Temperature impact on thermal evolution of advanced PVD ceramic

N° d’ordre 2015ISAL0009 Année

Thèse

Temperature impact on thermal evolution of advanced PVD ceramic and metallic glass thin films: physico-chemical and microstructural analysis

Présentée devant L’institut national des sciences appliquées de Lyon Pour obtenir Le grade de docteur École doctorale École doctorale materiaux de Lyon (ED 34) Par Mihai APREUTESEI (Ingénieur) Soutenue le 4 février 2015 devant la Commission d’examen

Jury. M. A. BILLARD Professeur (LERMPS-UTBM), Examinateur M. P. DJEMIA Professeur (Université Paris 13), Examinateur M. C. ESNOUF Professeur (INSA de Lyon), Examinateur M. L. JOLY-POTTUZ Maître de conférences, (INSA de Lyon), Co-directrice de thèse M. N. MARTIN Professeur (Institut FEMTO-ST), Rapporteur M. P. STEYEr Maître de conférences, HDR (INSA de Lyon),Directeur de thèse M. F. VAZ Associate professor, HDR (Univ. do Minho),Rapporteur

Laboratoire de recherche : MATEIS, CNRS UMR 5510 de l’INSA-Lyon

Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2015ISAL0009/these.pdf © [M. Apreutesei], [2015], INSA de Lyon, tous droits réservés

Abstract In the recent years the industrial requirements to develop new functional materials able to overcome the severe conditions during machining operation are continuously increasing. Researchers then must find novel solutions to respond to their severe industrial requirements. To coat the tool surface with advanced coatings is the most efficient solution. New nanostructured materials may nowadays exhibit unique mechanical, physical and chemical properties ensuring notable degradation resistance where the surface protection of materials against corrosion, wear, friction or oxidation is a key issue, particularly when operating in hostile environments. Within the scope of this Ph.D. thesis the influence of the temperature on the structural stability of two different PVD ceramic and metallic glass thin films is proposed. The main goal consists in the development of two distinct classes of thin films, with a wide range of properties. In order to prepare these films, the project will be focused on the study on the influence of PVD deposition conditions in the particular film’s growth characteristics: chemical composition, structure, morphology and the subsequent changes in the main properties of the thin films, namely oxidation and crystallization resistance, especially. For that purpose we adopted the multiscale approach. The first part is related to the ceramic CrN-based coatings to give new functionalities and improve the tools’ surface with the primary aim to increase their lifetime. Secondly, new protective materials able to better protect the exposed surfaces against high temperature oxidation have been proposed, namely CrAlN and CrAlYN coatings as will be evidenced in this manuscript. Adding aluminum into CrN leads to the formation of a solid solution and improvement of oxidation resistance. During high temperature oxidation, the hard structure is preserved and the oxide protective scale composed of Al2O3 and Cr2O3 phases protects the surface preventing its degradation. The second part of the manuscript is dedicated to the innovative Zr-Cu thin films metallic glasses prepared by a PVD magnetron co-sputtering method with the objective to investigate the amorphization ability and their structural properties. Their excellent properties at room temperature have recently attracted attention as a new class of materials with great potential for engineering applications due to unique mechanical and physico-chemical characteristics (high elastic strain limit, corrosion resistance…) The main objective was to perform a detailed (micro)structural, chemical investigation with emphasis on thermal and mechanical properties. The metallic glass films were amorphous in a wide composition range (33.3 – 89.1 at.% Cu) with fine surface morphologies and partially vein-like features. The crystallization resistance was different with respect to the initial composition and structural state. Whatever the characterization technique considered, films belonging to the 45 – 75 at.% Cu compositional zone presented the best properties, which was explained on the basis of crystallization ability. Finally, an important approach during the course of this thesis was the real time observation of the structure and surface modifications during heating by means of in situ methods. The thin films proposed during the course of the work could be straightforward used as surface engineering solutions to protect and extend the lifetime of the materials and components.


Résumé Ces dernières années, les exigences de l'industrie dans le développement de nouveaux matériaux fonctionnels en mesure de résister aux conditions difficiles pendant l'opération d'usinage sont en constante augmentation. Les chercheurs doivent donc trouver de nouvelles solutions pour répondre aux besoins industriels de plus en plus sévères. L’utilisation de revêtement protecteur à la surface de l’outil de coupe est une solution très efficace. Des nouveaux matériaux architecturés sont étudiés pour leurs propriétés mécaniques, physiques et chimiques uniques assurant une résistance aux dégradations de surface dues à la corrosion, l'usure, le frottement; en particulier lorsque ces outils sont utilisés dans des environnements hostiles. Dans le cadre de cette thèse de doctorat, l'influence de la température sur la stabilité structurale de deux types de films minces déposés par PVD a été étudiée. Des films céramiques et de verre métallique ont été envisagés. Afin de préparer et optimiser ces films, le projet s’est axé sur l'étude de l'influence des conditions de dépôt sur les caractéristiques de croissance du film: composition chimique, structure, morphologie, puis sur les changements ultérieurs des principales propriétés des films minces, à savoir la résistance à l’oxydation et à la cristallisation lors de leur utilisation à hautes températures. Une démarche multi-échelle a été développée pour caractériser au mieux les couches. La première partie du travail est liée aux revêtements céramiques à base de CrN pour donner de nouvelles fonctionnalités et améliorer la surface des outils de coupe dans le but essentiel d'accroître leur durée de vie. Deuxièmement, des nouveaux matériaux capables de mieux protéger les surfaces exposées à l'oxydation à haute température ont été étudiés. CrAlN et CrAlYN ont été envisagés car l’ajout d'aluminium dans CrN conduit à la formation d'une solution solide et à l'amélioration de la résistance à l'oxydation. Au cours de l'oxydation à haute température, la structure du film est préservée et une couche protectrice d’oxydes composée de Al2O3 et Cr2O3 protège sa surface. La seconde partie du manuscrit est dédiée aux films minces de verres métalliques de Zr-Cu préparés par un procédé de co-pulvérisation magnétron PVD. Le but de cette partie consiste en l’étude de la relation entre la structure amorphe de ces films et leurs propriétés mécaniques. La conservation du caractère amorphe de ces films en température présente également un caractère essentiel. Les verres métalliques ont récemment attiré un fort intérêt car ils présentent des propriétés mécaniques intéressantes à température ambiante. Ils présentent, de ce fait, un grand potentiel pour des applications d'ingénierie en raison de leurs caractéristiques mécaniques et physico-chimiques uniques (haute limite élastique, résistance à la corrosion ...) Une analyse détaillée de la microstructure, des propriétés thermiques et mécaniques a été réalisée sur une large gamme de composition de films de verres métalliques amorphes (33,3 -. 89,1 at% Cu). La résistance à la cristallisation était différente en fonction de la composition et de la structure initiale du film. Quelle que soit la technique de caractérisation utilisée, les films appartenant à la gamme 45 -. 75 at% Cu présentent les meilleures propriétés, ce qui a été expliqué par leur structure initiale et leur propension à la cristallisation.


Pour relier les propriétés mécaniques des couches à l’évolution de leurs microstructures, une partie importante de ce travail a porté sur l’observation de l’évolution de la couche au cours du chauffage au moyen de techniques de caractérisation in situ. Les films minces proposés au cours de ce travail peuvent être envisagées pour un large gamme d’application dans l’ingénierie de surface pour protéger les surfaces et améliorer la durée de vie des matériaux.


Table of contents

General introduction ............................................................................ iv

Chapter 1 State of the art for the PVD ceramic and metallic glass thin films Part A. Hard coatings developed for high temperature applications .................................... 2 1.1. Introduction ............................................................................................................................. 2 1.2. Short flash-back of PVD thin films’ development ............................................................. 2 1.3. Conventional CrN thin films: general considerations and studies ................................. 3 1.4. Effect of Al addition into CrN: chemical nature and properties improvement ............. 6 1.5. High temperature oxidation: nature of oxide scales........................................................ 11 1.6. Interest of a further yttrium incorporation ....................................................................... 17 Part B. New class of thin films: metallic glass like .................................................................. 23 1.7. Interest of metallic glass-based materials .......................................................................... 23 1.8. Bulk metallic glasses ............................................................................................................ 24

1.8.1. Basic concepts regarding bulk metallic glasses ......................................................... 24 1.8.2. Formation and design of BMGs................................................................................... 26 1.8.3. Properties and features of BMGs................................................................................. 27 1.8.4. Crystallization and structural stability of metallic glasses-based materials ......... 30

1.9. Thin film metallic glass (TFMG) ......................................................................................... 32 1.9.1 How to get a TFMG? ...................................................................................................... 32 1.9.2. Binary TFMG systems ................................................................................................... 35 1.9.3. Binary Zr-Cu .................................................................................................................. 37 1.9.4. Mechanical properties of TFMGs ................................................................................ 38

1.10. Advanced applications ...................................................................................................... 40 Chapter summary ........................................................................................................................ 44 List of references .......................................................................................................................... 46

Chapter 2 Description of the experimental details and techniques

2.1. Synthesis of thin films .......................................................................................................... 60 2.1.1. Short history ................................................................................................................... 60 2.1.2. Deposition techniques .................................................................................................. 60 2.1.3. Deposition conditions ................................................................................................... 63

2.2. Thin films characterization .................................................................................................. 67 2.2.1. Chemical and crystallographic analysis ..................................................................... 67 2.2.2. Multiscale microstructural characterization .............................................................. 70


2.3. Functional characterization ................................................................................................. 74 2.3.1. Thermal analysis ............................................................................................................ 74 2.3.1.1. Thermogravimetric analysis (TGA) ......................................................................... 74 2.3.1.2. Differential scanning calorimetry (DSC) ................................................................. 75 2.3.2. Corrosion resistance ...................................................................................................... 75 2.3.3. Mechanical properties ................................................................................................... 76 2.3.4. Elastic properties ........................................................................................................... 76

Chapter summary ........................................................................................................................ 79 List of references .......................................................................................................................... 80

Chapter 3 Influence of Al and/Y incorporation to CrN coatings on their high temperature oxidation resistance

3.1. Morphological characterization of the CrN-based coatings ........................................... 82 3.2. Chemical characterization ................................................................................................... 87 3.3. Mechanical characterization ............................................................................................... 90 3.4. Structural characterization of as-deposited ceramic thin films ...................................... 91 3.5. Microstructural and crystallographic investigation: solid solution state ..................... 95 3.6. High temperature coatings durability: oxidation resistance .......................................... 98

3.6.1. Real time and in situ monitoring ................................................................................. 98 3.6.1. Real time oxidation observation by Environmental-SEM ..................................... 101 3.6.3. Kinetic approach of the oxidation ............................................................................. 104 3.6.4. Qualitative approach of oxidation behavior ............................................................ 111 3.6.4.1. Structural and chemical analysis after isothermal oxidation at 950 °C ............. 113 3.6.5. Microstructural investigation of the CrAlYN coating oxidized at 950 °C ........... 119

Chapter summary ...................................................................................................................... 125 List of references ........................................................................................................................ 126

Chapter 4 Influence of copper additions on amorphization of Zr-cu thin films metallic glasses

4.1. Influence of the deposition parameters ........................................................................... 131 4.1.1. Deposition rate ............................................................................................................. 131

4.2. Chemical composition of the thin films........................................................................... 132 4.2.1. Current intensity on the Cu target ............................................................................ 132

4.3. Structural analysis of prepared films ............................................................................... 135 4.4. Morphology of as-deposited films ................................................................................... 139 4.5. Microstructural characterization ...................................................................................... 145


4.6. Phases’ formation and their thermal stability ................................................................ 149 4.6.1. Thermal structural evolution of the glassy structure ............................................. 149

4.7. Thermal stability characterization .................................................................................... 160 4.7.1. Differential scanning calorimetry.............................................................................. 160 4.7.2. Comparison between DSC and XRD results ........................................................... 162

4.8. Oxidation effect during in situ XRD heating .................................................................. 163 4.9. Towards a model of crystallization for Zr-Cu TFMGs .................................................. 165 4.10. Mechanical properties ...................................................................................................... 169

4.10.1. Mechanical characteristics ........................................................................................ 169 4.10.2. Elastic characteristics ................................................................................................ 171 4.10.3. Mechanical behavior ................................................................................................. 176

4.11. How to improve durability of surface with Zr-Cu TFMGs? ...................................... 179 Chapter summary ...................................................................................................................... 182 List of references ........................................................................................................................ 183

Chapter 5 Innovative approach in surface engineering to combine high temperature resistance with corrosion protection

5.1 Are ceramic and metallic films compatible? .................................................................... 189 5.2 What films do we associate? .............................................................................................. 190 5.3 Towards an optimized physico-chemical protection through a duplex CrAlYN/Zr - 44 at% Ni TFMG ............................................................................................................................. 194 Chapter summary ...................................................................................................................... 198

Conclusions ........................................................................................................................ 200

Outlook and perspectives ........................................................................................ 202


iv

Introduction This thesis reviews the scientific contents and experimental results relative to PhD research activities carried out by the author at the MATEIS Laboratory, INSA de Lyon on coatings deposited at the IRTES-LERMPS Laboratory, UTBM. Surface engineering is and will remain a growing industry in the next decades, since it allows increasing the performance of the materials, controlling their surface properties independently of the bulk. Most of materials used for high technology applications can be considered as composites, with a combination of various and sometimes conflicting properties. Engineering components are, in everyday practices, exposed to severe industrial conditions such as corrosion media, elevated temperatures, wear and friction. Even more, high mechanical stresses are applied on those components used in aggressive environments and especially at high temperatures. For all above-mentioned examples, a modification of the surface is often the advanced solution able to increases the in-service durability by applying a thin surface film. Nowadays, a particular challenge for the industry is to develop technological processes for production of coated parts with specific shapes or large sizes based on environmental-friendly technologies. Hence the development of such know-hows and a good understanding of the oriented-properties for the deposition of coatings are the key points for an effective application of new advanced technologies with respect to a continuous improvement of the materials durability. In the last years, radical progress has been perceived in the field of coatings fabricated by the physical vapor deposition (PVD) technology and nowadays many solutions in terms of deposition processes are projected according to the sought application and targeted characteristics. More recently, new PVD methods, like for instance the HIPIMS process, are currently developed to produce dense, homogenous and defect-free layers with excellent adhesion to the substrate, properties required for the production of coatings the industrial scale. However, for the moment these new technologies are still really expensive and not much developed at the industrial level and we limited our investigations to more conventional well-proven PVD processes. Considering the previous studies involved in the research group of the MATEIS laboratory in the framework of hard coatings systems and oxidation resistance, this thesis is willing to still go ahead in the better understanding of coatings’ damage in order to optimize their performance.


General introduction

v

The PhD studies of A. Mege-Revil1 on TiN-based and T. Schmitt on CrN-based2 ceramic coatings served as support for our investigations, as it will be presented in the current project. Their main findings related to the degradation mechanism of hard coatings during exposure at high temperatures and under mechanical stress provided many valuable insights on the oxidation mechanism and structural transformation of hard coatings. Besides, the recent work of Dr. J. Qiao3 and Prof. J. M. Pelletier on dynamic mechanical relaxation and atomic mobility in bulk metallic glasses (BMGs) from the same MATEIS Laboratory, have inspired us to assess such materials in a form of thin film. Their research activities addressed various issues from the development of Zr-based BMGs and mechanical properties investigations to the study of the influence of thermal treatments and crystallization kinetics. Within this context, the main purpose of this PhD project is to investigate some new thin films and to propose them to protect and improve the durability and performance of bulk materials, characteristics requested in severe industrial operations. The hard coatings are extensively studied from the earlier 80’s. As an evidence, more than 100 000 articles are proposed when we ask the Elsevier database using “hard coating” as keywords On the contrary, the metallic glass in a thin film form is a new coming category of materials with only slightly above 3 000 papers fulfilling the “thin film metallic glass” keywords. Despite of a still immature activity level the interest on such innovative thin films will be likely growing fast in the next years, since the metallic glass thin films have already confirmed their promising properties. Here we propose a motivating approach related to two categories of thin films produced by two different PVD methods, namely arc evaporation for hard coatings and magnetron sputtering for the metallic thin films glasses. The leading line of the PhD is to analyze the impact of temperature on the films’ structure and their properties. In particular, the knowledge of the oxidation resistance and the crystallization behavior will be of great importance. By investigating the relationship between composition, (micro)structure, thermal stability and mechanical properties, this thesis aims to contribute to a better understanding of the mechanisms which govern the physico-chemical and mechanical evolution. Two main systems were studied, representative of advanced coatings developed for their enhanced durability. The CrN-like coatings was chosen as a basis for the ceramic coatings examined in this work because of its intrinsic ability to sustain modifications due to oxidation at high temperatures. The already well known CrN reference allowed us to further explore the physico-chemical relationship when an outside energy tends to alter the microstructure and structural stability of the coatings. Still, due to more and more severe

1 A. Mege-Revil, PhD thesis, INSA de Lyon, 2008, N 2008-ISAL-0079 2 T. Schmitt, PhD thesis, INSA de Lyon, 2010, N. 2010-39 3 J. Qiao, PhD thesis, INSA de Lyon, 2013, N. 2013-ISAL-0008


vi

industrial conditions the CrN coating cannot respond to the industry needs anymore and therefore new solutions had to be proposed. In the current project we propose to more specifically study the next generation of oxidation-resistant coatings: CrAlN on the one hand, CrAlYN on the other hand. Objective of this part is to analyze through a multiscale approach, the role of both Al and Y alloying elements on the protectiveness of coatings in relation with their structure. The second thin films system represents a new class of materials, the metallic glasses, characterized by the lack of crystal-like structure and by a good chemical homogeneity. These characteristics give rise to outstanding mechanical properties together with interesting corrosion protectiveness. Although bulk metallic glasses are nowadays among the most studied materials, their counterparts in the form of coatings are very recent and their synthesis and properties not well understood. Due to a high glass forming ability, Zr-Cu alloys have been chosen as an ideal model system, and our objective is then to assess the opportunity, from a mechanical and physico-chemical viewpoint, of depositing these alloys on a surface. With the purpose to link the functional properties of films with their chemical nature, many compositions scanning the whole composition’s domain were investigated. The PhD thesis will be organized into five chapters. The chapter 1 reviews the state of the art related to CrN-like and Zr-Cu metallic glasses thin films, gathering also their respective applications. The experimental details and characterization techniques employed during the course of this project are included in the second chapter, where the deposition details, technologies used and the conditions are briefly explained. The wide range of information provided by the experiments presented in chapter 2, such as the chemical composition, morphology, phase and structure determinations are gathered into the three results chapters as follows. Chapter 3 is dedicated to CrN-based thin films. In this chapter is analysed the influence of Al and/or Y additions to the CrN matrix on their high temperature oxidation resistance. The chapter presents the microstructural evolution and structural changes experienced by the coatings during exposure at high temperatures. Chapter 4 describes the special features of Zr-Cu thin films from a materials science and engineering viewpoint. The films deposited in a wide compositional range (13.4 - 98 at.% Cu) allowed a deep understanding of this binary system from a macro-structural until a local microstructural level. A special focus is given to the measurement of mechanical properties through conventional as well as spectroscopic techniques. The thermal and corrosion resistance of films is also investigated. After the definition of the best candidate of each category, chapter 5 presents an innovative solution able to improve both corrosion and oxidation protection of the coated substrates. Such a synergy could be acquired by a duplex ceramic / metallic glass architecture. Finally, the thesis is intended to present new trends, results and to have an overlook about future developments and applications where coatings could respond to specific industrial requirements.


Content

Part A. Hard coatings developed for high temperature applications ........................................... 2

1.1. Introduction .................................................................................................................................... 2

1.2. Short flash-back of PVD thin films’ development ..................................................................... 2

1.3. Conventional CrN thin films: general considerations and studies ......................................... 3

1.4. Effect of Al addition into CrN: chemical nature and properties improvement .................... 6

1.5. High temperature oxidation: nature of oxide scales ............................................................... 11

1.6. Interest of a further yttrium incorporation ............................................................................... 17

Part B. New class of thin films: metallic glass like ......................................................................... 23

1.7. Interest of metallic glass-based materials ................................................................................. 23

1.8. Bulk metallic glasses .................................................................................................................... 24

1.9. Thin film metallic glass (TFMG) ................................................................................................ 32

1.10. Advanced applications .............................................................................................................. 40

Chapter summary ............................................................................................................................... 44

List of references ................................................................................................................................. 46

Chapter 1

State of the art for the PVD ceramic

and metallic glass thin films

Overview

This chapter will present the state of the art related to the two different ceramic and

metallic glass PVD thin films investigated during the course of this PhD project.

The main characteristics in terms of correlations involving the deposition method,

microstructure, oriented-properties and their durability in relation to the functional

properties are presented.

The chapter is divided into two different sections devoted to the PVD coatings developed

within the current PhD. Part A is dedicated to the ceramic coating with a special focus put

on nitride-based films, while the part B deals with innovative thin films metallic glass.


State of the art for the PVD for ceramic and metallic glass thin films

2

Part A. Hard coatings developed for high temperature applications

1.1. Introduction

This chapter covers some of the most important studies available in the literature related to

the CrN-based PVD hard coatings. Since the aim of the current section is to better

understand the oxidation performance, the mentioned studies here are focused mainly on

oxidation resistance, mechanical properties and structural transformations that may

undergo during heating.

The section begins with a short flash-back regarding previous PVD ceramic coatings,

followed by the need of continuous development of such materials for industrial

applications with specific examples from binary (CrN, TiN ) to more complicated systems

(CrAlN, CrAlYN, TiCrAlN). The coatings performance is influenced by many parameters

such as deposition conditions (bias, substrate type) or the chemical nature of the films.

Namely, the CrN coatings show a limitation at high oxidation temperatures while the

addition of Al significantly improves the oxidation resistance and mechanical properties.

Further incorporation of yttrium into the Cr(Al)N-like matrix leads to a finer structure and

change oxides scale morphology ensuring an even better oxidation resistance.

1.2. Short flash-back of PVD thin films’ development

Among the wide choice of layers, ceramic-based coatings give new functionalities

improving the tools performance in addition to a longer lifetime [1, 2]. Up to now, many

types of PVD hard coatings are applied as protective surfaces to extend the performance

capabilities of materials, particularly for gears or machining tools as presented in figure 1.1.

Hard nitride-based coatings are exploited for cutting, hot forming or casting applications

where they are subjected to high temperatures environments, prone to induce

microstructural modifications that may affect their properties [3 - 7].

Figure 1.1. Section of perforated mild steel sheet and punch used to cold-punch the holes in the

sheet (a): Cr – Al – N coated punch (b) and Ti – N coated punch (c), Ti – N coated punch showing

areas of wear and coating oxidation (d). The Ti – N coated punch shows signs of wear and

damage to the nose, whereas the Cr – Al – N coated punch shows no sign of wear and minimal

damage to the nose [4]


Chapter 1

3

Since the commercialization of the first TiN in early 1980s [8 - 10], and later of ternary

TiAlNs [11 - 14], transition metal nitrides based hard coatings have been continuously

developed, so that innovative protective layers are each year proposed to respond to the

severe industrial requirements. These systems represent the pioneer PVD systems in the

field of hard thin films, being exploited mainly for high speed machining applications, as

wear resistance coatings able to improve the oxidation resistance, to extend lifetime and

performance of machining operations, leading finally to an enhanced productivity.

However, their characteristics at elevated temperatures are limited, which was corrected by

a further addition of aluminum into the film [14 - 15]. For instance, it is now accepted that

when operated under dry machining, edge of the cutting tool can reach temperatures higher

than 800°C, which represents a critical challenge for conventional Ti-based coatings. These

drawbacks have significantly limited their practical applications. Although, in order to be

able to overcome the stern operating conditions the development of new materials with

better performance which combine tribological and corrosion resistance properties was

desired. The second generation of more elaborated films was (and still are) then governed

by the CrN-based systems. It is known that CrN is more efficient than TiN in high

temperature environments (caused by the physico-chemical conditions themselves or

resulting from a strong friction) [16 - 18]. Hence, this ceramic coating is widely applied to

protect the tools from corrosion and oxidation in plastic molding operations and high speed

dry machining.

