energy of the atoms in the coating [4, 22]. Therefore, at a certain ion energy the carbon ions are able to penetrate the surface layer [23]. With increasing carbon ion energy the penetration depth of the carbon ions increases. However, the required energy for the penetration is rather small compared to the total carbon ion energy. The excess en-ergy is used for the displacement of the car-bon atoms and is also converted to thermal energy. The influence of the energy fraction transformed to thermal energy is described in the model of thermal spikes [4, 20]. The thermal energy enables the diffusion of the atoms in the growing films additionally to the intermixing due to subplantation. Therefore, the combination of an increased penetration depth due to the increased carbon ion ener-gy and the thermal activated diffusion in the growing film can be assumed as the driving

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factors for the formation of the intermixing zone at the Cr/ta-C interfaces and the incor-poration of C in the Cr layers.

The thermal treatment showed the sharpen-ing of the Cr/ta-C interfaces in the multilayer stacks. The XRR and HRTEM investigations showed that the density of the Cr layers de-creased significantly and the formation of the metastable CrC is assumed at higher an-nealing temperatures in the Cr/ta-C multi-layers. According to our results, the thermal treatment initiated the interdiffusion process at the Cr/ta-C interfaces. The decrease of the mean carbon density indicated that no diffusion of Cr in the ta-C layer occurred. A decrease of the Cr layer density occurred at annealing temperature above 300 °C indicat-ing the diffusion of the carbon atom into the Cr layer of the recombination of the Cr layers and the incorporated C atoms. Therefore, the diffusion of Cr can be neglected. In contradic-tion to the thermodynamic prediction for the Cr-C system [24], the formation of the three stable chromium carbides was not observed. If the three stable chromium carbides would have formed at the interfaces, they would oc-cur as very thin layers, which could not be ob-served by XRR. However, the decrease of the Cr density indicated the formation of the metastable fcc-CrC carbide.

According to [18], we assume that the ion bombardment during the deposition of the Cr/ta-C multilayers and the formation of the high compressive stress are the driving factors for the stabilization of the metastable fcc-CrC phase during the thermal treatment. The metastable fcc-CrC phase contains the highest carbon amount of the chromium car-bides and it is likely formed at the Cr/ta-C in-terfaces. Due to the high sp³/sp² ratio in our Cr/ta-C multilayers, the compressive stresses were assumed to be high as well, which can shift the stability criteria for the metastable fcc-CrC phase. According to Bewilogua et al. [18], the metastable fcc-CrC phase trans-forms to the stable Cr3C2 during the thermal treatment, which could not be observed in the Cr/ta-C multilayer coatings. The phase tran-

sition from the metastable CrC to the stable Cr3C2 phase is connected with an increase of the mass density and with an increase of the molar volume. Therefore, the compres-sive stresses would further increase due to the phase transformation of the metastable CrC phase. The transformation of sp³ to sp² bonds in the ta-C layers is linked to the de-crease of the mass density and to an increase of the occupied molar volume so an increase of the compressive stress. Assuming the in-terdiffusion of carbon from the ta-C into the Cr layer to form CrC, the thickness of the Cr layer should increase by 40 % at a total deple-tion of the graphitic layer and a depletion of 70 % assuming a pure diamond layer. How-ever, such significant changes of the layer thicknesses during to the thermal treatment were not observed by XRR; neither for the ta-C layers nor for the Cr layers. Therefore, it can be assumed that the formation of the CrC phase in Cr layer due to the thermal treatment is rather caused by the ion implantation of C in the Cr layer during the deposition than by interdiffusion processes during the annealing from the ta-C layers into the Cr layers. How-ever as seen by HRTEM (compare Fig. 5 and Fig. 9), the ta-C layers in the multilayer stack after the thermal treatment appeared thin-ner than in the as-deposited state and their interfaces to the Cr layers are sharper. The carbon ion bombardment produced an inter-mixed zone in the as-deposited state as seen for DC3 (Fig. 5), which appeared in the XRR curves either as variation of the layer thick-ness or as additional interface roughness. Due to the blurred interfaces in the as-deposited state the decrease of the layer thickness could not be observed by XRR measurements. The sharpening of the interfaces after the ther-mal treatment of the multilayer stacks can be explained by diffusion processes in the orig-inally blurred interfaces. Due to the narrow homogeneity ranges, the gradual change of the carbon concentration has to be reshaped in order to achieve the concentration in the individual chromium carbide phases.

197

Conclusion

In the as-deposited Cr/ta-C multilayer coat-ings, the sp³ fraction in the ta-C layers was shown to modify the thermal stability of the ta-C layer and the amount of preserved sp³ bonds significantly. The increase of the sp³ fraction in the as-deposited state shifted the thermal activated transformation from sp³ to sp² to higher annealing temperatures due to the decelerated transformation of sp³ to sp² bonds due to a higher sp³ fraction. A model was proposed to explain the influence of the sp³ bonds and the compressive stress to ex-plain the density loss saturation depending on the sp³/sp² ratio in the as-deposited coatings. The multilayer structure of the Cr/ta-C coat-ings was still preserved after the annealing for 1h at 600 °C and a smoothening of the blurred interfaces were observed. We proposed the in-terdiffusion of carbon atoms due to the ther-mal treatment to explain the smoothening of the interfaces in the Cr/ta-C multilayers. Furthermore, the formation of metastable fcc-CrC is assumed.

Acknowledgement

This work was performed within the Clus-ter of Excellence “Structure Design of Nov-el High-Performance Materials via Atomic Design and Defect Engineering (ADDE)” that is financially supported by the Euro-pean Union (European Found for Regional Development) and by the Ministry of Sci-ence and Art of Saxony (SMWK).

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