Furthermore, among the advantages of such system it could be mentioned the chemical

inertness, low internal stresses levels, high hardness degrees, good wear resistance and

thermal stability [19, 20]. More recently, CrN followed the same evolution as TiN, with a

further aluminum alloying with the purpose of still improving its thermal protectiveness

[16, 21-24]. Al addition indeed leads to the formation, at high temperature, of a dense and

well-adherent alumina scale onto the surface, able to prevent the inward diffusion of oxygen

deeper into the coating [15].

1.3. Conventional CrN thin films: general considerations and studies

Most of papers are devoted to studies on the microstructure [19, 20], mechanical properties

[25, 26] or thermal stability [27 - 29] of this coating. For instance, Mayrhofer et al. [20] stated

that the microstructure of the CrN coating is influenced by ion bombardment conditions

applied during the deposition and that mechanical and thermal properties depends a lot on

the compressive state.

In the study of Oden et al. [19] the bias applied during deposition (-400 and -200 V) is

identified as the key parameter able to control residual stresses level (Figure 1.2.). These

changes have been considered to be responsible for the increase of hardness together with

the decrease of the critical load determined by the scratch test.



4

Figure 1.2. Influence of the bias applied during deposition on the residual stresses present in

CrN coatings [7]

In a comparative manner, Ernst and colleagues investigated the thermal decomposition

routes of CrN thin films deposited by the most common PVD techniques, namely arc

evaporation on the one hand and magnetron sputtering on the other hand [28]. Arced CrN

coating exhibits more growth defects resulting inhomogeneity and droplets compared to the

smoother sputter deposited coating (Figure 1.3. a, b). Both processes give rise to dense

microstructures with a more pronounced columnar growth type for sputtered coatings

(Figure 1.3. c, d). The main decomposition path for CrN coatings comprises two major

reactions dependent on the nitrogen release: CrN is first transformed into hexagonal-Cr2N

then this compound is still reduced into metallic chromium. According to their results and

the data from other studies the first decomposition occurs at about 600 °C, whereas the

ultimate decomposition step is reported to have an onset temperature above 700 °C [17, 27,

29].

Figure 1.3. Example of typical SEM micrographs of CrN coatings synthetized by reactive arc

evaporation (a, c) and magnetron sputtering (b, d) deposited on Si [28]

As one of the transition metal nitride films, CrN has been widely studied for material

protection due to its chemical inertness and oxidation resistance. Several reports are

available in literature dealing with the oxidation behavior and kinetics of such hard coating

[14, 16 – 18, 30 - 32]. In most of studies, it was reported that the onset oxidation temperature

is around 600 °C and the oxidation rate is significantly accelerated from temperatures above

700 °C.


Chapter 1

5

This temperature corresponds to the first decomposition reaction with formation of Cr2O3

oxide. Once the oxide top-layer is formed, oxide scale growth is accelerated in relation with

the continuous nitrogen release occurring for temperatures of 800 – 900 °C until the final

reaction give rise to Cr phase [18]. It has been stated that the mass gain due to oxidation of

the CrN hard coatings followed a quantitative parabolic law, suggesting linked reactions’

kinetic limited by the diffusion of reactants to the oxidation front [18, 31, 33- 35] as

exemplified in figure 1.4.

Figure. 1.4. Example of parabolic oxidation kinetics for the modulated pulsed power (MPP)

magnetron sputtering CrN coating oxidized at 700, 800 and 900 °C [33]

Thermal resistance of CrN is believed to be mainly dependent on the transformation of

cubic CrN to hcp-Cr2N phase. As mentioned by J. Lin et al. [33] during heating the columnar

boundary bonding between the grains are weakened, resulting in a less dense structure

which allows enough space for oxygen atoms to diffuse from the atmosphere toward the

substrate. Hence, the chromium mobilities are also enhanced due to the presence of many

crystalline defects and the non-compact structure. These modifications will make the

opened boundaries along the columnar grains to be preferential diffusion paths for the

outward diffusion of the Cr and N species.

Several authors reported that the oxidation behavior of CrN coating is an oxygen dependent

process directly controlled by the N-release and the diffusion of reactive species. However,

the nature of the reactive species is still under discussion. Based on the findings of Chim et

al. [14] for instance both Cr and O atoms would react, whereas for others authors [30, [30]

the transition metals 𝐶𝑟3+ and O2- ions would be involved in the oxidation mechanisms. A

recent study of Qi et al. [18] compared the oxidation resistance of CrN and of Cr2N in air at

temperatures ranging from 700 °C to 1000 °C.

They concluded that the Cr2O3 scales growth was controlled by both the outward diffusion

of chromium and inward diffusion of oxygen. The oxide scale exhibited denser structure

with a slightly larger thickness of the Cr2N film compared to CrN one. Some other authors

revealed in their studies that the oxide scale formed on the coating surface may sometimes

offer a limited protection because of the further presence of substrate derived elements into

the film [31] [30]. The main reported influencing element is iron. Unfortunately, this key-

issue is discussed less, since most of papers are based on coatings deposited into inert iron-

free substrates (MgO, Si, SiO2…).



6

Mayrhofer et al. [31] point out also the importance of the oxide scale morphology in its

protective efficiency. They showed that during dynamic oxidation some cavities/pores are

formed beneath the coating / oxide interface, making easier the oxygen-nitride contact

(Figure 1.5.). Oxygen may afterwards diffuse deeply into the coating’s thickness, limiting

the protective effect of the oxide layer.

Figure 1.5. Cross sectional SEM fracture of the CrN coating after dynamic TGA up to 1000 °C

using a heating rate of 5 K min in an Ar/O2 atmosphere The figure shows the formation of

cavities after oxidation [31]

It has been established that CrN could be efficient to protect substrates against oxidation.

Nevertheless, its protectiveness may be altered owing to a poor density. Moreover,

oxidation resistance is closely linked to intrinsic thermal transformation of CrN, susceptible

to evolve towards Cr2N and even Cr.

The next step in the CrN-based coatings improvement thus led to a further alloying,

involving aluminum.

1.4. Effect of Al addition into CrN: chemical nature and properties improvement

Extensively used in the last decades [14, 16, 17, 30, 31], the binary films seem to have

reached a certain limit to fulfill the new industrial needs (especially for applications at high

temperatures). Once exposed at temperatures higher than 700 °C, the mechanical properties

are quickly degraded owing to the formation of porous oxide at the film surface [17].

Recently, some solutions have been proposed by adding aluminium into the faced-centered

cubic (fcc) structure of CrN resulting in the formation of a metastable ternary CrAlN solid

solution [36 - 40]. The CrAlN coating has been reported to be deposit by several methods

such as magnetron sputtering [41 - 44] or arc evaporation [45 - 48]. Systems with different

phases are obtained depending on the aluminum content, as reported by Reiter et al. [49].

Below an Al content ~ 77 at.% [49 - 52] the AlN forms generally a solid solution in CrN

lattice. The dominant phase is indeed dependent on the aluminum content [49]. Al atoms

could substitute the Cr atoms in the fcc-CrN lattice for contents of 46 at.% Al, whereas for

contents of 71 at.%, Cr atoms substitute aluminum atoms in the fcc-AlN lattice. The

structure obtained above the theoretically determined limit (~ 77 at.%) is the hexagonal type.

The phase transition for AlN is not accurately identified, varying between 60 and 80 at.%

[50, 53].


Chapter 1

7

From a structural viewpoint, films are characterized by a strong {111} preferred orientation

[36, 39]. These ternary systems have been reported to exhibits improved hardness [24, 54, 55]

as well as oxidation resistance [15, 22, 23, 49, 56 - 59] compared with Al-free CrN or TiN.

Depending on the Cr content the reported values for hardness are in the range of 25 – 40

GPa [4, 24, 36, 38, 54] or sometime lower (18.8 GPa) as stated in the study of Shah and

Jayaganthan [39]. Le Bourhis et al. [24] made a report on some commercially available hard

CrAlN-based coatings and compared their hardness and residual stresses with the available

published work in the literature. Authors concluded that both the grain size and the

residual stresses of coatings are key-parameters linked to their density and mechanical

properties.

The effect of an aluminum addition in binary systems on different properties (wear

resistance, oxidation resistance, hardness….) has widely been studied. Results obtained are

summarizing in table 1.1. From all these results, a beneficial effect of the addition of Al was

highlighted.

Table 1.1. Examples of direct comparison between different types of coatings

Systems

compared

Oxidation

resistance

Wear

resistance Other properties

Substrate/

method Ref.

CrAlN,

TiAlN vs.

CrN, TiN

CrAlN,

TiAlN are

better

CrAlN,

TiAlN

are better

- CrAlN, TiAlN:

higher H/E’

- better wet drilling

test for CrAlN vs. TiN and

TiAlN

AISI H13

steel,

magnetron

sputtering

[37]

TiN,

TiAlN,

AlTiN and

CrAlN

- oxidation

resistance

of TiAlN

due to the

presence of

Al leads to

the lowest

wear rate

- lower wear

rate for TiN

and TiAlN

- high Al

content for

AlTiN and

CrAlN don't

present better

anti-wear

properties

than TiN and

TiAlN

- better

properties of

anti-spalling

and anti-

adhesion for

CrAlN than

TiAlN and

AlTiN

- Friction coefficients are at the

same level (around 0.7) and

can

be ranged as followed

AlTiN > CrAlN > TiAlN > TiN

AISI H13

steel,

magnetron

sputtering

[38]

AlCrN vs.

CrN

Better wear

resistance for

AlCrN

Better load-carrying capacity

for AlCrN

AISI 304 SS,

magnetron

sputtering

[39]



8

The comparison between the coatings of L. Aihua et al. [12] in terms of friction coefficients

are presented in figure 1.6.

Figure 1.6. Friction coefficient as function of number of cycles for different coatings [12]

The better tribological behavior of AlCrN coating is interpreted in terms of wear debris

emission and of load-carrying capacity as reported by Mo et al. [60] in figure 1.7.

Figure 1.7. Wear scar morphologies of the coatings when sliding against Si3N4 ceramic under

normal load of 10 N for CrN (a) and CrAlN (b) [60]

CrAlN has been of much interest and many studies presented its microstructure [36, 38, 39,

61, 62], mechanical and tribological properties [12, 24, 33, 37, 55, 60, 63], electrochemical

behavior [37, 40] or thermal stability [49, 52, 64, 65]. Even though, few studies [45, 66, 67]

present the influence of the substrate bias voltage on the deposited coatings, the final

microstructure obtained and their properties. This deposition parameter is determinant for

the growth rate, surface roughness and the composition of the thin films. Romero et al. [45]

highlighted that as the negative bias voltage was increased from – 50V to -400 V the growth

rate and the preferred orientation from {111} to {200} were changed.

The same observation is presented in the study of Y. Lv et al. [67] where the preferred

orientation of the film changed from {220} to {200} accompanied with the FWHM of the peak

widened as the bias voltage increases.


Chapter 1

9

The general trend in all the studies indicates that the hardness change as the negative bias

voltage increases. All authors agreed that the increase of the substrate bias voltage leads to:

- a decrease of the growth rate leading to a more compact and denser coating with small

grains and higher hardness,

- an improvement of the corrosion resistance.

Among the mechanical and tribological properties, Sanchez et al. [37] also investigated the

electrochemical behavior of Cr1-xAlxN coating. The behavior of the protective effect of this

coating is based on the substitution of Cr for Al when the power applied to aluminum target

increases.

Polarization resistance values decreased when Al content in CrAlN increased, as suggested

by their results presented in figure 1.8.

Figure 1.8. The Nyquist impedance diagrams (a) and Tafel curves (b) of Cr1-xAlxN coatings

grown on steel, as a function of the Al content (x values) [37]

The corrosion rate increases also as more Al is integrated in the matrix. The best corrosion

resistance was achieved for a 0.54 at.% Al content, which would correspond to a fine-

grained structure and a low surface roughness. Besides, for this aluminum enrichment in

the lattice, the study reports the highest hardness (30 GPa), lowest friction coefficient (0.45)

and a very high reduced elastic modulus (303 GPa) amongst all the investigated

composition (0.51 < x < 0.69).

Moreover, Barshilia et al. [16] proposed that during the chemical attack of Al, an Al2O3 layer

is formed on the surface of the coating. The exposed surface becomes passivated, which

hinders any further oxygen diffusion (corrosive species) through the pores inside the film.

The structural transformations of the Cr-based coatings have to be considered for a better

understanding of the oxidation process. The thermal stability of the CrAlN coating rules the

oxidation resistance and hence these structural modifications must be mentioned and

properly explained.



10

Many authors studied also the thermal stability of such kind of coatings as will be presented

in what follows. Reiter et al. [49] represented schematically a phase stability diagram of

AlCrN films for different Al contents, as shown on figure 1.9 a.

The XRD patterns presented by Reiter et al. [49] and Willmann et al. [52] showed that

initially the AlCrN is in its cubic form with the fcc-peaks located between the ones of fcc-

CrN and fcc-AlN. This structure is preserved up to 900 °C as it can be seen from figure 1.9.a.

The fcc-AlCrN transformed 100 °C higher where hcp-AlCrN and Cr2N phases are detected.

For some authors, the reduction of nitrogen may occur somehow before the precipitation of

AlN and the Cr2N transformation would be the first initiated reaction [17, 22].

The segregation of fcc-AlCrN to hcp-AlCrN leads to a gradual chromium enrichment of the

CrAlN matrix that promotes the formation of the Cr2N. For the CrN coating this

transformation takes place generally when the nitrogen is released during the heating above

800 °C.

For higher temperatures (1200 – 1300 °C) the spinodal decomposition occurs, expressed by

the concomitant appearance of hexagonal AlN combined with the metallic chromium phase

(Figure 1.9. b).

Figure 1.9. Structural stability of various AlCrN coatings determined by XRD spectra (a) [49] and

phases determined by XRD analysis for a Al0.7Cr0.3N film annealed at different temperatures [52]


Chapter 1

11

1.5. High temperature oxidation: nature of oxide scales

The oxidation protection and behavior at high temperatures of a thin nitride layer, such as

for instance TiN or CrN, revealed great potential when used to protect cutting tools in

different industries [68]. It is well recognized that CrAlN coating is a good candidate as a

better alternative to conventional CrN coatings especially for high temperature oxidation

resistance applications. The first studies of the oxidation performance, as presented on table

1.2., compared several nitride coatings.

Table 1.2. Example of direct comparison for different Al-containing films

Systems Oxidation resistance Substrate/method Ref.

CrN vs AlCrN Al addition improved the oxidation

resistance

glass, 304 SS, reactive

magnetron [17]

TiN vs CrN vs

TiAlN vs

CrAlN

with Al incorporation TiAlN and

CrAlN showed better oxidation

resistance

SS, vacuum arc deposition

[14]

CrAlN vs

CrAlSiN

vs AlCrSiN

CrAlN and AlCrSiN showed better

oxidation resistance FeCrAl alloy, arc evaporation [22]

TiAlN vs

CrAlN

Oxidation resistance improved with

the increase of Al content 304 SS, AIP method [15]

AlCrSiWN vs

AlTiN

vs AlCrN

AlCrSiWN > AlCrN > AlTiN WC– 1 2Co– 0. 9 5Cr3C2– 0. 3

5VC, arc evaporation [59]

TiAlN vs

AlCrN

TiAlN

> uncoated > AlCrN

2.25Cr-1Mo steel, arc

evaporation [13]

Focusing on the Al effect, some studies were dedicated to the oxidation performance of

AlCrN containing different Al contents [58] [56] [23] [57] [59] [58] [23]. Both the onset and

critical oxidation temperatures are shifted to higher values and the oxidation kinetics are

delayed.

For the CrN coating, the oxidation phenomenon is linked to the structural transformation

into the hexagonal Cr2N phase and formation of the Cr2O3 reaction [18]. J. Lin et al. [33]

confirmed this structural/physico-chemical correspondence, and identified the formation of

the Cr2N phase above 800 °C.

They determined that the matrix is gradually enriched in Cr due to the nitrogen release in

subsurface, which will promote the formation of the Cr2O3 (detected at the same

temperature) as shown in figure 1.10.



12

Figure 1.10. XRD patterns of Cr0.40Al0.60N films before and after air annealing at different

temperatures for 1h

In their thermal evolution mechanism, they point out the fact that alumina is preferentially

formed with respect to chromia, owing to its more favorable thermodynamic data (Gibbs

free energy for Al2O3 and Cr2O3 of - 378.2 kcal/mol and - 252.9 kcal/mol respectively). The

alumina usually forms earlier and more easily as revealed by the analysis from ref. [17]

where a mixture of both Al2O3/Cr2O3 oxidized phases was detected by XPS, XRD and EDS

results.

Based on these results, it is assumed that the alumina is formed at temperatures as early as

650 – 700 °C and can maintain its amorphous state until 900 °C. The presence of the oxide

layer on the top surface of the film was mentioned as well in ref. [22], and may contain both

chromium and aluminum as highlighted by figure 1.11.

Figure 1.11. Cross section of CrAlN film annealed to 1300 °C by SEM. Arrows show localized

areas analyzed by EDS [22]


Chapter 1

13

For Lin and coworkers, the role of Al addition to CrN on the improvement of the oxidation

resistance can find two origins [17]. The first explanation concerns the bonding’s nature

since the Al will occupy the Cr atomic sites and form the fcc-CrAlN metastable structure.

Authors stated that the fcc-AlN involves strong and covalent bondings, while CrN bondings

would rather present a metallic character. This could inhibit or delay the decomposition of

fcc-CrAlN with nitrogen loss because of the high bonding energy between Al and N atoms.

Hence, the Al atoms existing in the CrAlN lattice to bond with nitrogen atoms could

improve the thermal stability of the coating.

The covalent character of fcc-AlN and its beneficial effect on the films’ thermal stability is

also highlighted by Hamseh et al. [57]. Their XRD results showed that the film retained a

combination of hcp-AlN and hcp-Cr2N structure up to 800 C, temperature at which the

formation of Al2O3 and Cr2O3 oxides is initiated.

The second explanation given by Lin et al. is related to the formation of the dense oxide

layer on the films surface that could act as an effective diffusion barrier, slowing down the

inward diffusion of the oxygen at high temperatures. The dense protective layer is a mixture

scale (either amorphous or crystalline) with a more dense structure than the single

crystalline Cr2O3 for the CrN coating. Lastly, it is expected that the main parameter for the

oxidation of CrAlN thin film is the main phase structure change, namely precipitation of the

hexagonal AlN phase with formation of the Cr2N. Such spinodal decomposition would

correspond to the beginning of oxidation process.

On the other hand, Zhu et al. [69] reported the formation of a “solid oxide solution” (Cr,

Al)2O3 with pure Cr2O3 grains dispersed in the outer oxide layer (figure 1.12.) for AlCrN

with different Al/Cr atomic ratios. According to their findings the mixed (Cr,Al)2O3 scale is

formed predominately by the inward diffusion of oxygen without cations diffusion during

an oxidation process as follows. In the first steps of oxidation, the oxygen reacts with the

coating and form Cr2O3 and Al2O3 oxides on the surface. Once this thin oxide scale is

formed, the oxygen reaches the oxide scale/un-reacted coating to the grain boundaries.

The high affinity of Al for oxygen endorses the alumina to form rather first, while some Cr

can be also oxidized and develop Cr2O3.

The structural resemblances (corundum crystal structure) with a hexagonal symmetry

between the two type of oxides growth, have led to the final oxide formed to be a solid

solution (Cr,Al)2O3. Nitrogen is released during heating through some short diffusion paths

in the oxide scale. Some parts of Cr species diffused as well through the paths, being

oxidized at the outer layer of the oxide scale and formed pure Cr2O3 grains.

Authors explained that when the aluminum content in the coating is low, chromium may

diffuse easily through the relatively thin Al2O3 layer, developing Cr2O3 grains, which may

enrich the outer part of the oxide scale in chromium (Figure 1.12. a).

In the opposite situation, for Cr0.63Al0.37N coating the (Cr,Al)2O3 formed layer contained a

relatively higher amount of alumina, which makes more difficult the outward diffusion of

Cr.



14

Figure 1.12. Cross-section morphologies of three coatings (a) Cr0.82Al0.18N, (b) Cr0.74Al0.26N and (c)

Cr0.63Al0.37N after an oxidation of 20 hours at 1000 °C [69]

Banakh et al. [56] studied also the oxidation resistance of the CrAlN coating for a wide range

of Al contents (below the solubility limit) and they also corroborated the formation of a

mixed (Cr,Al)2O3 phase on the surface of the CrAlN. Their findings lead to an oxidation

model: from the room temperature CrAlN film incorporates some oxygen, the process being

favored by the aluminum alloying. When heated, the “real oxidation” carries on by the “de-

nitridation” of the structure (nitrogen release) with the beginning of the mixed Cr/Al oxide

phases’ formation. When the aluminum is present in high amount, a thin oxide fully

oxidized surface layer is developed, thereby acts as a diffusion barrier for oxygen, limiting

and slowing down further oxidation of the deeper nitride coating.

Other studies consider the diffusion of both Cr and Al cations responsible for the oxide

development [14, 17, 49, 67, 70]. Zhu and colleagues [69] consider that as the Al content

increases, the oxide scale becomes denser and may hinder the species flux (oxygen inward

as well as metallic cations outward diffusion).

In other studies, Galindo and coworkers [58] proposed an original approach involving

further intermediate TiN layers to limit the chromium outward diffusion in the case of

AlCrN oxidation. The outward diffusion of Cr is thereby decreased, which will promote the

formation of the protective scale mainly composed of alumina rather than chromium oxide.

The oxygen inward diffusion and migration of the substrate elements to the surface are

inhibited, trapped by the outer TiN layer. Similar experiments carried out by Endrino et al.

[23] with a TiN deposited sublayer at the surface and at the substrate showed improved

oxidation resistance for the AlCrN.

Figure 1.13. shows that TiN layer successfully delayed the Cr diffusion by forming the

alumina faster.


Chapter 1

15

Figure 1.13. Example of glow discharge optical emission spectroscopy depth-profiles of AlCrN

film oxidized at 900 °C (TiN layer near the surface) for 3h: a) general scan and b) oxidized layer

[23]

The GDOES profile indicates that the outward diffusion of Cr to the outer surface begins

below TiN layer (Figure 1.13. b). Once oxidized the depth profile shows the formation of a

50 nm Al-rich oxide film on the most outer surface. This demonstrates that the supply of Cr

atoms to the surface is limited and TiN proves its effectiveness as a diffusion barrier for Cr

within the AlCrN coating.

Finally, Reiter and coworkers [48] reported a systematic investigation of the oxidation

behavior of arc-evaporated AlCrN films, with Al compositions varying between 0.29 and

0.79 at.%. They analyzed the impact of heat treatments in inert atmosphere (Ar) and air on

both the structure and oxidation of the Al1-xCrxN coatings. When annealed in Ar they found

that Cr segregates to the films’ surface, followed by the decomposition of the ternary coating

into hcp-AlN and, due to the nitrogen release, to Cr2N and later Cr. In contrary, during

heating in oxidative conditions, a mixed Cr2O3-Al2O3 layer is formed due to inward/outward

diffusion of oxygen and Cr, Al, respectively. Growth of this oxide depends mainly on the

capacity of the nitrogen to be released. The last remark was that the oxide scale formed has a

protective effect against the decomposition of the metastable fcc-AlCrN phase during high

temperature exposure.

In figure 1.14. it can be observed a SIMS depth profile for the Al0.71Cr0.29N coating annealed

for 30 min in air at various temperatures. The compositional profiles show that Cr and Al

display a strong segregation as its content decreases down to a minimum and then increases

to its maximum close to the surface. The oxygen content at the coating surface was around

60 at. % and decreases in function of the annealing temperature. In the outer region, the

inward diffusion of oxygen and nitrogen depletion is clearly observed, associated with the

chromium external local enrichment.

The compositional variation in the region between Almax and Crmax together with the low N

concentrations and O contents could indicate the formation of a mixed oxide layer Cr2O3-

Al2O3 with a Cr enriched area due to more favorable outward diffusion of chromium.



16

Figure 1.14. Chemical SIMS depth profiles for Al0.71Cr0.29N annealed for 30 min in air at various

temperatures: (a) Cr, (b) Al, (c) N and (d) O chemical profiles [48]

Certain studies showed that among the ceramic materials crystallized into the AlN based

cubic structure which belongs to Fm-3m (225) space group, the AlCrSiN-based coatings [22]

[59] exhibits high hardness and good oxidation resistance. It can be noted that such a

beneficial effect of Si was already reported for more simple systems such as TiN [71] and

CrN [72].

To summarize and from a quantitative point of view, the incorporation of Al to CrN has a

significant influence on the oxidation behavior, more particular in the [600, 1200 °C]

temperature range. The onset temperature of the film oxidation indeed can be increased

from 600 to 650 °C and the temperature of the significant nitrogen loss is shifted from 800 °C

to 1000 °C. Moreover, the temperature at which the oxidation reaction becomes catastrophic

is changed from 1040 to 1180 °C [17].

Regarding the protective mechanism, it can be stated that the oxidation process of the

AlCrN coating is mainly controlled by the outward diffusion of Cr, Al and N species and by

the relative inwards diffusion of oxygen. Al2O3 formation is thermodynamically favored but

from a kinetic viewpoint Cr ions would be more mobile. Such compromised conditions

might explain the formation of the mixed (Al, Cr)2O3 compound, sometimes reported.


Chapter 1

17

1.6. Interest of a further yttrium incorporation

In the previous investigations on ternary CrAlN it was mentioned that thermal stability and

oxidation behavior were predominantly influenced by the transport of species to the

oxidation front.

More recently, some studies reported that the oxidation resistance of the metastable

(Ti,Cr)AlN coatings can be further improved by incorporating foreign elements such as Y or

Zr [73]. Main studies are focused on the influence of Y addition on the oxidation resistance

[68, 74 - 76], thermal stability [77, 78] and tribological properties [79, 81] and combinations of

those.

The reason for doping nitrides coatings with yttrium would be linked to its ability to

segregate on grain boundaries, susceptible to limit the mobility of species and to form faster

the stable α-alumina phase.

As usual, all attempts on the evolution of films were first tested to TiN-based coatings. It is

also the case to estimate the Y-doping efficiency.

Effect of an Y addition to TiN-based coatings

L. Zhu et al. [82] studied the effects of yttrium on the cathodic arc evaporated TiAlN coating

with emphasize on the structure and oxidation resistance. If the addition of 1 at.% Y does

not modify the NaCl-type structure, the resulting microstructure is however significantly

refined (grains of 12 nm for Y-containing film vs 20 nm for TiAlN).

The Y incorporation has the following effects:

- it promotes the formation of Y2O3 and Al2O3, which restrain the diffusion of the oxygen

and hence protect the coating from further oxidation;

- it hinders the grain boundary outward transport of Ti and Al cations. Y segregates faster to

the grain boundaries and plugs the diffusion paths that help to enhance the oxidation

resistance;

- Y incorporation in the cubic lattice leads to a grain refinement (12 nm for Y-containing film

and 20 nm for TiAlN) which shrinkages the distance to the grain boundaries

- the yttrium is also effective to improve the interfacial adhesion of the deposited films as

shown in fractured cross-section morphologies of the investigated films after oxidation at

900 °C for 210 min as presented in figure 1.15.

In the case of the TiAlN when exposed to an oxidative environment, the oxide formed inside

the film increases its volume leading to cracks as observed from figure 1.15.a., whereas for

the TiAlYN such phenomena is not seen. The better oxidation resistance of TiAlYN coating

is attributed therefore to the delay on growing the oxide scale characterized also by a fine-

grained structure.



18

Figure 1.15. SEM fractured cross section of TiAlN (a) and TiAlYN (b) after a 3.5 hour oxidation at

900 °C [82]

In more complicated system, namely (Ti,Cr,Al)N coating, Yamamoto et al. [83] performed

oxidation tests with the aim to understand the effect of Si+Y additions on the structure and

on mechanical properties. In their work it is mentioned only that the Y-containing film has

an improved oxidation resistance with no further explanations given. Another important

remark is linked to the presence of Al rich-oxide layer with needle-like structures in the

outermost part of the coating but again without any clear explanation about the involved

mechanism of the oxide formation.

However, interpretation of results is difficult because both Si and Y may influence the

oxidation process. A TEM cross section analysis indicates a Si enrichment corresponding to

a silicon-rich sublayer, prone to block species diffusion fluxes (Figure 1.16.). This elemental

chemical profile analysis also shows that Y is uniformly distributed throughout the layer.

Figure 1.16. Cross-sectional TEM images of near surface area of (TiCrAlSiY)N after the

isothermal oxidation test at 1000 °C for 30 min (a), detail of the frame on figure (b). EDX line

scan from point A to B (c), and electron diffraction pattern of needle like oxide on the surface (d)

[83]


Chapter 1

19

Given the good results obtained on the oxidation performance on TiAlN coatings when

doped by yttrium, the same kind of improvement was expected for CrAlN.

Effect of an Y addition to CrN-based coatings

Rojas et al. [73] explored effects of the dopant elements (2 at.% Zr or Y) on the microstructure

and the physico-chemical behavior of the CrAlN coating. Both coatings (CrAlZrN and

CrAlYN) develop on their surface an oxide top layer rich in Cr and Al oxides. In the case of

CrAlZrN the oxide scale is thicker, inferring that oxygen may easily progress with respect to

the other Y-containing film. A mechanism is suggested where Y atoms may diffuse out

concomitantly with Cr and Al elements forming a more efficient protective oxide layer,

whereas the mobility of Zr mobility is slowed down so that metallic atoms (or ions) remain

concentrated in the unaltered film structure. This can be observed from the EDS profiles

presented figure 1.17.

Figure 1.17. EDS spectra obtained from the topmost layer and inner part of the CrAlZrN (a) and

CrAlYN (b) coatings. Below are presented line scan profiles obtained across the CrAlZrN (c) and

CrAlYN (d) coatings [73]

Several authors have shown that the incorporation of large atoms such yttrium can

effectively delay the diffusion process resulting in a better thermal stability, improved

mechanical properties and oxidation resistance at high temperatures [68, 73 - 76].

Besides, yttrium addition is also motivated by its propensity to improve the adherence of

the oxide scale. Growth mechanisms of the oxide scale are modified by the slowed down

oxygen inward diffusion at the scale/gas interface along the scale boundaries.



20

Rovere et al. [70] studied the effect of yttrium (contents of 0, 1, 2 and 4 at.% ) on the

oxidation resistance of magnetron sputtering CrAlN under isothermal and dynamic

conditions up to 1400 °C. Coatings enriched with 1 and 2 at.% Y showed significantly lower

weight gains during isothermal oxidation denoting its beneficial protective effect regarding

the oxidation resistance owing to the formation of a dense and protective mixed α-

(Al,Cr)2O3 scale able to retard further oxidation.

The excellent oxidation protection of the Cr0.45Al0.53Y0.02N has been explained at a fine scale

considering the SEM (Figure 1.18) and TEM (Figure 1.19) results.

Figure 1.18. SEM Morphology of Cr0.45Al0.53Y0.02N film after oxidation at 1060 °C (a) and SEM-

EDX line scan profile (b) [70]

Figure 1.19. General view (a) and detail of TEM cross section (b) of Cr0.45Al0.53Y0.02N film after

oxidation at 1000 °C. The SAED pattern of the detail (c) and STEM-EDX line scan across oxide

scale (d) [70]

It is suggested that Y favors the preferential oxidation of stable Al2O3 over Cr2O3 due to

mobility differences between Cr and Al cations. The Y doping favors the formation of a

duplex oxide α-(Al,Cr)2O3 with an Al rich outer oxide layer on top of a Cr-rich oxinitride

layer which is crystalline and dense.


Chapter 1

21

The oxide scale provides outstanding protection of the unoxidized film before the beginning

of the diffusion onset (initiated by the nitrogen released).

Very recently, Qi et al. [84] have investigated as well the influence of the yttrium contents (0

to 2.3 at.%) on the hardness and oxidation resistance of the magnetron CrAlN coatings.

Compared to the previous report of Rovere and coworkers [70], they stated that the yttrium

has a beneficial effect for improving the oxidation resistance of CrAlN between 0.3 and 0.7

at.%.

Their results showed that Y atoms substitute Cr and/or Al atoms in the CrAlN lattice,

forming a metastable solid solution CrAlYN with a single fcc phase. The hardness increased

as more yttrium was added in the matrix due to the solid solution strengthening as

suggested by figure 1.20.

Figure 1.20. Evolution of average hardness for CrAlYN coatings as a function of Y content [84]

The oxide scales formed after oxidation at 1100 °C were thinner for the coatings containing

0.3 and 0.7 at.% Y, implying again a higher oxidation resistance that the one of Y-free CrAlN

coating (Figure 1.21).

Figure 1.21. The SEM cross-sectional micrographs of CrAlYN oxidized 2 hours in air at 1100 °C

[84]



22

As in the study of Rovere et al. [70] the formed oxide layer was a dense protective α-

(Al,Cr)2O3 type. For higher Y contents (> 1.3 at.%), a Y-Al based is formed. Growth of this

thick layer causes the rupture of the diffusion barrier, leading to a porous morphology and

poorer protectiveness.

The yttrium effect on oxidation resistance when it is incorporated into nitride-based

coatings is still under research, since as presented above, only few studies are dealing with

this motivating issue. Although many of the studies tried to optimize its content, its

efficiency and scientific role are less highlighted, so that further deeper investigations are

mandatory.

In the current thesis one of the main objectives related to ceramic coatings will be to better

understand the role of yttrium with respect to microstructural, chemical and structural

states during oxidation at high temperatures.


Chapter 1

23

Part B. New class of thin films: metallic glass like

1.7. Interest of metallic glass-based materials

The metallic glasses represent a frontier of the metal research, being today among the most-

actively studied materials since they are regarded as a model and likely future newcomers

for engineering applications. Due to the lack of dislocation-like crystalline defects and no

long-range structural periodicity, the bulk metallic glasses (BMGs) display an attractive

combination of physical and chemical properties such as amorphicity, extraordinary high

strength, soft magnetic characteristics and wear/corrosion resistance. Besides, the BMGs

have extraordinary thermal formability and superelasticity, comparable to those of

polymers and window glasses. The glass-forming materials with high thermal stability and

low critical cooling rates have attracted great attention for both fundamental and applied

research.

The initial interest and purpose for developing metallic glass materials has been for

industrial applications, owing to their excellent mechanical properties [85].

The efforts devoted to the continuous development and research of BMGs made possible

vapor-solid quenching during sputtering. Nowadays fully thin amorphous alloys called thin

film metallic glasses (TFMGs) are deposited by several physical vapor deposition

techniques. The co-sputtering by two or more targets seems to be an encouraging method in

exploring the composition, microstructure and orientated-properties of TFMGs.

The excellent adherence and the smooth film surface could for instance, have a noticeable

influence on the fatigue life [86]. These outstanding features can be exploited for

applications in nanotechnology, bio-nanotechnology and MEMS. The sub-nanoscale

structural homogeneity could be also of interest for optical applications [87 - 90].

This section covers the general properties of BMGs, based on their specific characteristics,

and recent progress in the development of thin metallic glass films with emphasis on the

binary Zr-Cu. The chapter begins with the description of the basic concepts related to the

formation and design of BMGs, followed by the influence of structural stability on their

specific properties. The second part is related more to the TFMGs, focusing on how to obtain

metallic glass thin films with examples of several multicomponent systems.

In the final part, is presented the binary TFMGs, particularly the Zr-Cu system investigated

during the course of this work. The state of art chapter ends by illustrating some of the

potential applications for such metallic glass-based materials.



24

1.8. Bulk metallic glasses

1.8.1. Basic concepts regarding bulk metallic glasses

In a general definition, metallic glasses can be defined as materials with crystal-free

structure and lack of long range atomic periodicity capable to vitrify when cooled at a

sufficient high rate (see figure 1.22.). The above definition nevertheless considers only the

structural viewpoint and it should be noted that the glasses may chemically have different

types of bonding (i.e. metallic, ionic or covalent). Among the extensive family of glasses, the

bulk metallic glasses are probably the most remarkable, owing to number of outstanding

characteristics such as extreme strength, high flexibility or amorphicity [87].

Besides, one of the most important features that distinguish them from general amorphous

materials is the glass transition that may transform supercooled liquids into a glassy state

when cooled.

On the other hand, it is acknowledged that the atomic configuration of molten metals and

alloys are generally disordered. When the liquids can be quenched at a sufficient high

cooling rate to suppress the formation of equilibrium crystalline phases, the non-crystalline

structure is expected to be retained up to a critical cooling rate of the order of 105 – 106 K s-1.

Metallic glasses could be therefore scientifically defined as amorphous alloys that exhibit

glass transition temperature (Tg), from which derive their unique properties of high viscous

flow ability and fracture toughness, along with physical and structural properties that

change abruptly at the Tg.

Figure 1.22. Metals typical atomic structure organization (body centered cubic arrangement) (a)

and disorder structure of a metallic glass (Cu35Zr65) (b) [91]

First metallic glasses were synthesized by Kramer in 1934 and Buckel et al. [92] in 1950 by

vapor deposition method. However, the quenched Au75Si25 discovered by Duwez and co-

workers in 1960 are generally recognized as the first metallic glass system vitrified by rapid

cooling from the molten state [93]. After this seminal discovery, there has been a systematic

research for metallic glass-forming compositions, seeking elemental combinations with

lower cooling rates for the preservation of the amorphous structure. From the Pd-Ni-P

system developed by the Turnbull’s group [94] in the early 1980s, to La-Al-Ni or La-Al-Cu

alloys studied by the Inoue’s group in late 1980s [95], the development of metallic glasses

has enabled to produce new structural materials with inimitable properties.

All these initial studies pave the way for finding more bulk metallic glasses with high glass

forming ability and thermal stability such as Mg-Cu-Y or Zr-, Fe-, Cu-, Pd-based [87, 96].


Chapter 1

25

Figure 1.23. illustrates few examples of first metallic glass materials considering their critical

casting thickness. Following the discovery of first metallic glasses, a large number of

systems were explored.

Invention of the melt spinning method as a rapid cooling technique [r2] made possible

quenching liquid alloys as 10 – 50 µm thick ribbons, which favored even more the

advancement in the metallic glass research.

Figure 1.23. Some of the first metallic glasses as a function of year and critical casting thickness

(a) [96] and Zr or Mg based BMGs rods and plates in various dimensions (b) [89]

The excellent glass former systems contain largely only transition metals, such as the Zr-Cu-

Al-based glasses [97, 98] or Zr-Al-Ni-Cu [99 - 101]. Peker and Johnson reported an excellent

multicomponent Zr41.2Ti13.8Cu12.5Ni10Be22.5 (at.%) BMG containing beryllium, the second

smallest metallic atom (0.111 nm) [102].

The different atomic sizes in this system are expected to limit the solubility in crystalline

phases, and to fill the empty space in the defective glass structure by stabilizing it more

efficiently.

To date, this is still one of the best glass formers that exemplify the compositional effect in a

multicomponent system able to ensure the densest structural packing in zirconium-based

BMGs. These research achievements have endorsed as well the progress in terms of casting

technology [87, 103, 104] and stimulated enthusiasm for metallic glass research.

Various alloys with excellent glass formation in iron-, nickel-, titanium-, and copper based

alloy systems are known today to vitrify upon cooling [85, 87, 89, 95, 96, 105].



26

1.8.2. Formation and design of BMGs

The general method for developing BMGs with desired properties is first to select a single,

proper, base component and then to match it to other glass-forming constituents.

For instance, for binary and ternary alloy systems low eutectic points have been widely used

as an indication for seeking good glass formers. It is quite difficult though to represent more

than three components on a phase diagram since the eutectic transitions in quaternary or

more complicated systems are rarely experimentally known. After the base composition is

found, some minor additions of a third/fourth element could enhance the glass forming

ability and plasticity of the metallic glass materials. For example, the addition of minor Cd

to a Mg-based bulk metallic glass increases the ΔTx (supercooled liquid region) to a

maximum of 70 K [106]. Most of the multicomponent metallic glasses have been discovered

by the empiric rule “mix and see” and the development of new BMGs requires considerable

experience. The ideal contents for the additions are basically found to be in a very-narrow

composition range. The selection of suitable additional materials and the method of

pinpointing the optimal fraction of the additional materials are thus quite important.

Several criteria have been proposed to predict the glass forming ability of alloys, based on

their thermodynamic parameters, coupled with kinetics of nucleation and/or growth [107]

[108][109][110]. Inoue proposed few empirical rules for choosing the elements susceptible to

lead to bulk metallic glasses from a chemical perspective [95], as follows:

Alloys must be constituted by a minimum of three elements,

A large atomic size difference is requested (at least 12%) between the major

elements,

Enthalpy of mixing between the main elements must be negative.

In spite of this, some authors consider that none of these criteria have been established to be

sufficiently predictive and necessary for bulk-glass formation [85, 111- 114].

The main reproach is stated that none of them can be used to control the properties of the

formed metallic glass. They proposed that elastic moduli of BMGs, having a correlation with

the glass features, mechanical and physical properties could give useful guidelines for

exploring the bulk-glass materials with desirable properties. Besides, some plausible

correlation between the kinetic, thermodynamic, physical and elastic properties of BMGs,

such as between fragility and the Poisson ratio or between Tg and elastic moduli, has indeed

been suggested [115 - 117].

Their research indicates that some features, glass formation and properties of metallic glass

materials, depend intensely of the elastic moduli of its components. Most of the best glass

formers follow the empirical rules, implying just certain physical principles that play vital

roles in the formation of the BMGs in multicomponent system, though these empirical rules

are not sufficient. In other words, they represent only the bare essentials for a glass

formation but are not enough for giving rise to new alloys with interesting tunable

properties.


Chapter 1

27

1.8.3. Properties and features of BMGs

BMGs display many unique and attractive physical and chemical properties.

From a practical perspective the amorphous structure of metallic glasses gives them two

specific properties. Firstly, they experience a glass transition into a supercooled liquid state

upon heating like for other kinds of glasses. In this stage the viscosity can be controlled over

a wide compositional range, creating great flexibility in shaping the glass [118 - 120].

Secondly, the absence of dislocations (crystalline defects) that rule most of the mechanical

properties in the common crystalline alloys, make them much stronger (3-4 times or more).

The combination of low stiffness and high strength gives to metallic glass high resilience,

namely the ability to store elastic strain energy and release it [114, 115, 121].

Their novel and unique features do deliver novel properties, or rather combinations of

properties that make them attractive for varied applications.

Ashby and Greer reported a detailed classification related to the correlation between

fracture toughness and elastic modulus for several types of materials as illustrated in figure

1.24.

For example, the metallic glass materials (zone IV) have higher fracture toughness than the

polymers (zone I) or classical alumina (zone III).

Figure 1.24. Relation between the fracture toughness and elastic modulus for different types of

materials I – polymers, II – oxides glasses, III – ceramics, IV –metallic glasses [121]



28

Moreover, amorphous alloys generally exhibit elastic moduli on the same order as

conventional engineering metals [122]. For instance, Co43Fe20Ta5.5Be31.5 BMG studied by

Inoue et al. has a Young’s modulus of order of 268 GPa and ultrahigh fracture strength of

about 5.185 MPa [123].

Figure 1.25. shows another classification representing the relationship between the Young’s

modulus evolution as a function of the strength for different materials.

Figure 1.25. Relationship between elastic modulus and yield strength for different materials [124]

From the figure it may be seen that BMGs are close to theoretical maximum with yield

strength higher than the silica glass or Al alloys. Several multicomponent materials from the

metallic glasses region are closer to the theoretical limit.

Over the last decade, mechanical properties of metallic glasses (MGs) have been

significantly investigated mainly in their bulk form [125, 126] but not so much as simple

binary thin film system [127]. Their mechanical performance is governed by shear bands

(SBs) formation achieved by localization of flow defects, i.e., shear transformation zone

(STZ) or free volume [122] and the inhomogeneous plastic deformation characteristics [128].

Mechanical performances of metallic glass-based materials are often explored using

nanoindentation measurements [129 - 133] for wear resistance [134], hardness and elastic

modulus [135 - 138], creep [133] or tensile behavior [139 - 143]. Several groups have reported

enhancements in plastic strain to failure resulting from relatively minor changes in alloy

composition [144 - 146].

Y. Huang et al. have conducted isothermal annealing to study the influence of the annealing

time on microstructure and fracture behavior of Zr-Cu-Ni-Al BMG (Figure 1.26.). They have

noted that after a 90-minute annealing time, some crystalline phases are precipitated within

the amorphous matrix. The main remark was that longer annealing time may result in

continuous growth of precipitated crystalline phases (i.e., Ni3Zr). The strain-rate

compressive strength was sensitive to micro-cracks initiation at crystalline-amorphous

interface due to the crystalline phase larger than 10 nm [147].

The enhancement of mechanical properties of BMGs follows generally an increasing trend

with the annealing time due to structural relaxation of the amorphous matrix.


Chapter 1

29

Figure 1.26. The stress strain curves for the as-cast and isothermally annealed Zr50.7Cu28Ni9Al12.5

BMG samples at different times at 743 K using a compressive strain rate of 1.0 x 103s-1 [147]

Even if the high toughness against fracture of metallic glasses is remarkable, there are

considered as best quasi-brittle materials, since they do not possess enough intrinsic

micromechanisms to mitigate the high stress concentrations at crack tips. The brittle fracture

of metallic glass occurs through the propagation of cracks (see figure 1.27a), whereas

engineering alloys generally fail by plastic yielding [122][34].

Figure 1.27. Illustrations of fracture toughness with crack propagation hindered through the

formation of multiple shear bands (a) [148] and fracture surface of La55Al25Cu10Ni5Co5 (b) [149]

During the plastic deformation shear bands are formed in the material. The stresses applied

to a sample may induce local yield that softens significantly the crack, which would reduce

the stress concentration and restrains further crack propagation. Once unstable fracture

initiates, this fracture is considered as “ductile” for most of metallic glasses. The specificity

of MGs is the fracture surface morphology that exhibits shear striations vein-, river-like

patterns such as the ones shown in figure 1.27.b [150].



30

1.8.4. Crystallization and structural stability of metallic glasses-based materials

Many studies on glass alloys [101, 151 - 154] have inferred that various properties of

amorphous alloys are directly related to their structural stability, evaluated by their

crystallization temperature. Amorphous alloys having lower crystallization temperatures

are considered to have inferior structural stability and therefore structural changes can

easily undergo when an extrinsic energy is applied (i.e. thermal, mechanical). [155].

Material’s stability is often evaluated by thermal techniques like differential scanning

calorimetry (DSC) or differential thermal analysis (DTA), able to detect crystallization

events and to measure thermal parameters such as activation energy, incubation time or

mechanism of nucleation and growth [156 - 159].

Typically, the crystallization can occur either during a continuous heating from low

temperature (non-isothermal mode) or during annealing at a given temperature (isothermal

mode). Different models have been proposed with respect to heating process, like the

Johnson-Mehl-Avrami’s (JMA) model [160], Kissinger’s method [161], or the Augis-

Bennett’s model [162]. The JMA is usually the starting point for the explanation of the

isothermal crystallization kinetics [163 - 165].

In figure 1.28. is shown a typical DSC curve for a metallic glass with all the corresponding

characteristic temperatures: glass transition temperature (Tg), crystallization temperature

(Tx), solidus temperature (Ts), liquidus temperature (Tl). Some metallic glasses can have two

crystallization peaks, thus corresponding to a multistage crystallization process [152].

Figure 1.28. DSC curve of Pd79Ag3.5P6Si9.5Ge2 glass with Tg =613 K, Tx= 644 K, Ts= 967 K,

Tl = 1065 K [166]

On the other hand, the most actively studied temperature domain is situated between Tg

and Tx, called the supercooled region (ΔTx), since the temperature during applications is

often situated around Tg [85, 89].

In the literature several reports can be found on the deformation [119, 120, 167], mechanical

behavior [122, 168, [169], crystallization and thermal stability [170 - 173], corrosion resistance

[174, 175] of metallic glasses within the supercooled liquid region.


Chapter 1

31

Most studies demonstrate that the crystallization of metallic glasses involves often several

elementary processes influenced by the thermodynamics (i.e., enthalpy) and annealing

temperature [152, 176].

For instance, Lee et al. [171] reported different activation energy values for the primary

crystallization for Cu-Zr-Al BMG and the whole crystallization mechanism. They showed

that crystallization is an interface-controlled three-dimensional isotropic growth process

with an early-stage saturation of nucleation.

Zhang and coworkers [173] studied as well the Zr-Cu-Al and they proposed that this system

could be regarded as one of the most stable BMGs based on their comparative study.

Atomic diffusion of Al atoms formed for Zr50Cu40Al10 BMG alloys makes a higher

homogenous closed packed structure with respect to the other compositions. It was also

found that unusual B2-type Zr-Cu crystalline phase is mainly formed when glass is rapidly

heated up, as exemplified in figure 1.29.

Figure 1.29. XRD profiles obtained when heated to 700 °C at different heating rates for Zr-Cu-Al

BMG [173]

The long-range diffusion cannot take place sufficiently in the case of high heating rates and,

thus, the B2-type ZrCu phase possessing a similar composition to that of the glass matrix is

formed. This phase is a high-temperature phase and the driving force for the crystallization

toward the B2 phase will be lower than that of the equilibrium phase (i.e., Zr2Cu).

The long-range diffusion is furthermore predictable for the latter crystallization process

including here, as well, the phase separation. This can be an origin of the high thermal

stability and main reason for delaying the crystallization of Zr-Cu-Al BMG.

Similar data are available in the study of Jang et al. [177] studying Zr-Al-Cu-Ni-Si alloys

with different silicon contents from 0 to 10 at.%, and for Surreddi et al. [159] on Al-Gd-Ni-Co

glassy powders. Studies on Mg-based [153, 163] or Fe-based [185, 162] BMGs deal also with

crystallization aspects for better understanding the growth kinetics.



32

1.9. Thin film metallic glass (TFMG)

The first metallic glass in a form of thin film was fabricated, in 1977, by Leamy and Dirks

using a sputtering process [178]. Nowadays, several physical vapor deposition techniques

(PVD) such as pulsed laser [179], arc plasma [180] and sputtering [132, 137] are used for

generating such amorphous alloys.

The TFMGs are usually deposited using pure metal targets by DC magnetron sputtering

system at low temperatures, without bias in a typical configuration as presented in figure

1.30.

Figure 1.30. Example of targets’ configuration inside a deposition chamber [181]

The PVD enables a rapid deposition of complex multicomponent metallic-glass systems.

By controlling sputtering parameters, multicomponent metallic glass films may be

fabricated using single-target magnetron with fully homogenous composition, amorphous

structure and wide supercooled liquid region. The roughness and glass forming ability of

deposited films can be influenced by the deposition conditions.

1.9.1 How to get a TFMG?

Taking advantage of the rapid cooling, PVD techniques are very efficient to deposit glassy

films. Sputtering allows the amorphous or nanocrystalline phases to form during the vapor-

solid deposition with a potential fine control of the chemical composition, ensuring that the

mixing of species takes place on the film surface [179].

The co-sputtering process involving two or three targets appears to be a promising way to

form TFMGs with a wide variation of composition and microstructure.

Chemical composition is well controlled and therefore the properties can be somehow tuned

for the final applications. The sputtering technique is an efficient and cost-effective approach

which can significantly contribute to discover new advanced metallic glass-like systems.

So far, numerous TFMGs based on multicomponent alloys [88], like Zr-based (Zr-Cu-Al [98],

Zr-Cu-Ti [182], Zr-Cu-Al-Ni [184, 185], Zr-Cu-Al-Ag [186]), Pd-based (Pd-Cu-Si [187]) Mg-

based [127] (Mg-Cu-Zr [128]) have been developed and investigated.


Chapter 1

33

Several authors reviewed the Zr-Cu-Al-Ni TFMGs deposited by multicomponent single

targets [181, 188 - 191]. For instance, in the study of Liu et al. [191] they obtained exceptional

flat surface with the desired chemical composition. They used different deposition

conditions (i.e., pressure, power) and revealed that sputtering is a stable process able to

reduce the materials costs. Higher sputtering powers combined with low Ar gas pressures

benefit the film smoothening and flattens (roughness lower than 1 nm) as shown in figure

1.31.

Figure 1.31. Morphology evolution of TFMGs evolution with deposition time and under

different argon pressures for a Zr-Cu-Al-Ni TFMG [191]

Similar findings are presented in the study of Chu et al. [189] where a multicomponent

single target was used to deposit Zr-Cu-Ni-Al alloy thin films. Their interest was to study

the influence of the substrate heating on films’ amorphization and their subsequent

mechanical properties. Results showed that during substrate heating, the thermal energy

absorbed by the deposited molecules can locally rearrange atoms, forming a more stable

structure.

The same interpretation was given by Liu et al. who presented dense columnar-like growth

of coatings with a specific grain-like surface morphology. Films’ microstructure, in surface

as well as in cross-section, is directly associated with the substrate temperature.

According to authors, a more stable structure would be achieved with polygon-like grains,

when the substrate temperature is hold at 400 °C (Figure 1.32.).



34

Figure 1.32. Influence of substrate heating on surface and cross section morphologies for Zr-Cu-

Al-Ni TFMG deposited at different substrate temperatures [189]

On the other hand, Hsiao et al. [188] deposited Zr-Cu-Al-Ni-B films using a multicomponent

Zr-Cu-Al-Ni target and a pure boron one. They reported as well a columnar microstructure

of films. Besides, they also observed a morphology change from a columnar to a fine and

featureless microstructure, when boron is introduced into the TFMGs.

Some authors consider that this structure/morphology change can be explained also by

using models such as the free volume theory [192] or configuration potential energy [188].

At the same time, there are no studies on the TFMGs that explain clearly the influence of the

deposition parameters on the film’s morphology changes, often observed.

On the other hand, as presented in previous section, it is accepted that BMGs exhibit shear

striation with typical vein patterns when they are deformed. A closed and detailed

microscopic observation (Figure 1.33.) of such phenomenon of deformed Zr-based BMGs

reveals a wealth of information on the deformation characteristics and serrated flow from

the study of Torre et al. [150].

Figure 1.33. SEM images showing the shear strains and vein-like patterns for a

Zr57.9Cu22Fe8Al12Pd0.1 BMG [150]


Chapter 1

35

Results of the major shear steps indicate that they are laterally separated into many smaller

shear surfaces and areas with viscous fracture type and vein-like patterns. The spacing of

these smaller shear planes have a size of several tens of nanometers. Similar feature has been

observed for Mg-based BMGs with Cd contents up to 3 at.%, in the study of Sun et al. [106]

and Hsieh et al. for Mg-Cu-Y-Gd BMGs [126]. Shear-banding features and micro-cracks are

also described by Yu et al. [193] for Zr-based metallic glass and Song et al. [194] for Cu-Zr-Ag

BMG.

Of great interest are recent papers on binary Zr-Ni [195] which also reveal the same shear

striations and partially vein pattern that implies similar deformation mechanism as for

BMGs. Matteo et al., studying effects of films’ thickness of a Zr-Ni alloy (between 300 – 900

nm) observed only at a thickness of 500 nm the characteristics river-like features.

The vein patterns are the typical characteristic of localized plastic flow which occurs in the

shear bands for the glassy alloys indicating enhanced toughness [122, 196, 197].

The initiation of periodic striations is supposed to be the result of self-assembled cavities

when their size reaches a critical value. This is an indication of activation of plastic flow and

the deformation in the compression process [196].

1.9.2. Binary TFMG systems

In the previous section general studies and examples were given for more complicated

TFMGs in a general way for familiarizing with the metallic glass-like films.

In what follows, since the purpose of the current thesis is the investigation of binary Zr-Cu

films, a more detailed bibliography research will be spotlighted for this specific system.

In the literature, only few studies on simple binary TFMGs can be found.

For instance, binary Mg-Cu thin films are studied for their lighter weight and special

application in the biomedical area. Chou et al. [127] investigated this system in a

composition range varying from 38 to 82 at.% and they found partial amorphous films for

some compositions.

R. Banerfee et al. [198] studied the relation between microstructure, phase stability and

electrical transport properties of Cu-Nb alloy thin films (composition from 5 to 90 at.% Nb).

At low Nb and Cu contents (<25 at.% of the alloying composition) only single body centered

cubic (bcc) and face centered cubic (fcc) solid solutions are formed. Films’ composition

corresponding to the middle of the investigated range (30 – 70 at.% Nb) exhibit an

amorphous structure with low temperature coefficients of resistivity (TCR).

For instance, 90.4 μ.ohm.cm for Cu–45 at. % Nb and 96.2 μ.ohm.cm for Cu–70 at. % Nb. The

onset temperatures for the superconducting transitions are also deduced from their

experiments.

M. Ghidelli et al. reported some results on Zr-Ni TFMGs in a chemical composition ranged

between 42 and 85 at.% [199]. The local atomic structure was investigated by TEM and XRD,

and a method based on the Nagel and Tauc’s criterion [200] was applied to determine the

(meta)stability of Zr-Ni TFMGs. The mechanical size effect in Zr65Ni35 TFMG was also

described by Brillouin spectroscopy and nanoindentation measurements [195].



36

Mechanical properties assessed by the nanoindentation technique were studied as well by F.

Zeng et al. [138] for Cu-Ta amorphous films with Ta content from 25.6 to 96 at.%.

They demonstrated that the hardness and Young’s modulus vary linearly with the Ta

concentration.

Chen et al. [201] focused their research on binary Zr-Ti and Zr-Cu thin films. For the Zr-Cu

system, fully amorphous TFMGs (Figure 1.34. a) were got for various deposition conditions,

suggesting its intrinsic high vitrification tendency. In the case of Zr-Ti (figure 1.34. b) low

sputtering powers and substrate temperatures are required to form amorphous films.

TEM bright field plan view image reveals the existence of spots on the amorphous ring

indicating nevertheless the presence of a further fine nanocrystalline (~ 2 nm in diameter)

phase.

Figure 1.34. TEM bright field plan-view image of the Zr45Cu55 (a) and Zr54Ti46 (b) co-sputtered

films with an inserted diffraction pattern [201]

These studies evidence that a correlation exists between the composition of the films, their

structure and properties. Zr-Cu thin films present an amorphous structure in a large range

of composition and may present interesting properties depending on the composition.

A further study of this promising system is one of the objectives of this work.


Chapter 1

37

1.9.3. Binary Zr-Cu

Extensive work has been carried out on the synthesis [202 - 204], microstructure [201, 205,

206], mechanical and functional properties [137, 207, 208] and simulations [209] on the

binary Zr-Cu TFMG system.

Among the first studies, Musil and Zeman [208], Dudonis et al. [202] and Minemura et al.

[203] explored the formation and thermal stability of this system in a wide compositional

range using different deposition substrates and conditions. Their findings clearly showed

the composition dependence of this system with similar tendencies for crystallization [203],

microhardness [208], reflectivity and sheet resistance [202].

Based on crystallization temperature, as related by Minemura et al., the Zr-Cu system shows

an optimal thermal resistance around the composition of Zr40Cu60 (Figure 1.35.).

Figure 1.35. Composition dependence of crystallization temperature for Zr-Cu BMG [203]

Besides, Wang and co-authors found a similar composition-dependency using

nanoindentation values [137]. Their results demonstrate that chemical compositions

corresponding to hardness peaks match well with the composition densities peaks.

For them the hardness peaks can be qualitatively interpreted in terms of denser atomic

packing model and free-volume theory. The method provides quantitative experimental

data to establish new structural models and could be a fast and direct way to search for new

glass-forming alloys.

Finally, the composition influence has been also investigated in the study of P. Coddet et al.

[207] where authors explored the influence of deposition conditions, and of the Cu content

on the hardness and Young’s Modulus values. The composition can be easily controlled just

by changing the applied current intensities on the metallic targets as can be seen in figure

1.36. for a binary Zr-Cu TFMG [207].



38

Figure 1.36. Composition of Zr–Cu binary alloys as a function of the current intensity on the Cu

target [207]

Moreover, as previously reported for other binary systems [138] the mechanical

characteristics for the Zr-Cu TFMGs increase with the Cu content in an almost linear way as

illustrated in figure 1.37.

Figure 1.37. Evolution of hardness and Young's modulus values of binary Zr–Cu films versus the

Cu content [207]

1.9.4. Mechanical properties of TFMGs

Some recent works have been published on thin film metallic glasses as well for

understanding the mechanical behavior [88, 132, 168, 185, 195, 207] in terms of nanoscale

creep deformation [210], plastic flow [184], hardness [138] [211] and friction coefficient [212].

For instance, the reason for the linear enhancement of the elastic modulus of Cu-Ta

amorphous films with Ta content (Figure 1.38.) was attributed to the disordered state in the

interface of the amorphous matrix.

Figure 1.38. Evolution of mechanical properties of binary Cu-Ta amorphous films. E’ in this

figure represents the calculated elastic moduli of the mechanical mixtures based on the relative

elastic moduli of different materials [138]


Chapter 1

39

Furthermore, a mechanism in which the inter-atomic potential of a disordered state was

proposed. The model considers the inter-atomic distance and the nature of the inter-atomic

potential. They stated that an increase of the inter-atomic distance will induce a decrease of

the elastic modulus if the nature of the potential does not change. The enhanced modulus of

their films could not come from the variation of the inter-atomic distance since the average

distance increases after amorphization [138].

Moreover, they consider that due to the differences in free energy for the constituent

elements a part is stored in the alloy after amorphization. This energy modifies the nature of

the interatomic potential that will later influence the elastic modulus. However, further

work still has to be considered to measure the actual inter-atomic potential and to confirm

their model.

Cao and coworkers gave information about the strain rate deformation of magnetron Ni-Nb,

Cu-Zr-Al, Zr-Cu-Ni-Al MG thin films [146]. In a recent paper, Chuang et al. reported that

the nitrogen content was the key factor for increased values of mechanical properties,

surface roughness and coefficient of friction for Zr-Ni-Al-Si TFMG [212].

The presence of fine nanocrystalline Mg2Cu particles in the structure could also affect the

hardness as reported by Chou et al. for binary Mg-Cu thin films [127]. In other study they

prepared micropillars of Zr55Cu31Ti14 amorphous films with Ta layers and investigated the

mechanical response at room temperature (Figure 1.39). It has been concluded that

mechanical properties can be improved via inserting of high-stiffness Ta layers to form

composite-like microstructure.

Figure 1.39. SEM images of deformed monolithic ZrCuTi (a), 50 nm-ZrCuTi/5 (b) nm-Ta and 50

nm-ZrCuTi/50 nm-Ta (d) and the engineering stress-strain curves (c) [127]

The underlying physical mechanisms are related to the material geometry and composition.

For the ZrCuTi/Ta composite micropillars, the deformation is linked to the amorphous

geometric volume, whereas for different-thickness of Ta layer the share behavior is

connected with the formation of an equal-thickness plastic zone in the amorphous layers

[213].



40

A set of micro-pillars with varying thickness of Zr-based TFMGs were also prepared by Ye

et al. [184] to explain the relation between yield strength and Young’s Modulus.

The results from micro-compression show that plastic flows in the TFMG-based micro-pillar

prove a strong size-and-shape dependency.

This is evidenced by smoother plastic deformation transition from the inhomogeneous to

homogeneous mode when the TFMG-based micropillars (with a submicron-scale film

thickness) are deformed into the shape of a low aspect ratio.

The mechanical properties of Zr-based TFMGs depend a lot on the deposition conditions.

For instance, Chuang et al. discussed such issues during a complex analysis of the

mechanical characteristics of Zr-based TFMGs with a nitrogen content ranging from 0 to

17.7 at.%. The surface roughness, hardness, coefficient of friction and critical load increased

with the increase of the nitrogen content.

On the other hand, the mechanical characteristics of TFMG are affected as well by the local

atomic bonding and less by microstructure. For example, the elastic modulus and hardness

are strongly influenced by the Ta content for Zr-Cu-Ti TFMG. Their values (elastic modulus

up to ~140 GPa and hardness 10 GPa) increased with the Ta content in the film.

Compared to Zr-, Cu- and Ti-based amorphous alloys these values are much higher, which

makes the quaternary Zr-Cu-Ti-Ta promising for innovating applications [211].

1.10. Advanced applications

Metallic glasses in the form of thin film are evolving as alternative coatings with a great

potential as structural materials. Their versatility facilitates their exploitation for a wide

range of applications starting from biomedical, micro and opto-electronics, nano-devices to

military components (Figures 1.40 – 1.41). For instance, due to their lighter weight Mg-Cu

TFMG can be used in the biomedical field [127].

Furthermore, Zr-based TFMGs could be exploited for medical tools to improve the

sharpness of surgical blades, suture needles, and surgery scissors or for dental components.

The liquid crystal displays (LCDs) performance can be improved when a ZrCu-based

TFMG/ITO bilayer is applied compared to conventional indium tin oxide (ITO) transparent

electrode [88].

In addition, TFMGs have no grain boundary in opposite to conventional crystalline

materials do. Such a structural plus allowed Zr-Cu-Al to be the first exploited for MEMS

applications [214]. It is yielded the use of Zr-based metallic glass film to improve the

tribological behavior of 316L stainless steel. It was concluded that the high adhesion

strength of interfacial layer between the film/steel and low residual stress promotes

metallurgical bonding, hence improving the mechanical properties [182, 185].


Chapter 1

41

Figure 1.40. Some of the main applications and direct product-targeted for BMG-based materials.

Adapted from ref. [121]

Figure 1.41. Examples of metallic glass-based materials used for micro-tools (A) [215], hard drive

components, an outer appearance of the self-made Coriolis flowmeter using the Ti–Zr–Cu–Ni–Sn

glassy alloy pipe (C) [90] and golf clubs (D)

The TFMG-based materials were recently developed for their antimicrobial protection

against pathogens as Staphylococcus aureus, Escherichia coli, Pseudomonas or mixed-microbial

infections that are usually transmitted from the hospital environment (instruments or

equipment) by direct hand-to-hand contact [88, 216, 217].

The medical environmental surfaces such as door handles, surgical instruments, and push

plates are generally made of stainless steel.

Chu et al. [218] have reported that 200 nm thick Cu-based and Zr-based TFMGs exhibit high

efficiency levels as antimicrobial materials against the growth of E. coli and S. aureus

bacteria’s than bare silicon wafer (Figure 1.42.).



42

Figure 1.42. Photos showing the antibacterial rate (AR) evaluation of TFMGs against pathogens

on bare and coated surfaces after incubation for 24 h at 37 °C [217]

The bacterial adhesion of pathogens to both Zr- and Cu- based TFMG is hindered to

different extents that of the bare Si wafer. Both TFMG show antibacterial rate (AR) values of

almost 100% against E. coli, whereas the Cu-based TFMG is less effective on S. aureus. Chen

et al. studied the microstructure and antimicrobial capabilities of Zr-Cu-Al-Ag TFMG

against C. albicans, E. coli and P. aeruginosa.

As compared to the polished stainless steel substrate, the TFMG investigated reveal an

effective antimicrobial activity and show an inverse relationship between surface roughness

and growth area. By silver addition the microbial activity and the viability of mature biofilm

for strain of P. aeruginosa are completely inhibited on TFMGs after the incubation time of 72

h [219].

Chiang et al. investigated also the antimicrobial effects of Zr-Al-Ni-Cu on most common

nosocomial infection pathogens [217]. Their results showed a positive effect, reducing

colonization and biofilm’s formation on surfaces. Similar approach can be found in the

study of Kaushik et al. [220] on Ti-Cu-Ni TFMG.

Another promising outlet for TFMGs is represented by the corrosion protectiveness

afforded by the modified surfaces. MGs-based materials can indeed be used for applications

where a protection against corrosion attacks is required, thanks to their structure and

chemistry homogeneities.

For instance, according to the studies of Zander et al. [221] and Chen et al. [222] the corrosion

mechanisms of Cu-based BMGs in NaCl solution is governed by a stable passive film.

Liu et al. [223] stated that an addition of small amounts of Cr, Mo or W had a further

beneficial effect on corrosion resistance of Cu-based BMGs in both acidic and alkaline

media.

Passive films enriched in TiO2 and ZrO2 formed on Ti- and Zr-based BMGs, are more stable

and denser, explaining their better protective effect [224, 225].

Xie et al. [174] investigated the electrochemical behavior in the Hank’s solution of Ti-based

BMGs. They concluded that such microstructural heterogeneity may affect the corrosion

resistance, favoring crevice corrosion mechanisms.


Chapter 1

43

The strong passive ability yields Zr-based alloys to be still one of the most investigated bulk

glass formers The zirconium belongs to the “valve metal” class, is indeed prone to form a

thin passive film playing the role of barrier layer with low ionic and electronic

conductivities [85, 87, 89, 105].

Some works concern binary alloys [226 - 230] with regards to corrosion and crystallization

resistances or mechanical properties, and nearly none for TFMGs, apart from more

complicated systems such as Zr-Cu-Ni-Al-based [231, 232], Zr-Cu-Al-Ag [233] or Cu-Hf-Al

[234]. For instance, C.Y. Chuang et al. [231] showed that Zr-based TFMG exhibits a better

corrosion resistance compared with AISI 420 substrate due to its amorphous structure as

exemplified in figure 1.43. below.

Figure 1.43. Example of a potentiodynamic polarization curves of the Zr-based thin film metallic

glass and AISI 420 substrate [231]

For the TFMG a higher potential (Ecorr) and lower corrosion current density (Icorr) are

observed as compared with AISI 420 substrate denoting a better corrosion resistance.

They concluded that Zr-based TFMGs plays a good protection role against corrosion attack

by HCl aqueous solution due to its amorphous and defect-free structure.



44

Chapter summary

In this chapter it was presented the state of the art on the two different classes of PVD thin

films investigated, namely hard CrN-based coatings and Zr-Cu thin films metallic glasses.

In part A, the presented studies are dealing with the latest developments in terms of high

temperature oxidation prevention of the machining tools protected by nitride-based

coatings. The development of ceramic films has intensively increased during the past twenty

years from the already classical TiN-based systems to newly CrN-based, and today to

ternary systems such as CrAlN. Literature dedicated to ceramic hard coatings is very

prolific in relation with the large development of PVD processes since 1990s in both

academic and industrial domains. In this study, we have then limited our investigations to

recent papers linked to nitride-based coatings and more specifically to CrN-based ones.

The first generation of TiN has highlighted the importance and efficiency of ceramic

coatings to protect the machining tools used in severe environments. The development of

PVD coatings made today possible to improve the oxidation performance of such surfaces

by applying CrN-based coatings. However, the industry requirements are in a continuous

changing trend and new solutions had to be initiated as a response to their demand,

pushing the research towards new more complex ways, such as CrAlN and lately CrAlYN

systems. The chemical nature, structure and physical properties are significantly enhanced

by incorporation of Al and/or Y into the fcc matrix. When the Al content is below the

solubility limit (i.e. lower than 77 at. %), the structure formed is a solid solution where

lattice distortions are induced.

The good oxidation resistance is mainly linked to the structural stability of the film

(transformation into Cr2N phase), and to hindered transport of the species (outward process

for the metals and inward diffusion into the coatings for oxygen).

The findings suggest that in the case of CrAlN and CrAlYN the protective oxide scale

formed on the surface would be composed by a stable mixture of both Cr2O3 and Al2O3

oxides, or by a single mixed (Al,Cr)2O3 oxide layer. Such a complex layer ensures a higher

protectiveness as compared with single Cr2O3 grown onto CrN due to their higher chemical

stability.

Among the compared CrN-based coatings, CrAlYN has the best oxidation resistance which

is attributed to the delaying of oxide layer growth and grain refinement of the structure.

The addition of Al indeed shifts the onset and critical oxidation temperatures to higher

values compared with the classical CrN. The beneficial effect of Al addition is related

mainly to the formation of a dense Al2O3, which restrain the diffusion of the oxygen.

The further addition of large yttrium atoms seems efficient to retard the diffusion, resulting

in an increased thermal stability and better oxidation resistance at high temperatures.

If the literature focused on the already “old” CrN coating, papers devoted to the ternary and

quaternary Al-containing films are more limited. In particular, very few thermodynamic

data are available on these advanced recent systems. Therefore, our objective will be to

analyze the role of each Al and Y alloying element on the oxidation behavior of films.

For this study, we propose a multiscale analysis based on a qualitative and quantitative

approach associated with in situ and post-mortem experiments.


Chapter 1

45

In the second part, some general notions related to the bulk metallic glasses and newly thin

metallic glass films are highlighted. The main characteristics of such materials are the

crystals-free structure and distinctive properties such as high thermal stability and strength

limits, glass forming ability, excellent elastic and mechanical properties.

Several examples of BMGs and multicomponent TFMGs were given in order to present the

state of the art of these materials. Most of the highlighted studies deal with the formation

and deposition of the metallic glass-like materials with the focus put on their structural

stability, crystallization and mechanical properties.

The TFMGs are successfully deposited by many deposition techniques such as pulsed laser,

arc plasma and most common sputtering, using pure metals or multicomponent targets.

The binary Zr-Cu TFMG system can provide a first-rate system for structural models and a

base for future thin metallic glass-based alloys due to its high glass forming ability and

structural stability. Up to now, most of the studies on TFMGs available in literature are

dedicated to the compositional range where the film remains amorphous and their

relationship between the deposition parameters and specific properties (e.g. mechanical,

electrical).

The thin metallic glass films can be considered as new class of coatings for applications

where structural and chemical homogeneity is mandatory. Their adaptability permits the

exploitation for a complex variety of applications from micro and opto-electronics to

nanodevices or biomedicine.

Finally, the chapter provides some studies on several types of thin films developed for

advanced future applications in the field of corrosion or as bactericide layers for instance.

It can also be underlined that most of studies are focused on Zr-Cu films characterized by a

balanced chemical composition around of 50 at.% Cu. Moreover, few thermal

characterizations are published and there is also a lack of information regarding the

coating’s physico-chemical damage.

Therefore, originality of our work relies on the wide composition range of investigated films

scanning the whole phase diagram. Moreover, we will not only study the films’ structure in

their amorphous state, but also beyond the glass transition temperature.

The opportunity to use TFMGs as corrosion protection will also be assessed.



46

List of references

[1] Y. Benlatreche, C. Nouveau, R. Marchal, J.-P. Ferreira Martins, and H. Aknouche,

“Applications of CrAlN ternary system in wood machining of medium density fibreboard

(MDF),” Wear, vol. 267, no. 5–8, pp. 1056–1061, 2009.

[2] C. Nouveau, C. Labidi, R. Collet, Y. Benlatreche, and M.-A. Djouadi, “Effect of surface

finishing such as sand-blasting and CrAlN hard coatings on the cutting edge’s peeling

tools’ wear resistance,” Wear, vol. 267, no. 5–8, pp. 1062–1067, 2009.

[3] B. Navinšek, P. Panjan, and I. Milošev, “Industrial applications of CrN (PVD) coatings,

deposited at high and low temperatures,” Surf. Coat. Technol., vol. 97, no. 1–3, pp. 182–191,

1997.

[4] E. Spain, J. C. Avelar-Batista, M. Letch, J. Housden, and B. Lerga, “Characterisation and

applications of Cr–Al–N coatings,” Surf. Coat. Technol., vol. 200, no. 5–6, pp. 1507–1513,

2005.

[5] I. Konyashin, G. Fox-Rabinovich, and A. Dodonov, “TiN thin films deposited by filtered

arc-evaporation: structure, properties and applications,” J. Mater. Sci., vol. 32, no. 22, pp.

6029–6038, 1997.

[6] K.-D. Bouzakis, N. Michailidis, S. Gerardis, G. Katirtzoglou, E. Lili, M. Pappa, M.

Brizuela, A. Garcia-Luis, and R. Cremer, “Correlation of the impact resistance of variously

doped CrAlN PVD coatings with their cutting performance in milling aerospace alloys,”

Surf. Coat. Technol., vol. 203, no. 5–7, pp. 781–785, 2008.

[7] J. L. Endrino, G. S. Fox-Rabinovich, and C. Gey, “Hard AlTiN, AlCrN PVD coatings for

machining of austenitic stainless steel,” Surf. Coat. Technol., vol. 200, no. 24, pp. 6840–6845,

août 2006.

[8] C. Mitterer, P. . Mayrhofer, and J. Musil, “Thermal stability of PVD hard coatings,”

Vacuum, vol. 71, no. 1–2, pp. 279–284, 2003.

[9] F. Vaz, P. Machado, L. Rebouta, P. Cerqueira, P. Goudeau, J. P. Rivière, E. Alves, K.

Pischow, and J. de Rijk, “Mechanical characterization of reactively magnetron-sputtered

TiN films,” Surf. Coat. Technol., vol. 174–175, pp. 375–382, 2003.

[10] F. Vaz, J. Ferreira, E. Ribeiro, L. Rebouta, S. Lanceros-Méndez, J. A. Mendes, E. Alves, P.

Goudeau, J. P. Rivière, F. Ribeiro, I. Moutinho, K. Pischow, and J. de Rijk, “Influence of

nitrogen content on the structural, mechanical and electrical properties of TiN thin films,”

Surf. Coat. Technol., vol. 191, no. 2–3, pp. 317–323, 2005.

[11] F. Vaz, L. Rebouta, M. Andritschky, M. F. da Silva, and J. C. Soares, “Thermal oxidation of

Ti1 − xAlxN coatings in air,” J. Eur. Ceram. Soc., vol. 17, no. 15–16, pp. 1971–1977, 1997.

[12] L. Aihua, D. Jianxin, C. Haibing, C. Yangyang, and Z. Jun, “Friction and wear properties

of TiN, TiAlN, AlTiN and CrAlN PVD nitride coatings,” Int. J. Refract. Met. Hard Mater.,

vol. 31, pp. 82–88, 2012.

[13] V. Chawla, A. Chawla, B. S. Sidhu, S. Prakash, and D. Puri, “Oxidation Behavior of

Nanostructured TiAlN and AlCrN Thin Coatings on ASTM-SA213-T-22 Boiler Steel,” J.

Miner. Mater. Charact. Eng., vol. 09, no. 11, p. 1037, Nov. 2010.

[14] Y. C. Chim, X. Z. Ding, X. T. Zeng, and S. Zhang, “Oxidation resistance of TiN, CrN,

TiAlN and CrAlN coatings deposited by lateral rotating cathode arc,” Thin Solid Films,

vol. 517, no. 17, pp. 4845–4849, 2009.

[15] M. Kawate, A. Kimura Hashimoto, and T. Suzuki, “Oxidation resistance of Cr1−XAlXN

and Ti1−XAlXN films,” Surf. Coat. Technol., vol. 165, no. 2, pp. 163–167, 2003.

[16] H. C. Barshilia, N. Selvakumar, B. Deepthi, and K. S. Rajam, “A comparative study of

reactive direct current magnetron sputtered CrAlN and CrN coatings,” Surf. Coat.

Technol., vol. 201, no. 6, pp. 2193–2201, 2006.


Chapter 1

47

[17] J. Lin, B. Mishra, J. J. Moore, and W. D. Sproul, “A study of the oxidation behavior of CrN

and CrAlN thin films in air using DSC and TGA analyses,” Surf. Coat. Technol., vol. 202,

no. 14, pp. 3272–3283, 2008.

[18] Z. B. Qi, B. Liu, Z. T. Wu, F. P. Zhu, Z. C. Wang, and C. H. Wu, “A comparative study of

the oxidation behavior of Cr2N and CrN coatings,” Thin Solid Films, vol. 544, pp. 515–520,

2013.

[19] M. Odén, C. Ericsson, G. Håkansson, and H. Ljungcrantz, “Microstructure and

mechanical behavior of arc-evaporated Cr–N coatings,” Surf. Coat. Technol., vol. 114, no. 1,

pp. 39–51, 1999.

[20] P. H. Mayrhofer, G. Tischler, and C. Mitterer, “Microstructure and mechanical/thermal

properties of Cr–N coatings deposited by reactive unbalanced magnetron sputtering,”

Surf. Coat. Technol., vol. 142–144, pp. 78–84, 2001.

[21] M. A. Djouadi, C. Nouveau, P. Beer, and M. Lambertin, “CrxNy hard coatings deposited

with PVD method on tools for wood machining,” Surf. Coat. Technol., vol. 133–134, pp.

478–483, 2000.

[22] T. Polcar and A. Cavaleiro, “High temperature properties of CrAlN, CrAlSiN and

AlCrSiN coatings – Structure and oxidation,” Mater. Chem. Phys., vol. 129, no. 1–2, pp.

195–201, 2011.

[23] J. L. Endrino, G. S. Fox-Rabinovich, A. Reiter, S. V. Veldhuis, R. Escobar Galindo, J. M.

Albella, and J. F. Marco, “Oxidation tuning in AlCrN coatings,” Surf. Coat. Technol., vol.

201, no. 8, pp. 4505–4511, 2007.

[24] E. Le Bourhis, P. Goudeau, M. H. Staia, E. Carrasquero, and E. S. Puchi-Cabrera,

“Mechanical properties of hard AlCrN-based coated substrates,” Surf. Coat. Technol., vol.

203, no. 19, pp. 2961–2968, 2009.

[25] Y. H. Cheng, T. Browne, and B. Heckerman, “Mechanical and tribological properties of

CrN coatings deposited by large area filtered cathodic arc,” Wear, vol. 271, no. 5–6, pp.

775–782, 2011.

[26] J. J. Zhang, M. X. Wang, J. Yang, Q. X. Liu, and D. J. Li, “Enhancing mechanical and

tribological performance of multilayered CrN/ZrN coatings,” Surf. Coat. Technol., vol. 201,

no. 9–11, pp. 5186–5189, février 2007.

[27] P. H. Mayrhofer, F. Rovere, M. Moser, C. Strondl, and R. Tietema, “Thermally induced

transitions of CrN thin films,” Scr. Mater., vol. 57, no. 3, pp. 249–252, août 2007.

[28] W. Ernst, J. Neidhardt, H. Willmann, B. Sartory, P. H. Mayrhofer, and C. Mitterer,

“Thermal decomposition routes of CrN hard coatings synthesized by reactive arc

evaporation and magnetron sputtering,” Thin Solid Films, vol. 517, no. 2, pp. 568–574,

Nov. 2008.

[29] C. Paternoster, A. Fabrizi, R. Cecchini, M. E. Mehtedi, and P. Choquet, “Thermal stability

of CrN x nanometric coatings deposited on stainless steel,” J. Mater. Sci., vol. 43, no. 10,

pp. 3377–3384, May 2008.

[30] D. . Lee, Y. . Lee, and S. . Kwon, “High temperature oxidation of a CrN coating deposited

on a steel substrate by ion plating,” Surf. Coat. Technol., vol. 141, no. 2–3, pp. 227–231,

2001.

[31] P. . Mayrhofer, H. Willmann, and C. Mitterer, “Oxidation kinetics of sputtered Cr–N hard

coatings,” Surf. Coat. Technol., vol. 146–147, pp. 222–228, 2001.

[32] A. H. Liu, J. X. Deng, H. B. Cui, J. Zhao, and X. Ai, “Oxidation Resistance of CrN and

CrAlN Coating Tools,” Adv. Mater. Res., vol. 189–193, pp. 137–141, 2011.

[33] J. Lin, N. Zhang, W. D. Sproul, and J. J. Moore, “A comparison of the oxidation behavior

of CrN films deposited using continuous dc, pulsed dc and modulated pulsed power

magnetron sputtering,” Surf. Coat. Technol., vol. 206, no. 14, pp. 3283–3290, 2012.



48

[34] A. Mège-Revil, P. Steyer, G. Thollet, R. Chiriac, C. Sigala, J. C. Sanchéz-Lopéz, and C.

Esnouf, “Thermogravimetric and in situ SEM characterisation of the oxidation

phenomena of protective nanocomposite nitride films deposited on steel,” Surf. Coat.

Technol., vol. 204, no. 6–7, pp. 893–901, 2009.

[35] A. Mège-Revil, P. Steyer, J. Fontaine, J. F. Pierson, and C. Esnouf, “Oxidation and tribo-

oxidation of nanocomposite Cr–Si–N coatings deposited by a hybrid arc/magnetron

process,” Surf. Coat. Technol., vol. 204, no. 6–7, pp. 973–977, 2009.

[36] G. S. Kim and S. Y. Lee, “Microstructure and mechanical properties of AlCrN films

deposited by CFUBMS,” Surf. Coat. Technol., vol. 201, no. 7, pp. 4361–4366, 2006.

[37] J. E. Sanchéz, O. M. Sanchéz, L. Ipaz, W. Aperador, J. C. Caicedo, C. Amaya, M. A. H.

Landaverde, F. E. Beltran, J. Muñoz-Saldaña, and G. Zambrano, “Mechanical, tribological,

and electrochemical behavior of Cr1−xAlxN coatings deposited by r.f. reactive magnetron

co-sputtering method,” Appl. Surf. Sci., vol. 256, no. 8, pp. 2380–2387, 2010.

[38] J. Lin, B. Mishra, J. J. Moore, and W. D. Sproul, “Microstructure, mechanical and

tribological properties of Cr1−xAlxN films deposited by pulsed-closed field unbalanced

magnetron sputtering (P-CFUBMS),” Surf. Coat. Technol., vol. 201, no. 7, pp. 4329–4334,

2006.

[39] H. N. Shah and R. Jayaganthan, “Influence of Al Contents on the Microstructure,

Mechanical, and Wear properties of Magnetron Sputtered CrAlN Coatings,” J. Mater. Eng.

Perform., vol. 21, no. 9, pp. 2002–2009.

[40] O. M. S. Quintero, W. A. Chaparro, L. Ipaz, J. E. S. Barco, F. E. Beltrán, and G. Zambrano,

“Influence of the microstructure on the electrochemical properties of Al-Cr-N coatings

deposited by co-sputtering method from a Cr-Al binary target,” Mater. Res., vol. 16, no. 1,

pp. 204–214, 2013.

[41] B. Tlili, N. Mustapha, C. Nouveau, Y. Benlatreche, G. Guillemot, and M. Lambertin,

“Correlation between thermal properties and aluminum fractions in CrAlN layers

deposited by PVD technique,” Vacuum, vol. 84, no. 9, pp. 1067–1074, 2010.

[42] W. Z. Li, Q. Z. Chen, T. Polcar, R. Serra, and A. Cavaleiro, “Influence of Zr alloying on the

mechanical properties, thermal stability and oxidation resistance of Cr–Al–N coatings,”

Appl. Surf. Sci., vol. 317, pp. 269–277, 2014.

[43] C. Nouveau, B. Tlili, H. Aknouche, Y. Benlatreche, and B. Patel, “Comparison of CrAlN

layers obtained with one (CrAl) or two targets (Cr and Al) by magnetron sputtering,”

Thin Solid Films, vol. 520, no. 7, pp. 2932–2937, 2012.

[44] C. Nouveau, C. Labidi, J.-P. F. Martin, R. Collet, and A. Djouadi, “Application of CrAlN

coatings on carbide substrates in routing of MDF,” Wear, vol. 263, no. 7–12, pp. 1291–1299,

2007.

[45] J. Romero, M. A. Gómez, J. Esteve, F. Montalà, L. Carreras, M. Grifol, and A. Lousa,

“CrAlN coatings deposited by cathodic arc evaporation at different substrate bias,” Thin

Solid Films, vol. 515, no. 1, pp. 113–117, 2006.

[46] M. Mühlbacher, R. Franz, J. Paulitsch, H. Rudigier, P. Polcik, P. H. Mayrhofer, and C.

Mitterer, “Influence of Fe impurities on structure and properties of arc-evaporated AlCrN

coatings,” Surf. Coat. Technol., vol. 215, pp. 96–103, 2013.

[47] R. Franz, J. Neidhardt, R. Kaindl, B. Sartory, R. Tessadri, M. Lechthaler, P. Polcik, and C.

Mitterer, “Influence of phase transition on the tribological performance of arc-evaporated

AlCrVN hard coatings,” Surf. Coat. Technol., vol. 203, no. 8, pp. 1101–1105, 2009.

[48] A. E. Reiter, C. Mitterer, and B. Sartory, “Oxidation of arc-evaporated Al1−xCrxN

coatings,” J. Vac. Sci. Technol. A, vol. 25, no. 4, pp. 711–720, 2007.

[49] A. E. Reiter, V. H. Derflinger, B. Hanselmann, T. Bachmann, and B. Sartory, “Investigation

of the properties of Al1−xCrxN coatings prepared by cathodic arc evaporation,” Surf.

Coat. Technol., vol. 200, no. 7, pp. 2114–2122, 2005.


Chapter 1

49

[50] A. Kimura, M. Kawate, H. Hasegawa, and T. Suzuki, “Anisotropic lattice expansion and

shrinkage of hexagonal TiAlN and CrAlN films,” in Surface & coatings technology, vol. 169–

70, pp. 367–370.

[51] X. Ding, X. T. Zeng, Y. C. Liu, F. Z. Fang, and G. C. Lim, “Cr1−xAlxN coatings deposited

by lateral rotating cathode arc for high speed machining applications,” Thin Solid Films,

vol. 516, no. 8, pp. 1710–1715, 2008.

[52] H. Willmann, P. H. Mayrhofer, P. O. Å. Persson, A. E. Reiter, L. Hultman, and C. Mitterer,

“Thermal stability of Al–Cr–N hard coatings,” Scr. Mater., vol. 54, no. 11, pp. 1847–1851,

2006.

[53] Y. Makino and K. Nogi, “Synthesis of pseudobinary Cr-Al-N films with B1 structure by

rf-assisted magnetron sputtering method,” Surf. Coat. Technol., vol. 98, no. 1–3, pp. 1008–

1012, 1998.

[54] K. Bobzin, N. Bagcivan, P. Immich, S. Bolz, R. Cremer, and T. Leyendecker, “Mechanical

properties and oxidation behaviour of (Al,Cr)N and (Al,Cr,Si)N coatings for cutting tools

deposited by HPPMS,” Thin Solid Films, vol. 517, no. 3, pp. 1251–1256, 2008.

[55] H. Hasegawa, K. Ohashi, S. Tsukamoto, T. Sato, and T. Suzuki, “Characterization of

quaternary (Cr,Al)N-based films synthesized by the cathodic arc method,” Surf. Coat.

Technol., vol. 202, no. 4–7, pp. 786–789, 2007.

[56] O. Banakh, P. . Schmid, R. Sanjinés, and F. Lévy, “High-temperature oxidation resistance

of Cr1−xAlxN thin films deposited by reactive magnetron sputtering,” Surf. Coat. Technol.,

vol. 163–164, pp. 57–61, 2003.

[57] S. Khamseh, “Oxidation Resistance of CrAlN Films with Different Microstructures

Prepared by Pulsed DC Balanced Magnetron Sputtering System,” Mater. Trans., vol. 51,

no. 2, p. 271, 2010.

[58] R. E. Galindo, J. L. Endrino, R. Martínez, and J. M. Albella, “Improving the oxidation

resistance of AlCrN coatings by tailoring chromium out-diffusion,” Spectrochim. Acta Part

B At. Spectrosc., vol. 65, no. 11, pp. 950–958, 2010.

[59] Y. Feng, L. Zhang, R. Ke, Q. Wan, Z. Wang, and Z. Lu, “Thermal stability and oxidation

behavior of AlTiN, AlCrN and AlCrSiWN coatings,” Int. J. Refract. Met. Hard Mater., vol.

43, pp. 241–249, 2014.

[60] J. L. Mo and M. H. Zhu, “Sliding tribological behaviors of PVD CrN and AlCrN coatings

against Si3N4 ceramic and pure titanium,” Wear, vol. 267, no. 5–8, pp. 874–881, 2009.

[61] Y. X. Wang, S. Zhang, J.-W. Lee, W. S. Lew, and B. Li, “Influence of bias voltage on the

hardness and toughness of CrAlN coatings via magnetron sputtering,” Surf. Coat.

Technol., vol. 206, no. 24, pp. 5103–5107, 2012.

[62] J. Lin, M. Pinkas, W. D. Sproul, and J. J. Moore, “The phase and microstructure of CrAlN

films deposited by pulsed dc magnetron sputtering with synchronous and asynchronous

bipolar pulses,” Thin Solid Films, vol. 520, no. 1, pp. 166–173, 2011.

[63] Y. Birol, “Sliding wear of CrN, AlCrN and AlTiN coated AISI H13 hot work tool steels in

aluminium extrusion,” Tribol. Int., vol. 57, no. 0, pp. 101–106, 2013.

[64] M. Bartosik, R. Daniel, C. Mitterer, and J. Keckes, “Thermally-induced formation of

hexagonal AlN in AlCrN hard coatings on sapphire: Orientation relationships and

residual stresses,” Surf. Coat. Technol., vol. 205, no. 5, pp. 1320–1323, 2010.

[65] Y. Birol and D. İsler, “Response to thermal cycling of CAPVD (Al,Cr)N-coated hot work

tool steel,” Surf. Coat. Technol., vol. 205, no. 2, pp. 275–280, 2010.

[66] Y. Chunyan, T. Linhai, W. Yinghui, W. Shebin, L. Tianbao, and X. Bingshe, “The effect of

substrate bias voltages on impact resistance of CrAlN coatings deposited by modified ion

beam enhanced magnetron sputtering,” Appl. Surf. Sci., vol. 255, no. 7, pp. 4033–4038,

2009.



50

[67] Y. Lv, L. Ji, X. Liu, H. Li, H. Zhou, and J. Chen, “Influence of substrate bias voltage on

structure and properties of the CrAlN films deposited by unbalanced magnetron

sputtering,” Appl. Surf. Sci., vol. 258, no. 8, pp. 3864–3870, 2012.

[68] T. Dudziak, P. K. Datta, P. H. Mayrhofer, and F. Rovere, “High Temperature Oxidation

Resistance of CrAlYN-Coated Ti45Al8Nb,” Oxid. Met., vol. 75, no. 5–6, pp. 359–376, 2011.

[69] M. Zhu, M. Li, and Y. Zhou, “Oxidation resistance of Cr1−xAlxN (0.18 ≤ x ≤ 0.47) coatings

on K38G superalloy at 1000–1100 °C in air,” Surf. Coat. Technol., vol. 201, no. 6, pp. 2878–

2886, 2006.

[70] F. Rovere, P. H. Mayrhofer, A. Reinholdt, J. Mayer, and J. M. Schneider, “The effect of

yttrium incorporation on the oxidation resistance of Cr–Al–N coatings,” Surf. Coat.

Technol., vol. 202, no. 24, pp. 5870–5875, 2008.

[71] P. Steyer, D. Pilloud, J. F. Pierson, J.-P. Millet, M. Charnay, B. Stauder, and P. Jacquot,

“Oxidation resistance improvement of arc-evaporated TiN hard coatings by silicon

addition,” Surf. Coat. Technol., vol. 201, no. 7, pp. 4158–4162, 2006.

[72] T. Schmitt, P. Steyer, J. Fontaine, N. Mary, C. Esnouf, M. O’Sullivan, and F. Sanchette,

“Cathodic arc deposited (Cr,Six)N coatings: From solid solution to nanocomposite

structure,” Surf. Coat. Technol., vol. 213, pp. 117–125, 2012.

[73] T. C. Rojas, S. El Mrabet, S. Domínguez-Meister, M. Brizuela, A. García-Luis, and J. C.

Sánchez-López, “Chemical and microstructural characterization of (Y or Zr)-doped

CrAlN coatings,” Surf. Coat. Technol., vol. 211, pp. 104–110, 2012.

[74] H. Scheerer and C. Berger, “Wear Mechanisms of (Cr,Al,Y)N PVD Coatings at Elevated

Temperatures,” Plasma Process. Polym., vol. 6, no. S1, pp. S157–S161, 2009.

[75] M. Moser, P. H. Mayrhofer, and H. Clemens, “Oxidation Resistance and Ductility of a

Coated γ-TiAl Based Alloy,” BHM Berg- Hüttenmänn. Monatshefte, vol. 153, no. 7, pp. 268–

272, 2008.

[76] S. Mato, G. Alcalá, M. Brizuela, R. Escobar Galindo, F. J. Pérez, and J. C. Sánchez-López,

“Long-term high temperature oxidation of CrAl(Y)N coatings in steam atmosphere,”

Corros. Sci., vol. 80, pp. 453–460, 2014.

[77] M. Moser, D. Kiener, C. Scheu, and P. H. Mayrhofer, “Influence of Yttrium on the

Thermal Stability of Ti-Al-N Thin Films,” Materials, vol. 3, no. 3, pp. 1573–1592, 2010.

[78] F. Rovere and P. H. Mayrhofer, “Thermal stability and thermo-mechanical properties of

magnetron sputtered Cr-Al-Y-N coatings,” J. Vac. Sci. Technol. A, vol. 26, no. 1, pp. 29–35,

2008.

[79] M. Brizuela, A. García-Luis, P. Corengia, D. González-Santamaría, R. Muñoz, and J.-J.

González, “Microstructural, Mechanical and Tribological Properties of CrA1YN Coatings

Deposited by Magnetron Sputtering,” Plasma Process. Polym., vol. 6, no. S1, pp. S162–S167,

2009.

[80] F. Rovere and P. H. Mayrhofer, “Impact of yttrium on structure and mechanical

properties of Cr–Al–N thin films,” J. Vac. Sci. Technol. A, vol. 25, no. 5, pp. 1336–1340,

2007.

[81] T. Miyake, A. Kishimoto, and H. Hasegawa, “Tribological properties and oxidation

resistance of (Cr,Al,Y)N and (Cr,Al,Si)N films synthesized by radio-frequency magnetron

sputtering method,” Surf. Coat. Technol., vol. 205, Supplement 1, pp. S290–S294, 2010.

[82] L. Zhu, Y. Zhang, W. Ni, and Y. Liu, “The effect of yttrium on cathodic arc evaporated

Ti0.45Al0.55N coating,” Surf. Coat. Technol., vol. 214, pp. 53–58, 2013.

[83] K. Yamamoto, S. Kujime, and G. Fox-Rabinovich, “Effect of alloying element (Si,Y) on

properties of AIP deposited (Ti,Cr,Al)N coating,” Surf. Coat. Technol., vol. 203, no. 5–7, pp.

579–583, 2008.

[84] Z. B. Qi, Z. T. Wu, and Z. C. Wang, “Improved hardness and oxidation resistance for

CrAlN hard coatings with Y addition by magnetron co-sputtering,” Surf. Coat. Technol.


Chapter 1

51

[85] W. H. Wang, “Bulk Metallic Glasses with Functional Physical Properties,” Adv. Mater.,

vol. 21, no. 45, pp. 4524–4544, 2009.

[86] C. L. Chiang, J. P. Chu, F. X. Liu, P. K. Liaw, and R. A. Buchanan, “A 200 nm thick glass-

forming metallic film for fatigue-property enhancements,” Appl. Phys. Lett., vol. 88, no. 13,

pp. 131902–131902–3, 2006.

[87] M. Chen, “A brief overview of bulk metallic glasses,” NPG Asia Mater., vol. 3, no. 9, pp.

82–90, Sep. 2011.

[88] J. P. Chu, J. S. C. Jang, J. C. Huang, H. S. Chou, Y. Yang, J. C. Ye, Y. C. Wang, J. W. Lee, F.

X. Liu, P. K. Liaw, Y. C. Chen, C. M. Lee, C. L. Li, and C. Rullyani, “Thin film metallic

glasses: Unique properties and potential applications,” Thin Solid Films, vol. 520, no. 16,

pp. 5097–5122, 2012.

[89] J. C. Huang, J. P. Chu, and J. S. C. Jang, “Recent progress in metallic glasses in Taiwan,”

Intermetallics, vol. 17, no. 12, pp. 973–987, 2009.

[90] A. Inoue and A. Takeuchi, “Recent development and application products of bulk glassy

alloys,” Acta Mater., vol. 59, no. 6, pp. 2243–2267, 2011.

[91] N. Mattern, P. Jóvári, I. Kaban, S. Gruner, A. Elsner, V. Kokotin, H. Franz, B. Beuneu, and

J. Eckert, “Short-range order of Cu–Zr metallic glasses,” J. Alloys Compd., vol. 485, no. 1–2,

pp. 163–169, 2009.

[92] W. Buckel and R. Hilsch, “Zur Supraleitung von Kupfersulfid,” Z. Für Phys., vol. 128, no.

2, pp. 324–346, 1950.

[93] W. Klement, R. H. Willens, and P. Duwez, “Non-crystalline Structure in Solidified Gold–

Silicon Alloys,” Nature, vol. 187, no. 4740, pp. 869–870, 1960.

[94] H. W. Kui, A. L. Greer, and D. Turnbull, “Formation of bulk metallic glass by fluxing,”

Appl. Phys. Lett., vol. 45, no. 6, pp. 615–616, 1984.

[95] A. Inoue, “Stabilization of metallic supercooled liquid and bulk amorphous alloys,” Acta

Mater., vol. 48, no. 1, pp. 279–306, 2000.

[96] J. F. Löffler, “Bulk metallic glasses,” Intermetallics, vol. 11, no. 6, pp. 529–540, 2003.

[97] D. Kawase, A. P. Tsai, A. Inoue, and T. Masumoto, “Crystallization on supercooled liquid

in metallic Zr-Cu-Al glasses,” Appl. Phys. Lett., vol. 62, no. 2, p. 137, 1993.

[98] Q. Guo, J. H. Noh, P. K. Liaw, P. D. Rack, Y. Li, and C. V. Thompson, “Density change

upon crystallization of amorphous Zr–Cu–Al thin films,” Acta Mater., vol. 58, no. 10, pp.

3633–3641, 2010.

[99] Y. J. Sun, D. D. Qu, Y. J. Huang, K.-D. Liss, X. S. Wei, D. W. Xing, and J. Shen, “Zr–Cu–Ni–

Al bulk metallic glasses with superhigh glass-forming ability,” Acta Mater., vol. 57, no. 4,

pp. 1290–1299, février 2009.

[100] A. Cai, X. Xiong, Y. Liu, W. An, J. Tan, and Y. Pan, “Design of new Zr–Al–Ni–Cu bulk

metallic glasses,” J. Alloys Compd., vol. 468, no. 1–2, pp. 432–437, 2009.

[101] Y.-L. Gao, J. Shen, J.-F. Sun, G. Wang, D.-W. Xing, H.-Z. Xian, and B.-D. Zhou,

“Crystallization behavior of ZrAlNiCu bulk metallic glass with wide supercooled liquid

region,” Mater. Lett., vol. 57, no. 13–14, pp. 1894–1898, 2003.

[102] A. Peker and W. L. Johnson, “A highly processable metallic glass:

Zr41.2Ti13.8Cu12.5Ni10.0Be22.5,” Appl. Phys. Lett., vol. 63, no. 17, pp. 2342–2344, 1993.

[103] A. L. Greer, “Metallic Glasses,” Science, vol. 267, no. 5206, pp. 1947–1953, 1995.

[104] A. L. Greer, “Confusion by design,” Nature, vol. 366, no. 6453, pp. 303–304, 1993.

[105] W. L. Johnson, “Bulk metallic glasses — a new engineering material,” Curr. Opin. Solid

State Mater. Sci., no. 3, pp. 383–386.

[106] Y. D. Sun, Q. R. Chen, and G. Z. Li, “Enhanced glass forming ability and plasticity of Mg-

based bulk metallic glass by minor addition of Cd,” J. Alloys Compd., vol. 584, pp. 273–278,

2014.



52

[107] Z. P. Lu and C. T. Liu, “Glass Formation Criterion for Various Glass-Forming Systems,”

Phys. Rev. Lett., vol. 91, no. 11, p. 115505, 2003.

[108] Y. Zhang, D. Q. Zhao, M. X. Pan, and W. H. Wang, “Glass forming properties of Zr-based

bulk metallic alloys,” J. Non-Cryst. Solids, vol. 315, no. 1–2, pp. 206–210, 2003.

[109] Z. P. Lu and C. T. Liu, “A new glass-forming ability criterion for bulk metallic glasses,”

Acta Mater., vol. 50, no. 13, pp. 3501–3512, 2002.

[110] B. DONG, S. ZHOU, D. LI, C. LU, F. GUO, X. NI, and Z. LU, “A new criterion for

predicting glass forming ability of bulk metallic glasses and some critical discussions,”

Prog. Nat. Sci. Mater. Int., vol. 21, no. 2, pp. 164–172, 2011.

[111] W. L. Johnson and K. Samwer, “A universal criterion for plastic yielding of metallic

glasses with a (T/Tg) 2/3 temperature dependence,” Phys. Rev. Lett., vol. 95, no. 19, p.

195501, 2005.

[112] A. Cai, X. Xiong, Y. Liu, W. An, and J. Tan, “Artificial neural network modeling of

reduced glass transition temperature of glass forming alloys,” Appl. Phys. Lett., vol. 92, no.

11, p. 111909, 2008.

[113] W. H. Wang, P. Wen, D. q. Zhao, M. x. Pan, and R. J. Wang, “Relationship between glass

transition temperature and Debye temperature in bulk metallic glasses,” J. Mater. Res.,

vol. 18, no. 12, pp. 2747–2751, 2003.

[114] W. H. Wang, “Elastic moduli and behaviors of metallic glasses,” J. Non-Cryst. Solids, vol.

351, no. 16–17, pp. 1481–1485, 2005.

[115] V. N. Novikov and A. P. Sokolov, “Correlation of fragility and Poisson’s ratio: Difference

between metallic and nonmetallic glass formers,” Phys. Rev. B, vol. 74, no. 6, p. 064203,

2006.

[116] B. Yang, C. T. Liu, and T. G. Nieh, “Unified equation for the strength of bulk metallic

glasses,” Appl. Phys. Lett., vol. 88, no. 22, p. 221911, 2006.

[117] T. Scopigno, “Is the Fragility of a Liquid Embedded in the Properties of Its Glass?,”

Science, vol. 302, no. 5646, pp. 849–852, 2003.

[118] J. Lu, G. Ravichandran, and W. L. Johnson, “Deformation behavior of the

Zr41.2Ti13.8Cu12.5Ni10Be22.5 bulk metallic glass over a wide range of strain-rates and

temperatures,” Acta Mater., vol. 51, no. 12, pp. 3429–3443, 2003.

[119] H.-J. Jun, K. S. Lee, and Y. W. Chang, “Deformation behavior and formability of a Ti–Zr–

Ni–Be bulk metallic glass within supercooled liquid region,” Intermetallics, vol. 18, no. 8,

pp. 1537–1543, 2010.

[120] J. N. Mei, J. L. Soubeyroux, J. J. Blandin, J. S. Li, H. C. Kou, H. Z. Fu, and L. Zhou,

“Homogeneous deformation of Ti41.5Cu37.5Ni7.5Zr2.5Hf5Sn5Si1 bulk metallic glass in

the supercooled liquid region,” Intermetallics, vol. 19, no. 1, pp. 48–53, 2011.

[121] M. F. Ashby and A. L. Greer, “Metallic glasses as structural materials,” Scr. Mater., vol. 54,

no. 3, pp. 321–326, 2006.

[122] C. A. Schuh, T. C. Hufnagel, and U. Ramamurty, “Mechanical behavior of amorphous

alloys,” Acta Mater., vol. 55, no. 12, pp. 4067–4109, 2007.

[123] A. Inoue, B. Shen, H. Koshiba, H. Kato, and A. R. Yavari, “Cobalt-based bulk glassy alloy

with ultrahigh strength and soft magnetic properties,” Nat. Mater., vol. 2, no. 10, pp. 661–

663, 2003.

[124] T. Burgess and M. Ferry, “Nanoindentation of metallic glasses,” Mater. Today, vol. 12, no.

1–2, pp. 24–32, 2009.

[125] B. F. Lu, J. F. Li, L. T. Kong, and Y. H. Zhou, “Correlation between mechanical behavior

and glass forming ability of Zr–Cu metallic glasses,” Intermetallics, vol. 19, no. 7, pp. 1032–

1035, 2011.


Chapter 1

53

[126] P. J. Hsieh, S. C. Lin, H. C. Su, and J. S. C. Jang, “Glass forming ability and mechanical

properties characterization on Mg58Cu31Y11−xGdx bulk metallic glasses,” J. Alloys

Compd., vol. 483, no. 1–2, pp. 40–43, 2009.

[127] H. S. Chou, J. C. Huang, Y. H. Lai, L. W. Chang, X. H. Du, J. P. Chu, and T. G. Nieh,

“Amorphous and nanocrystalline sputtered Mg–Cu thin films,” J. Alloys Compd., vol. 483,

no. 1–2, pp. 341–345, 2009.

[128] S. Y. Kuan, H. S. Chou, and J. C. Huang, “Mechanical characteristics of Mg–Cu–Zr thin

film metallic glasses,” Surf. Coat. Technol. 231, 58-61, 2013

[129] S. Nowak, P. Ochin, A. Pasko, O. Maciejak, P. Aubert, and Y. Champion,

“Nanoindentation analysis of the mechanical behavior of Zr-based metallic glasses with

Sn, Ta and W additions,” J. Alloys Compd., vol. 483, no. 1–2, pp. 139–142, 2009.

[130] B.-G. Yoo, J.-H. Oh, Y.-J. Kim, K.-W. Park, J.-C. Lee, and J. Jang, “Nanoindentation

analysis of time-dependent deformation in as-cast and annealed Cu–Zr bulk metallic

glass,” Intermetallics, vol. 18, no. 10, pp. 1898–1901, 2010.

[131] B. Yang and T. G. Nieh, “Effect of the nanoindentation rate on the shear band formation

in an Au-based bulk metallic glass,” Acta Mater., vol. 55, no. 1, pp. 295–300, 2007.

[132] J.-J. Pang, M.-J. Tan, K. M. Liew, and C. Shearwood, “Nanoindentation study of size effect

and loading rate effect on mechanical properties of a thin film metallic glass

Cu49.3Zr50.7,” Phys. B Condens. Matter, vol. 407, no. 3, pp. 340–346, 2012.

[133] Y. H. Chen, J. C. Huang, L. Wang, and T. G. Nieh, “Effect of residual stresses on

nanoindentation creep behavior of Zr-based bulk metallic glasses,” Intermetallics, vol. 41,

pp. 58–62, 2013.

[134] Y. Huang, Y. L. Chiu, J. Shen, Y. Sun, and J. J. J. Chen, “Mechanical performance of

metallic glasses during nanoscratch tests,” Intermetallics, vol. 18, no. 5, pp. 1056–1061,

2010.

[135] J. Gu, M. Song, S. Ni, S. Guo, and Y. He, “Effects of annealing on the hardness and elastic

modulus of a Cu36Zr48Al8Ag8 bulk metallic glass,” Mater. Des., vol. 47, pp. 706–710,

2013.

[136] U. Ramamurty, S. Jana, Y. Kawamura, and K. Chattopadhyay, “Hardness and plastic

deformation in a bulk metallic glass,” Acta Mater., vol. 53, no. 3, pp. 705–717, 2005.

[137] Z. T. Wang, K. Y. Zeng, and Y. Li, “The correlation between glass formation and hardness

of the amorphous phase,” Scr. Mater., vol. 65, no. 9, pp. 747–750, 2011.

[138] F. Zeng, Y. Gao, L. Li, D. M. Li, and F. Pan, “Elastic modulus and hardness of Cu–Ta

amorphous films,” J. Alloys Compd., vol. 389, no. 1–2, pp. 75–79, 2005.

[139] Z. . Zhang, J. Eckert, and L. Schultz, “Difference in compressive and tensile fracture

mechanisms of Zr59Cu20Al10Ni8Ti3 bulk metallic glass,” Acta Mater., vol. 51, no. 4, pp.

1167–1179, février 2003.

[140] H. J. Pei, C. J. Lee, X. H. Du, Y. C. Chang, and J. C. Huang, “Tension behavior of metallic

glass coating on Cu foil,” Mater. Sci. Eng. A, vol. 528, no. 24, pp. 7317–7322, 2011.

[141] H. J. Pei, S. Y. Kuan, M. C. Liu, and J. C. Huang, “Tensile behavior of

amorphous/nanocrystalline ZrCu/Cu multilayered films with graded interfaces,”

Intermetallics, vol. 31, pp. 191–195, 2012.

[142] G. Chen, J. L. Cheng, and C. T. Liu, “Large-sized Zr-based bulk-metallic-glass composite

with enhanced tensile properties,” Intermetallics, vol. 28, pp. 25–33, 2012.

[143] L. S. Huo, H. Y. Bai, X. K. Xi, D. W. Ding, D. Q. Zhao, W. H. Wang, R. J. Huang, and L. F.

Li, “Tensile properties of ZrCu-based bulk metallic glasses at ambient and cryogenic

temperatures,” J. Non-Cryst. Solids, vol. 357, no. 16–17, pp. 3088–3093, 2011.

[144] W. Ma, H. Kou, J. Li, H. Chang, and L. Zhou, “Effect of strain rate on compressive

behavior of Ti-based bulk metallic glass at room temperature,” J. Alloys Compd., vol. 472,

no. 1–2, pp. 214–218, 2009.



54

[145] C.-C. Wang, Y.-W. Mao, Z.-W. Shan, M. Dao, J. Li, J. Sun, E. Ma, and S. Suresh, “Real-

time, high-resolution study of nanocrystallization and fatigue cracking in a cyclically

strained metallic glass,” Proc. Natl. Acad. Sci., vol. 110, no. 49, pp. 19725–19730, 2013.

[146] Q. P. Cao, Y. Ma, C. Wang, X. D. Wang, and J. Z. Jiang, “Effect of temperature and strain

rate on deformation behavior in metallic glassy films,” Thin Solid Films. 561, 60-69, 2014

[147] Y. Huang, W. Zheng, H. Fan, D. Wang, J. Shen, and J. Mi, “The effects of annealing on the

microstructure and the dynamic mechanical strength of a ZrCuNiAl bulk metallic glass,”


[148] A. L. Greer, “Metallic glasses: Damage tolerance at a price,” Nat. Mater., vol. 10, no. 2, pp.

88–89, 2011.

[149] N. Nagendra, U. Ramamurty, T. T. Goh, and Y. Li, “Effect of crystallinity on the impact

toughness of a La-based bulk metallic glass,” Acta Mater., vol. 48, no. 10, pp. 2603–2615,

2000.

[150] F. H. Dalla Torre, A. Dubach, J. Schällibaum, and J. F. Löffler, “Shear striations and

deformation kinetics in highly deformed Zr-based bulk metallic glasses,” Acta Mater., vol.

56, no. 17, pp. 4635–4646, 2008.

[151] D. Xu, B. Lohwongwatana, G. Duan, W. L. Johnson, and C. Garland, “Bulk metallic glass

formation in binary Cu-rich alloy series – Cu100−xZrx (x=34, 36, 38.2, 40 at.%) and

mechanical properties of bulk Cu64Zr36 glass,” Acta Mater., vol. 52, no. 9, pp. 2621–2624,

mai 2004.

[152] X. Cui, F. Q. Zu, W. J. Zhang, Z. Z. Wang, and X. Y. Li, “Phase competition of Cu64Zr36

and its effect on glass forming ability of the alloy,” Cryst. Res. Technol., vol. 48, no. 1, pp.

11–15, 2013.

[153] Y. T. Cheng, T. H. Hung, J. C. Huang, P. J. Hsieh, and J. S. C. Jang, “Thermal stability and

crystallization kinetics of Mg–Cu–Y–B quaternary alloys,” Mater. Sci. Eng. A, vol. 449–451,

pp. 501–505, 2007.

[154] D. Okai, Y. Shimizu, N. Hirano, T. Fukami, T. Yamasaki, and A. Inoue, “Isothermal

crystallization in supercooled liquid state for Ca50Mg22.5Cu27.5 metallic glass,” J. Alloys

Compd., vol. 504, Supplement 1, pp. S247–S250, 2010.

[155] H.-R. Wang, Y.-L. Gao, Y.-F. Ye, G.-H. Min, Y. Chen, and X.-Y. Teng, “Crystallization

kinetics of an amorphous Zr–Cu–Ni alloy: calculation of the activation energy,” J. Alloys

Compd., vol. 353, no. 1–2, pp. 200–206, 2003.

[156] J. C. Qiao and J. M. Pelletier, “Crystallization kinetics in Cu46Zr45Al7Y2 bulk metallic

glass by differential scanning calorimetry (DSC),” J. Non-Cryst. Solids, vol. 357, no. 14, pp.

2590–2594, 2011.

[157] K. G. Prashanth, S. Scudino, K. B. Surreddi, M. Sakaliyska, B. S. Murty, and J. Eckert,

“Crystallization kinetics of Zr65Ag5Cu12.5Ni10Al7.5 glassy powders produced by ball

milling of pre-alloyed ingots,” Mater. Sci. Eng. A, vol. 513–514, pp. 279–285, 2009.

[158] D. R. dos Santos and D. S. dos Santos, “Crystallization kinetics of Fe-B based amorphous

alloys studied in-situ using X-rays diffraction and differential scanning calorimetry,”

Mater. Res., vol. 4, no. 1, pp. 47–51, 2001.

[159] K. B. Surreddi, S. Scudino, M. Sakaliyska, K. G. Prashanth, D. J. Sordelet, and J. Eckert,

“Crystallization behavior and consolidation of gas-atomized Al84Gd6Ni7Co3 glassy

powder,” J. Alloys Compd., vol. 491, no. 1–2, pp. 137–142, 2010.

[160] C. W. Price, “Use of Kolmogorov-Johnson-Mehl-Avrami kinetics in recrystallization of

metals and crystallization of metallic glasses,” Acta Metall. Mater., vol. 38, no. 5, pp. 727–

738, mai 1990.

[161] R. T. Yang and M. Steinberg, “Reaction kinetics and differential thermal analysis,” J. Phys.

Chem., vol. 80, no. 9, pp. 965–968, 1976.


Chapter 1

55

[162] M. Stoica, S. Kumar, S. Roth, S. Ram, J. Eckert, G. Vaughan, and A. R. Yavari,

“Crystallization kinetics and magnetic properties of Fe66Nb4B30 bulk metallic glass,” in

Journal of alloys and compounds, vol. 483, pp. 632–637.

[163] F. Xu, Y. Du, P. Gao, Z. Han, G. Chen, S. Wang, and J. Jiang, “Crystallization of melt-spun

Mg63Ni22Pr15 amorphous alloy ribbon,” J. Alloys Compd., vol. 441, no. 1–2, pp. 76–80,

2007.

[164] J. T. Zhang, W. M. Wang, H. J. Ma, G. H. Li, R. Li, and Z. H. Zhang, “Isochronal and

isothermal crystallization kinetics of amorphous Fe-based alloys,” Thermochim. Acta, vol.

505, no. 1–2, pp. 41–46, 2010.

[165] X. Wu, L. Meng, W. Zhao, Z. Suo, Y. Si, and K. Qiu, “Crystallization Kinetics of Misch

Metal Based Bulk Metallic Glasses,” J. Rare Earths, vol. 25, no. 2, pp. 189–193, 2007.

[166] M. D. Demetriou, M. E. Launey, G. Garrett, J. P. Schramm, D. C. Hofmann, W. L. Johnson,

and R. O. Ritchie, “A damage-tolerant glass,” Nat. Mater., vol. 10, no. 2, pp. 123–128, 2011.

[167] Q. Wang, S. Gravier, J. J. Blandin, J. M. Pelletier, and J. Lu, “Deformation and

crystallization of a Zr41.2Ti13.8Cu12.5Ni10Be22.5 bulk metallic glass in the supercooled

liquid region,” Mater. Sci. Eng. A, vol. 435–436, pp. 405–411, 2006.

[168] Y. Liu, S. Hata, K. Wada, and A. Shimokohbe, “Thermal, Mechanical and Electrical

Properties of Pd-Based Thin-Film Metallic Glass,” Jpn. J. Appl. Phys., vol. 40, no. Part 1,

No. 9A, pp. 5382–5388, 2001.

[169] D. H. Bae, H. K. Lim, S. H. Kim, D. H. Kim, and W. T. Kim, “Mechanical behavior of a

bulk Cu–Ti–Zr–Ni–Si–Sn metallic glass forming nano-crystal aggregate bands during

deformation in the supercooled liquid region,” Acta Mater., vol. 50, no. 7, pp. 1749–1759,

2002.

[170] T. Fukami, K. Okabe, D. Okai, T. Yamasaki, T. Zhang, and A. Inoue, “Crystal growth and

time evolution in Zr–Al–Cu–Ni glassy metals in supercooled liquid,” Mater. Sci. Eng. B,

vol. 111, no. 2–3, pp. 189–196, 2004.

[171] K. S. Lee, Y. M. Jo, and Y.-S. Lee, “Crystallization and high-temperature deformation

behavior of Cu49Zr45Al6 bulk metallic glass within supercooled liquid region,” J. Non-

Cryst. Solids, vol. 376, pp. 145–151, 2013.

[172] S. Wei, B. Ding, T. Lei, and Z. Hu, “Crystallization of amorphous Zr65Cu27Al8 alloy with

wide supercooled liquid region,” Mater. Lett., vol. 37, no. 4–5, pp. 263–267, 1998.

[173] S. Zhang, T. Ichitsubo, Y. Yokoyama, T. Yamamoto, E. Matsubara, and A. Inoue,

“Crystallization Behavior and Structural Stability of Zr50Cu40Al10 Bulk Metallic Glass,”

Mater. Trans., vol. 50, no. 6, pp. 1340–1345.

[174] G. Xie, F. Qin, S. Zhu, and D. V. Louzguine-Lugzin, “Corrosion behaviour of porous Ni-

free Ti-based bulk metallic glass produced by spark plasma sintering in Hanks’ solution,”


[175] J. Fornell, N. Van Steenberge, A. Varea, E. Rossinyol, E. Pellicer, S. Suriñach, M. D. Baró,

and J. Sort, “Enhanced mechanical properties and in vitro corrosion behavior of

amorphous and devitrified Ti40Zr10Cu38Pd12 metallic glass,” J. Mech. Behav. Biomed.

Mater., vol. 4, no. 8, pp. 1709–1717, 2011.

[176] N. Mattern, A. Schöps, U. Kühn, J. Acker, O. Khvostikova, and J. Eckert, “Structural

behavior of CuxZr100−x metallic glass (x = 35−70),” J. Non-Cryst. Solids, vol. 354, no. 10–11,

pp. 1054–1060, 2008.

[177] J. S. C. Jang, Y. W. Chen, L. J. Chang, H. Z. Cheng, C. C. Huang, and C. Y. Tsau,

“Crystallization and fracture behavior of the Zr65−xAl7.5Cu17.5Ni10Six bulk amorphous

alloys,” Mater. Chem. Phys., vol. 89, no. 1, pp. 122–129, 2005.

[178] H. J. Leamy and A. G. Dirks, “The microstructure of amorphous rare-earth/transition-

metal thin films,” J. Phys. Appl. Phys., vol. 10, no. 8, pp. L95–L98, 1977.



56

[179] G. A. Almyras, G. M. Matenoglou, P. Komninou, C. Kosmidis, P. Patsalas, and G. A.

Evangelakis, “On the deposition mechanisms and the formation of glassy Cu–Zr thin

films,” J. Appl. Phys., vol. 107, no. 8, pp. 084313–084313–6, 2010.

[180] R. Yamauchi, S. Hata, J. Sakurai, and A. Shimokohbe, “Combinatorial Search for Low

Resistivity Pd–Cu–Si Thin Film Metallic Glass Compositions,” Jpn. J. Appl. Phys., vol. 45,

no. 7, pp. 5911–5919, 2006.

[181] Y. P. Deng, Y. F. Guan, J. D. Fowlkes, S. Q. Wen, F. X. Liu, G. M. Pharr, P. K. Liaw, C. T.

Liu, and P. D. Rack, “A combinatorial thin film sputtering approach for synthesizing and

characterizing ternary ZrCuAl metallic glasses,” Intermetallics, vol. 15, no. 9, pp. 1208–

1216, 2007.

[182] H. S. Chou, J. C. Huang, and L. W. Chang, “Mechanical properties of ZrCuTi thin film

metallic glass with high content of immiscible tantalum,” Surf. Coat. Technol., vol. 205, no.

2, pp. 587–590, 2010.

[183] H. S. Chou, J. C. Huang, L. W. Chang, and T. G. Nieh, “Structural relaxation and

nanoindentation response in Zr–Cu–Ti amorphous thin films,” Appl. Phys. Lett., vol. 93,

no. 19, pp. 191901–191901–3, Nov. 2008.

[184] J. C. Ye, J. P. Chu, Y. C. Chen, Q. Wang, and Y. Yang, “Hardness, yield strength, and

plastic flow in thin film metallic-glass,” J. Appl. Phys., vol. 112, no. 5, pp. 053516–053516–9,

2012.

[185] F. X. Liu, F. Q. Yang, Y. F. Gao, W. H. Jiang, Y. F. Guan, P. D. Rack, O. Sergic, and P. K.

Liaw, “Micro-scratch study of a magnetron-sputtered Zr-based metallic-glass film,” Surf.

Coat. Technol., vol. 203, no. 22, pp. 3480–3484, 2009.

[186] P. H. Tsai, Y. Z. Lin, J. B. Li, S. R. Jian, J. S. C. Jang, C. Li, J. P. Chu, and J. C. Huang,

“Sharpness improvement of surgical blade by means of ZrCuAlAgSi metallic glass and

metallic glass thin film coating,” Intermetallics, vol. 31, no. 0, pp. 127–131, 2012.

[187] B. E. Schuster, Q. Wei, T. C. Hufnagel, and K. T. Ramesh, “Size-independent strength and

deformation mode in compression of a Pd-based metallic glass,” Acta Mater., vol. 56, no.

18, pp. 5091–5100, 2008.

[188] T.-P. Hsiao, Y.-C. Yang, J.-W. Lee, C.-L. Li, and J. P. Chu, “The microstructure and

mechanical properties of nitrogen and boron contained ZrCuAlNi thin film metallic glass

composites,” Surf. Coat. Technol. 237, 276-283, 2013

[189] J.-H. Chu, H.-W. Chen, Y.-C. Chan, J.-G. Duh, J.-W. Lee, and J. S.-C. Jang, “Modification of

structure and property in Zr-based thin film metallic glass via processing temperature

control,” Thin Solid Films. 561, 38-42, 2014

[190] L. He, J. P. Chu, C.-L. Li, C.-M. Lee, Y.-C. Chen, P. K. Liaw, and P. M. Voyles, “Effects of

annealing on the compositional heterogeneity and structure in zirconium-based bulk

metallic glass thin films,” Thin Solid Films. 561, 87-92, 2014

[191] Y. H. Liu, T. Fujita, A. Hirata, S. Li, H. W. Liu, W. Zhang, A. Inoue, and M. W. Chen,

“Deposition of multicomponent metallic glass films by single-target magnetron

sputtering,” Intermetallics, vol. 21, no. 1, pp. 105–114, 2012.

[192] P. Murali and U. Ramamurty, “Embrittlement of a bulk metallic glass due to sub-Tg

annealing,” Acta Mater., vol. 53, no. 5, pp. 1467–1478, 2005.

[193] G. S. Yu, J. G. Lin, M. Mo, X. F. Wang, F. H. Wang, and C. E. Wen, “Effect of relaxation on

pressure sensitivity index in a Zr-based metallic glass,” Mater. Sci. Eng. A, vol. 460–461,

pp. 58–62, 2007.

[194] K. K. Song, S. Pauly, Y. Zhang, B. A. Sun, J. He, G. Z. Ma, U. Kühn, and J. Eckert,

“Thermal stability and mechanical properties of Cu46Zr46Ag8 bulk metallic glass and its

composites,” Mater. Sci. Eng. A, vol. 559, pp. 711–718, 2013.


Chapter 1

57

[195] M. Ghidelli, A. Volland, J.-J. Blandin, T. Pardoen, J.-P. Raskin, F. Mompiou, P. Djemia,

and S. Gravier, “Exploring the mechanical size effects in Zr65Ni35 thin film metallic

glasses,” J. Alloys Compd. 615, S90-S92, 2014

[196] H. Peng, S. S. Li, Y. P. Qi, and T. Y. Huang, “Mg–Ni–Gd–Ag bulk metallic glass with

improved glass-forming ability and mechanical properties,” Intermetallics, vol. 19, no. 7,

pp. 829–832, 2011.

[197] J. Eckert, J. Das, S. Pauly, and C. Duhamel, “Mechanical properties of bulk metallic glasses

and composites,” J. Mater. Res., vol. 22, no. 02, pp. 285–301, 2007.

[198] R. Banerjee, S. Bose, A. Genc, and P. Ayyub, “The microstructure and electrical transport

properties of immiscible copper-niobium alloy thin films,” J. Appl. Phys., vol. 103, no. 3,

pp. 033511–033511–7, 2008.

[199] M. Ghidelli, S. Gravier, J.-J. Blandin, T. Pardoen, J.-P. Raskin, and F. Mompiou,

“Compositional-induced structural change in ZrxNi100−x thin film metallic glasses,” J.

Alloys Compd. 615, S348-S351, 2014

[200] S. Nagel and J. Tauc, “Nearly-Free-Electron Approach to the Theory of Metallic Glass

Alloys,” Phys. Rev. Lett., vol. 35, no. 6, pp. 380–383, 1975.

[201] C. J. Chen, J. C. Huang, H. S. Chou, Y. H. Lai, L. W. Chang, X. H. Du, J. P. Chu, and T. G.

Nieh, “On the amorphous and nanocrystalline Zr–Cu and Zr–Ti co-sputtered thin films,”

J. Alloys Compd., vol. 483, no. 1–2, pp. 337–340, 2009.

[202] J. Dudonis, R. Bručas, and A. Miniotas, “Synthesis of amorphous Zr-Cu alloys by

magnetron co-sputtering,” Thin Solid Films, vol. 275, no. 1–2, pp. 164–167, 1996.

[203] T. Minemura, J. J. van den Broek, and J. L. C. Daams, “Formation and thermal stability of

amorphous Cu‐Zr thin films deposited by coevaporation,” J. Appl. Phys., vol. 63, no. 9, pp.

4426–4430, 1988.

[204] K. Kondoh, J. Fujita, J. Umeda, and T. Serikawa, “Estimation of Compositions of Zr-Cu

Binary Sputtered Film and Its Characterization,” Res. Lett. Mater. Sci., vol. 2008, pp. 1–5,

2008.

[205] J. Qin, X. Yong, Z. Xin-Yu, L. Gong, L. Li-Xin, X. Zhe, M. Ming-Zhen, and L. Ri-Ping, “Zr-

Cu Amorphous Films Prepared by Magnetron Co-sputtering Deposition of Pure Zr and

Cu,” Chin. Phys. Lett., vol. 26, no. 8, p. 086109, 2009.

[206] J.-H. Shin, D.-I. Kim, K. Cho, H. Suematsu, K. H. Kim, and R. Nowak, “Plasticity-induced

nanocrystallization in a Zr65Cu35 thin film metallic glass,” J. Non-Cryst. Solids, vol. 362,

pp. 65–68, 2013.

[207] P. Coddet, F. Sanchette, J. C. Rousset, O. Rapaud, and C. Coddet, “On the elastic modulus

and hardness of co-sputtered Zr–Cu–(N) thin metal glass films,” Surf. Coat. Technol., vol.

206, no. 17, pp. 3567–3571, 2012.

[208] J. Musil and P. Zeman, “Structure and microhardness of magnetron sputtered ZrCu and

ZrCu-N films,” Vacuum, vol. 52, no. 3, pp. 269–275, 1999.

[209] L. Xie, P. Brault, A.-L. Thomann, and L. Bedra, “Molecular dynamic simulation of binary

ZrxCu100−x metallic glass thin film growth,” Appl. Surf. Sci., vol. 274, pp. 164–170, juin

2013.

[210] F. Wang, J. M. Li, P. Huang, W. L. Wang, T. J. Lu, and K. W. Xu, “Nanoscale creep

deformation in Zr-based metallic glass,” Intermetallics, vol. 38, pp. 156–160, 2013.


metallic glass with high content of immiscible tantalum,” Surf. Coat. Technol., vol. 205, no.

2, pp. 587–590, 2010.

[212] C.-Y. Chuang, J.-W. Lee, C.-L. Li, and J. P. Chu, “Mechanical properties study of a

magnetron-sputtered Zr-based thin film metallic glass,” Surf. Coat. Technol., no. 215, 312-

321, 2013



58

[213] H. S. Chou, X. H. Du, C. J. Lee, and J. C. Huang, “Enhanced mechanical properties of

multilayered micropillars of amorphous ZrCuTi and nanocrystalline Ta layers,”


[214] J. P. Chu, J. C. Huang, J. S. C. Jang, Y. C. Wang, and P. K. Liaw, “Thin film metallic

glasses: Preparations, properties, and applications,” JOM, vol. 62, no. 4, pp. 19–24,. 2010.

[215] G. Kumar, A. Desai, and J. Schroers, “Bulk Metallic Glass: The Smaller the Better,” Adv.

Mater., vol. 23, no. 4, pp. 461–476, 2011.

[216] J. O. Noyce, H. Michels, and C. W. Keevil, “Potential use of copper surfaces to reduce

survival of epidemic meticillin-resistant Staphylococcus aureus in the healthcare

environment,” J. Hosp. Infect., vol. 63, no. 3, pp. 289–297, 2006.

[217] P.-T. Chiang, G.-J. Chen, S.-R. Jian, Y.-H. Shih, J. S.-C. Jang, and C.-H. Lai, “Surface

Antimicrobial Effects of Zr61Al7.5Ni10Cu17.5Si4 Thin Film Metallic Glasses on

Escherichia coli, Staphylococcus aureus, Pseudomonas aeruginosa, Acinetobacter

baumannii and Candida albicans,” Fooyin J. Health Sci., vol. 2, no. 1, pp. 12–20, 2010.

[218] J. P. Chu, T.-Y. Liu, C.-L. Li, C.-H. Wang, J. S. C. Jang, M.-J. Chen, S.-H. Chang, and W.-C.

Huang, “Fabrication and characterizations of thin film metallic glasses: Antibacterial

property and durability study for medical application,” Thin Solid Films, vol 561, 102-107,

561

[219] H.-W. Chen, K.-C. Hsu, Y.-C. Chan, J.-G. Duh, J.-W. Lee, J. S.-C. Jang, and G.-J. Chen,

“Antimicrobial properties of Zr–Cu–Al–Ag thin film metallic glass,” Thin Solid Films, vol

561, 98 – 101, 2014

[220] N. Kaushik, P. Sharma, S. Ahadian, A. Khademhosseini, M. Takahashi, A. Makino, S.

Tanaka, and M. Esashi, “Metallic glass thin films for potential biomedical applications,” J.

Biomed. Mater. Res. B Appl. Biomater., p. n/a–n/a, Mar. 2014.

[221] D. Zander, B. Heisterkamp, and I. Gallino, “Corrosion resistance of Cu–Zr–Al–Y and Zr–

Cu–Ni–Al–Nb bulk metallic glasses,” J. Alloys Compd., vol. 434–435, pp. 234–236, 2007.

[222] S. F. Chen, S. L. Lin, J. K. Chen, and Y. L. Lin, “Thermal stability and corrosion behavior

of Cu–Zr–Al–Y bulk metallic glass,” Intermetallics, vol. 18, no. 10, pp. 1954–1957, 2010.

[223] B. Liu and L. Liu, “The effect of microalloying on thermal stability and corrosion

resistance of Cu-based bulk metallic glasses,” Mater. Sci. Eng. A, vol. 415, no. 1–2, pp. 286–

290, 2006.

[224] C.-F. Hsu, W. Kai, H.-M. Lin, C.-K. Lin, and P.-Y. Lee, “Fabrication and corrosion

behavior of Ti-based bulk metallic glass composites containing carbon nanotubes,” J.

Alloys Compd., vol. 504, Supplement 1, pp. S176–S179, 2010.

[225] C. Li, W. Chen, Q. Jiang, L. Wang, and D. Luo, “Corrosion resistance of Ti-based metallic

glass coating in concentrated nitric acid,” Mater. Chem. Phys., vol. 143, no. 3, pp. 900–903,

2014.

[226] H. B. Lu, L. C. Zhang, A. Gebert, and L. Schultz, “Pitting corrosion of Cu–Zr metallic

glasses in hydrochloric acid solutions,” J. Alloys Compd., vol. 462, no. 1–2, pp. 60–67, 2008.

[227] H. Bala and S. Szymura, “Acid corrosion of amorphous and crystalline Cu-Zr alloys,”

Appl. Surf. Sci., vol. 35, no. 1, pp. 41–51, 1988.

[228] M. Naka, K. Hashimoto, and T. Masumoto, “Corrosion behavior of amorphous and

crystalline Cu50Ti50 and Cu50Zr50 alloys,” J. Non-Cryst. Solids, vol. 30, no. 1, pp. 29–36,

1978.

[229] B.-P. Zhang, A. Kawashima, H. Habazaki, K. Asami, and K. Hashimoto, “The effect of

structural heterogeneity on the pitting corrosion behavior of melt-spun amorphous Ni-Zr

alloys,” Corros. Sci., vol. 39, no. 10–11, pp. 2005–2018, 1997.

[230] A. Jukic, J. Piljac, and M. Metikoš-Hukovic, “Electrocatalytic behavior of the Co33Zr67

metallic glass for hydrogen evolution,” J. Mol. Catal. Chem., vol. 166, no. 2, pp. 293–302,

février 2001.


Chapter 1

59

[231] C.-Y. Chuang, Y.-C. Liao, J.-W. Lee, C.-L. Li, J. P. Chu, and J.-G. Duh, “Electrochemical

characterization of Zr-based thin film metallic glass in hydrochloric aqueous solution,”

Thin Solid Films, vol 529, 338 – 341, 2013

[232] C. W. Chu, J. S. C. Jang, S. M. Chiu, and J. P. Chu, “Study of the characteristics and

corrosion behavior for the Zr-based metallic glass thin film fabricated by pulse magnetron

sputtering process,” Thin Solid Films, vol. 517, no. 17, pp. 4930–4933, 2009.

[233] B. Subramanian, S. Yugeswaran, A. Kobayashi, and M. Jayachandran, “Fabrication of

amorphous Zr48Cu36Al8Ag8 thin films by ion beam sputtering and their corrosion

behavior in SBF for bio implants,” J. Alloys Compd., vol. 572, pp. 163–169, 2013.

[234] B. Zhang, Y. G. Chen, and H. B. Guo, “Electrochemical Behavior of Cu-Hf-Al Amorphous

Films,” ECS Electrochem. Lett., vol. 3, no. 1, pp. C1–C3, 2014.


Content

2.1. Synthesis of thin films ....................................................................................................................... 60 2.2. Thin films characterization ............................................................................................................... 67 2.3. Functional characterization .............................................................................................................. 74 Chapter summary ..................................................................................................................................... 79 List of references ....................................................................................................................................... 80

Chapter 2

Description of the

experimental details and techniques

Overview

The aim of this work was the development of two kinds of innovative films and the

improvement of their characteristics (mechanical properties, resistance against severe

environments…). It was, thus, necessary to control the synthesis process but also to

characterize the films before, during or after testing their properties.

This chapter presents all the experimental techniques used during this study and is

divided in three main parts devoted namely to the films’ synthesis, to the metallurgical

characterization of layers and to their functional characterization.


Description of the experimental details and techniques

60

2.1. Synthesis of thin films

2.1.1. Short history

A new state of the matter, the plasma, has been explored by M. Faraday at the beginning of

the 19th century from the discharge into a tube of glass under vacuum [1]. The term and

process which today is called “sputtering” is credited to W.R. Grove [2] after he successfully

sputtered oxides species onto polished silver, copper or platinum plates in 1852. Later on,

his pioneering results encouraged other researchers to carry on the development of the thin

films deposition techniques. In 1858 A. W. Wright [3] used an “electrical deposition

apparatus” to create mirrors. A significant progress was the “electron trap” model,

proposed by Penning in the 1930s, which results from combination of electric and magnetic

fields known as Penning’s discharge. This study was quite essential for further

improvement of the magnetron sputtering sources configurations as more developed by

Mattox et al. in 1970 [4].

Among the first systems of deposition using the sputtering are the dielectric films sputtered

by Davidse et al in 1966 [5] and the Cermet resistor films using metals (i.e., Ni, Al, Cu)

prepared by Hohenstein [6]. Focusing the scientific efforts to develop such methods had led

to reactively sputtered hard coatings on tools in the mid-1970s, and to be commercially

available from the early 1980s.

2.1.2. Deposition techniques

Basic concepts

The deposition techniques can be generally classified taking into account the nature of the

deposition process, through either a chemical reaction (chemical vapor deposition) or a

physical mechanism (physical vapor deposition). There are numerous schemes that classify

these techniques from different points of view. For instance, the chemical approach given by

Campbell [7] considers the overlap between physical and chemical methods (i.e.,

evaporation and thermal growth). A deeper classification was proposed by Schiller et al. [8]

based on the vacuum deposition techniques, specifically physical vapor deposition (PVD)

and chemical vapor deposition (CVD). Bunshah and Mattox [9] proposed a different

classification where the concern is the dimensions of the depositions species (i.e.,

atoms/molecules, liquid droplets) in an atomistic deposition process.

Figure 2.1. illustrates the widely accepted classification adopted by Schiller.

Figure 2.1. Classification of vacuum deposition techniques [10]


Chapter 2

61

In the following part the focus will be on the PVD methods since they were indeed used for

the deposition of our films as highlighted in figure 2.1. The most common and exploited

PVD practices are evaporation and sputtering [10] which could be described in three general

steps as described below in figure 2.2.

Figure 2.2. Schematic representation of a PVD process

For the current thesis, investigated films have been prepared by PVD magnetron sputtering

technique in a laboratory-sized deposition system (metallic glasses) and by arc evaporation

into an industrial chamber (ceramic coatings).

Both techniques rely on the same mechanism (see figure 2.2.) allowing the deposition of

materials with potential good electrical conductivity, thermal conductance and high

homogeneity. On one hand, sputtering is referred to a physical process as it is simply a

momentum exchange between charged particles generated under plasma.

The material is therefore removed from a target (i.e., metal, ceramic) due to an energetic ion

bombardment with no chemical reactions. This method enables obtaining a high throughput

of the coating deposition processes, at relatively low running-costs if compared to other

methods. Some important advantages of this method are the ability of putting down the

thin film and the versatility of large areas, easy control of the process and the simplicity of

the equipment. Magnetron sputtering is widely used in the semiconductor photovoltaic,

recording, medical and automotive industries.

On the other hand, during the cathodic arc evaporation (CAE) process a high intensity

current is applied to the cathode and auxiliary electrode, forming an arc that impacts the

target. This shock results in the formation of a few-micrometers large and highly energetic

spot. The interaction between electrons and the atoms create a high degree of cations,

attracted toward the negatively charged substrate with an important kinetic energy on the

basis of their outstanding adherence. The main inconvenient of this technique results from

the high energy associated with the formation of the ion plasma and higher temperatures

than the melting temperature of the material. As a result, some molten metal pits in form of

droplets could be deposited on the substrate surface or on the growing coating,

incorporating itself in the film.

The coating homogeneity is highly altered and structural and properties modifications

should be expected. New methods are already developed for limiting the droplets

implantation during the deposition by applying filters, for instance capable to separate the

plasma flux [11 – 13].



62

Due to the high electric field between the target and the substrate, kinetic of cations is high

leading to a high adherence of the coating. The overall principle for the two deposition

methods are represented in figure 2.3.

Figure 2.3. Schematic representation of the (a) sputtering process and arc

evaporation (b) [14]

The thin film metallic glasses were deposited in a laboratory-size reactor at the IRTES-

LERMPS laboratory, Université de Technologie de Belfort-Montbéliard. The ceramic

coatings were prepared at Ionbond industrial company specialized in such types of thin

films using the optimized industrial conditions. Generally, a PVD system is mainly

composed by a deposition chamber, a vacuum system, a gas flow control system, an

electrical system and a control unit.

In figure 2.4. the magnetron co-sputtering system used for deposition of the Zr-Cu metallic

glass films is presented, together with the main components.

Figure 2.4. PVD magnetron co-sputtering deposition system: 1 –deposition chamber, 2 – rotative

sample holder, 3 – programming and control unit, 4 – inside deposition chamber with different

metal targets, 5 – generators for applying current on targets


Chapter 2

63

Compared with laboratory-designed deposition machines, the industrial systems are

distinguished by their bigger deposition chamber size (volume of about 1.2 m3) and capacity

to deposit parts by different shapes and sizes in the same time. The system disposes of

double pumping system (primary and secondary), heating system and allows applying a

tension (bias) on sample holder and consequently on the samples to be coated. The bias

plays an important role during the deposition step since it accelerates the ions generated by

the target. The bias combined with the deposition temperature is among the most

imperative parameter which could influences the coating structure, adhesion and the final

properties and applications. The triple rotation of the sample holder ensures a minimal

periodicity effect and a higher homogeneity of the coating. The multicomponent targets are

obtained using the powder metallurgy technology, namely, grinding and sintering of pure

metals. In figure 2.5. the arc-PVD deposition equipment used for deposition of CrN-based

coatings is presented, with the main components.

Figure 2.5. Arc PVD industrial equipment [photos Ionbond], 1 –programming and control unit, 2

– deposition chamber, 3 – carousel samples holder, 4 – radiant heating system, 5 – target of

different sizes

2.1.3. Deposition conditions

Substrates

The films were deposited on different substrates depending on the characterization and

property targeted. Table 2.1. presents the different categories of substrate and the main

experimental methods used for analyzing the films. The Si wafers were in general used for

the SEM fractured cross section obtained perpendicularly to film, annealing and

nanoindendation tests due to lower roughness. Free-substrate films were obtained from the

films deposited on glass after emerging in an ultrasonic cleaning machine into ethanol.

It is well known that machining tools are usually made of iron-based materials. The ceramic

coatings have been considered mainly for their protective surfaces for the cutting tools

devoted to industrial applications. The various substrates used during the deposition and

investigations endorsed a base research level on one hand and a more practical/industrial

approach on the other hand.

Considering the as-deposited films it has been confirmed that films structure was

independent on the nature of the substrate (Figure 2.6.).



64

Figure 2.6. Example of the CrN coating deposited on different substrates

Table 2.1. Different natures of substrates and their objective

Substrate type Investigation target / technique

Si wafer SEM, EDS, SEM-cross section, thermal annealing,

nanoindendation

Glass DSC, thickness (profilometer), electrochemical tests, EDS

Steel (HSS M2 Steel,

316L)

oxidation tests (TGA), FIB foils for TEM, ESEM tests,

tribological tests

SiO2 in situ XRD, SEM

Chemical composition of the M2 steel used typically for cutting tools is presented further in

table 2.2.

Table 2.2. Chemical composition of the M2 steel substrate

Element Content (%) Element Content (%)

C 0.78-1.05 W 5.50-6.75

Mn 0.15-0.40 V 1.75-2.20

Si 0.20-0.45 Cu 0.25

Cr 0.20-0.45 P 0.03

Ni 0.3 S 0.03

Mo 4.50-5.50

One of the most important factors influencing adhesion in the deposition process is the

surface preparation. To ensure optimum bond performance and long-term environmental

resistance, substrates must be free of organic (i.e., grease, dirt, oil) and inorganic

contaminants (i.e., rust, scale, oxide layers). These can be cleaned by either mechanical or

chemical processes, or a combination of both.


Chapter 2

65

The procedure for obtaining a proper film could be divided in three principal steps as

follows: - ex situ cleaning of the substrates, in situ cleaning of the substrates and deposition

of the coatings.

The ex situ cleaning of the steel substrates consists of a manual grinding with SiC paper from

400 to 1200 grad. For obtaining the mirror like surface, final polishing with 1-3 μm diamond

paste using a tightly woven cloth was applied. All the other substrates were degreased into

acetone-based solution for 10 minutes using the ultrasonic bath.

The second way to clean substrates (etching) is done in the deposition chambers and was

intended to remove the small contamination layer that has been developed after the primary

cleaning process. This step consists in bombarding the surface with incidents ions

accelerated by a high electric field under high vacuum and argon plasma. In the same time

during this step some micro roughness in the substrates surface are created that will, then,

enhance the coating adhesion by a simple “entrapment”.

Zr-Cu thin film metallic glasses

The metallic glass films were prepared in a static mode (substrates fixed in the front of the

target), using a gas atmosphere composed of argon by means of a 40-liter cylinder Alcatel

450 SCM sputtering reactor. The base pressure was always below 10-4 Pa before refilling

with argon at the desired pressure around ~0.5 Pa.

The study was mainly based on Zr-Cu TFMG in a wide compositional range (13 – 98 at. %

Cu). The target with a diameter of 145 mm and thickness of about 6 mm and high purity

(99.6 at.%) were powered by an Advanced Energy Pinnacle and MDX500 generators.

We deliberately used only one copper target (for two zirconium ones) to correct the easier

sputtering tendency with respect to zirconium. The deposition conditions are summarized

in the table 2.3.

Table 2.3. Summary of the deposition parameters for metallic-like films

Parameter Value

Argon flow 50 sccm

Bias ground

Target current intensity between 0.1 and 1.5 A

Target tension intensity between ~250 and 550 V

Target power intensity between ~0.70 and 450 W

Total pressure during deposition between ~0.4 and 0.5 Pa

Deposition time between ~ 30 to 120 min

Deposition rate 3 µm/h

Rotation speed 3 rpm



66

CrN-based hard coatings

The hard coatings were synthesized by arc-evaporation using pure chromium and

multicomponent (Al0.6Cr0.4, Al0.63Cr0.35Y0.02) targets using a constant evaporation intensity of

80 A into an industrial equipment. The equipment is used for the daily depositions made at

Ionbond company (Chassieu, France) for different parts from automotive, medical and

decorative industries. In the case of these coatings an ionic etching was applied prior to the

deposition.

During these steps the ion energy generated by the cathode under inert gas is further

accelerated in order to eliminate any further oxides formed on the surface. For the targets,

the pre-deposition pulverization allows to clean their surfaces from any contamination. The

deposition conditions are summarized in the table 2.4.

Table 2.4. Summary of the deposition parameters for hard coatings

Parameter Value Comments

Deposition pressure 3.8 Pa regulation with N2

Bias between ~ -30 and -60 V

Target current intensity ~ 100 A

Target tension intensity between ~18 and 22 V

Deposition temperature 470 °C

Deposition time between ~ 70 to 140 min

Deposition rate ~ 1.5 µm/h

Rotation speed 2 rpm spindle speed = 5.5 carousel rotation


Chapter 2

67

2.2. Thin films characterization

2.2.1. Chemical and crystallographic analysis

The primary features that have to be controlled are the chemical composition, crystalline

phase and crystallinity, microstructure, surface and interface structure, and stresses.

The structure of the obtained films depends on the processing parameters like polarization

tension of the substrate, deposition temperature or working gas pressure. There are few

theoretical models that relate the resultant structure of the coatings, and their physical

properties, considering the deposition parameters. Among these, the model proposed by J.

Thornton [15], must be distinguished, since it is widely accepted as being the most

consistent, presenting a general overview of the developed films´ structure prepared by

PVD methods. In this model, there are two main parameters, whose variation induces the

development of a certain structure and morphology (pressure and substrate temperature) as

shown in figure 2.7.

Figure 2.7. Schematic representation of Thornton thin film structural zone diagram (Td

represents the deposition temperature and Tf the melting point of the coating material)

From a qualitative point of view, four zones can be distinguished in this model,

corresponding to four different types of microstructure. The materials grown with features

of these zones are characterized for the predominance of adatoms diffusion, depending on

the Td/Tf ratio.

Despite of this complex diagram, the well-known substrate polarization (bias voltage) is not

considered within this model. This factor becomes extremely important for the

microstructure evolution and thus for all physical/chemical properties of the deposits.

The influence of this parameter in film properties is well documented in several research

works [16, 17].



68

The coatings crystallographic parameters are analyzed by the x-ray diffraction (XRD) non-

destructive technique widely used in the materials science area and which relies on Bragg’s

law, expressed as:

λ n = 2 dhkl sin [Eq. 2.1.]

where λ is the wavelength of the incident beam, n is an integer, dhkl is the interplanar

distance and the incidence angle.

The diffraction profiles are obtained using the x-ray diffraction Brucker D8 Advance System

apparatus equipped with Cu Kα radiation (λ = 0.154 nm). Different configurations can be

employed to find the crystallographic arrangements using a complex methodology as it is

explained in the following:

- Bragg-Brentano,

- grazing incidence,

- in situ heating

The classical Bragg-Brentano (Figure 2.8.a) is not always well adapted for studying the

films structure since interferences with the substrate materials might be expected.

As an alternative for thin films analysis the grazing incidence method (Figure 2.8.b) allows

the diffraction of a stationary incident beam impacting the surface with an angle that can be

as low as 0.5°. The emitter source is fixed and only the detector rotates and parallel beams

are detected. As an example for a CrN coating when such method is applied at 0.5° incident

degree only 0.1654 µm of the film is analyzed. The enormous advantage of this technique is

that no substrate effect is detected on the diffractograms since the beam depths can be

easily controlled as a function of the film’s thickness.

Figure 2.8. Examples of X-ray diffraction setup (a) classical Bragg-Brentano and (b) grazing

incidence [18]

Furthermore, except the information related to the crystalline planes, the grain size and

lattice parameters can be estimated from the peaks broadening. The Debye-Scherrer formula

gives a relation between the peak width and the size of the crystallites, assuming no internal

stresses, expressed as:

𝐷ℎ𝑘𝑙 =0.89 𝜆

𝛽 cos (𝜃) [Eq. 2.2]


Chapter 2

69

With Dhkl the average crystallite size in the <hkl> direction, β the angular size at the half-

height of the diffraction peak, the incidence angle in radians and λ the incident beam

wavelength [19].

The diffraction profiles obtained were indexed using the International Centre for Diffraction

Data (ICDD) database and further fitted by software (TOPAS) for acquiring the lattice

parameters in a quantitative manner.

The structural evolution and thermal stability of the films were followed by means of in situ

heating XRD.

Figure 2.9. presents the X-ray diffraction Brucker system in the high temperature

configuration with all its components.

Figure 2.9. The X-ray diffraction Brucker system under its in situ high temperature configuration:

1- X-ray source, 2 – HTK 1200N heating furnace 3 – detector, 4 – gas bottle (Helium), 5– control

unit, 6 – turbo pump

Principle

For the analysis of the thermal structural evolution, the in situ high temperature XRD

chamber is used from room up to 600 °C and 1200 °C for metallic glass films and for hard

coatings, respectively, with a heating step of 30 °C. The advantage of such analysis

compared to the standard post-annealing testing is that the structural evolution is more

dynamically followed. Prior launching the XRD experiment the oven is heated up to 700 °C

without any sample inside under vacuum at a typical pressure of 2.4 × 10− 2 Pa.

The temperature is maintained here for several minutes and then cooled to room

temperature. This procedure was repeated several times in order to evacuate all the reactive

gasses and especially the oxygen from the heating chamber.

At the last step the helium was smoothly introduced in the chamber and the in situ XRD

experiment is launched.



70

2.2.2. Multiscale microstructural characterization

The microstructure of the thin films is explored through a multiscale approach starting from

the base metallographic level as the porosity rates and morphological view throughout

grain orientation. To begin with, the microstructure of a thin film depends on a variety of

factors, such as mechanism of nucleation and growth, effect of buffer layers, chemical

composition, temperature and so on. Thin films microstructure can be quite different than

the bulk material with respect to grain size and texture.

Porosity evaluation

For the ceramic coatings it is quite essential to know the porosity rates and to understand

their corrosion behavior since these defects could form short-diffusion paths for the

oxidation species. The porosity rate is estimated using the mixed potential theory as

explained in details in reference [20]. The model considers the coating as being nobler than

the bare steel (substrate).

Hence, the porosity rate is evaluated as a mathematical ratio between the corrosion current

density of the coated substrate and the corrosion current density of the bare steel at the

corrosion potential of the coated substrate. In the following a more detailed example will be

given.

Assuming two samples with the corrosion potentials 𝐸𝑐𝑜𝑟𝑟𝑠𝑡𝑒𝑒𝑙 for the bare steel and Ecorr

mixed for

the coating different current densities 𝑖𝑐𝑜𝑟𝑟𝑠𝑡𝑒𝑒𝑙 and 𝑖𝑐𝑜𝑟𝑟

𝑚𝑖𝑥𝑒𝑑 will be obtained as plotted in the

figure 2.10.

Figure 2.10. Schematic comparison of polarization curves of coated steel and uncoated steel in

the case of a nobler coating on steel substrate (a) and example of tested sample (b) [20])

Determination of the porosity rate, is given by: estimated as:

P(%) =icorrmixed

icorrsteel at mixed x 100 [Eq. 2.3.]


Chapter 2

71

The porosity rate (P), then corresponds to the area ratio of steel effectively exposed to the

electrolyte. The setup consists in a homemade measuring cell with three electrodes, namely

a working electrode, a reference electrode and the counter electrode, connected to a

potentiostat as illustrated in figure 2.11. The measurements of the linear polarization in the

range of -0.25 V versus open circuit to 0.9 at a scanning rate of 0.16mV/s and 1mV increment

are performed using the potentiostat/galvanostat Princeton Applied Research model 273.

The tests are done in 30g/L NaCl ang 10 g/L Na2SO4 solutions.

The corrosion current density (icorr), corrosion potential (Ecorr) are determined using the

Tafel’s method.

Figure 2.11. Schematic diagram of the glass 0.5 L cell setup for the electrochemical tests

It is well known that the coatings obtained by CAE technology have usually lots of pores

and a high quantity of droplets inherent to the deposition process. After the initial

evaluation of these characteristics the scanning electron microscope (SEM) was used to

explore and identify the topography and morphology of the thin films. During the course of

this work, a ZEISS Supra55VP SEM apparatus was used to obtain high resolution

micrographs of the samples surface.

The machine can operate in different modes depending on needed information. Secondary

Electron Imaging (SEI) mode is used to produce high resolution images of the sample

surface and showing morphology and relief. The back-scattered electrons (BSE) mode

permits the “identification” of the chemical elements related to the atomic number (Z) of the

specimens. For the heavier elements a lighter image is obtained, whereas lighter elements

appear darker. Using this mode the information related to the distribution of the chemical

elements constituting the sample can be easily obtained. The energy dispersive X-ray (EDX)

can be also performed to determine the chemical species on a spot or an area of the analyzed

sample surface.

In the context of this study, SEM is used to obtain surface micrographs and to identify the

defects and oxides formed on the surface. The transversal cross-sectional morphologies

obtained by perpendicularly fracture of the films allows studying the growth type, in-depth

defects, and thickness of the films.



72

Figure 2.12. displays an example of typical micrographs obtained by SEM for AlCrYN

coating in top surface (a) and Zr-Cu thin film metallic glass in cross section (b).

It is interesting to highlight the droplets and pores seen for the ceramic coating and the fine

columnar-like growth for the Zr-Cu film.

Figure 2.12. Example of SEM micrographs for AlCrYN ceramic coating (a) and cross section of

Zr-Cu thin film metallic glass (b)

Nowadays, the evolution of the materials surface during oxidation for instance, can be

pursued in a real time in situ observation by means of an environmental scanning electron

microscope (ESEM). The particularity of such equipment is that a low pressure of gas could

be reached inside the analysis chamber which allows studying organic samples of even

insulating organic specimens. The system is able to use the reactive to produce the electronic

signal employed for generating micrographs. Of course, a compromise between the working

pressure (some Torr) and the expected resolution has to be generally adjusted.

In our case, the ESEM is employed for monitoring in situ and in real time the evolution of

the oxidation phenomena of the hard coatings (i.e., CrN) at high temperatures.

Figure 2.13. shows the inside ESEM chamber and its components.

Figure 2.13. Image from inside ESEM equipped with the heating chamber: 1 - heating stage, 2 -

thermal protective shield, 3 - column and gaseous SE detector, 4 - cooling circuit


Chapter 2

73

The samples analyzed were deposited on M2 steel with round disk-shape of about 3 mm in

diameter under a water pressure of few Torr from room temperature up to 1000 °C.

The images are taken at several magnifications from 250 up to 8000 at a heating step of 50°C

/min. The experiments have been mainly performed by Annie Malchère from the MATEIS

Laboratory.

As it has been stated at the beginning of section 2.3. the multiscale investigation consists as

well on local investigation of the structure at a deeper level. The microstructure of the

coatings is investigated at a deeper scale by means of the JEOL 2010F transmission electron

microscope (TEM).

The ultra-thin specimens are prepared in two ways with respect to film’s type and targeted

investigation. Specimens for metallic glass films are prepared by ion machining using a

(Gatan PIPS apparatus) using argon ions with a primary polishing for about 50 minutes at 8°

incidence and beam energy of 4 keV. The final polishing takes 5 minutes at a 6 ° incidence

angle and 2.7 keV beam energy. For the hard films, the TEM thin foils are prepared by

focused ion beam (FIB) technique using a Zeiss Nvision 40 dual beam microscope.

Different step of preparation are presented in figure 2.14.

Thin foils’ preparation has been performed by Thierry Douillard from the MATEIS

Laboratory.

Figure 2.14. Schematic representation of a TEM sample prepared by FIB: cutting stage (a),

bonding and extracting steps (b) and the final step as TEM thin foil (c)

The TEM microscope is operating at a 200 kV accelerating voltage .High resolution images

(HRTEM) and Selected Area Electron Diffraction (SAED) can be obtained. Regarding the Zr-

Cu films, SAED allows to identify the crystallographic (or amorphous state) structure of

films. This structure can be then confirmed at a deeper scale by HRTEM. For the CrN

coatings, imaging at different magnifications can be performed to check the general

structure of the grains and the grain boundaries.

The TEM observations have been mainly performed by Claude Esnouf from the MATEIS

Laboratory.



74

2.3. Functional characterization

The investigations presented in the previous sections allow characterizing materials from a

metallurgical approach, as reflected by the coatings deposition step, morphology, nature

and structure of films. The way how the films are deposited plays an important role on the

final functional applications for which the coatings are devoted. This section deals with the

measurement of the coatings’ properties, able to a better understanding their in-service

behavior.

2.3.1. Thermal analysis

2.3.1.1. Thermogravimetric analysis (TGA)

The oxidation resistance of the ceramic coatings is determined using the thermogravimetric

technique using Setaram TGA92 thermobalance as presented in figure 2.15. The coating

deposited on M2 steel is hold by a platinum wire via its drilled hole.

Figure 2.15. Schematic diagram of the Setaram TGA 92 assembly (a) and sample used for testing

(b)

The oxidation resistance is characterized in several modes: isothermal, cyclic and dynamic.

Isothermal oxidation is carry out at temperatures ranging from 750 °C to 1000 °C every 50

°C with a heating ramp of 40 °C/min. The experiment consists in heating at the desired

temperature, keeping that temperature for 4 hours with a final cooling down to room

temperature. The mass variation is recorded during the experiment, but only the variation

during the isothermal period is considered as a function of time.

In dynamic oxidation mode the sample is heated with 50 °C/min up to 500 °C/min and at 1

°C/min from 500 to 1200 °C/min. It must be highlighted that we controlled that oxidation

did not start significantly below 500 °C. All the oxidation experiments are performed in dry

air atmosphere.

Finally, the thermodynamic parameters such as activation energy (Ea), kinetic factors (kp)

and critical temperature of oxidation (Tc) are achieved from the interpretation of the

thermograms.


Chapter 2

75

2.3.1.2. Differential scanning calorimetry (DSC)

The thermal behavior was investigate during a standard differential scanning calorimetry

(DSC) commercial instrument (Perkin Elmer, DSC-7) under high purity dry nitrogen at a

flow rate of 20 mL/min. Prior to the DSC tests , substrate -free films were obtained using

ultra-sonic cleaning into ethanol. Aluminum pans were used as sample holders.

All samples had a similar mass, of about 5 mg. The procedure of DSC experiments is as

following: initially, the samples were heated up to 850 K at a constant heating rate of 20

K/min.

Then, the samples were cooled d own to the room temperature at a cooling rate of 50 K/min.

The same thermal parameters are used to determine the baseline during a second identic

heating program.

2.3.2. Corrosion resistance

The metal parts used during the industrial processes may undergo various extreme

conditions, such as high temperatures and corrosive environments. The oxidation, wear and

corrosion may impact the nature and the lifetime of pieces. Coatings are known to be a

reliable solution for protecting parts against such attacks. The corrosion is a major problem

in the industry and therefore the surface engineering protection has to be better understood

for successfully improving the lifespan of the components.

In order to evaluate the corrosion behavior of the thin films deposited on M2 steel substrate

electrochemical polarization tests in saline solutions are conducted. First of all, the corrosion

potential and corrosion kinetics are of main importance for the finding of the best coating

able to fulfill the desired anticorrosion application.

The electrochemical experiments were carried out at room temperature in a typical three-

electrode glass cell. Graphite probe is used as counter electrode with a surface of 2.5 cm² and

a saturated calomel electrode (SCE) as reference electrode (0 VSCE = 0.248 VENH). All

potential are reported from SCE potential. Only glass coated was used as working electrode

with a surface of 0.28 cm2. The reason of this insulating substrate is to study only the

intrinsic electrochemical behavior of the coating without any galvanic coupling involved

with a conductive metallic substrate.

The electrolyte used in this study was a naturally aerated sulfate sodium solution (Na2SO4,

10g/L) solution with a natural pH of 6.6 prepared from reagent-grade chemicals and

distilled water. The tests were performed with a Biologic SP-300 potentiostat from Biologic®

with ultra-low current cables into a Faraday’s cage. Open circuit potential (OCP) was

measured for a time period of three hours at the beginning of each experiment to ensure the

steady state required for a proper measurement. The polarization curves were recorded

with a scan rate of 0.166 mV/s from 0.5 V below up to 0.9 V above the corrosion potential.



76

2.3.3. Mechanical properties

The values of mechanical properties are achieved by nanoindendation using the continuous

stiffness measurement method with a Berkovich indenter under load of 20 mN. Prior to the

measurement, the contact area function was calibrated using a standard fused silica

material. The values of hardness and Young's modulus were obtained from load/unload

curves at room temperature and consisted of 15-20 measurements for each sample.

One of the advantages of this method is that it is able to characterize the stiffness and to

calculate the hardness and Young's modulus of the thin films directly according to the

Oliver and Pharr's model. The indentations are performed only in the 1/10 of film’s

thickness in order to eliminate the substrate influence. The films roughness is characterized

by contact profilometry using the AltiSurf 500 profilometer or nanoindentation device.

2.3.4. Elastic properties

The elastic properties for Zr-Cu thin films have been measured with the Brillouin light

scattering (BLS) and picosecond ultrasonic (PU) as presented in figure 2.16.

The common required structural information to extract elastic constants from PU and BLS

measurements of sound velocities is the mass density achieved from a fit to X-ray

Reflectivity (XRR) scans using an optical model [21].

Figure 2.16. Schematic representation of the picosecond ultrasonic (a) and Brillouin light

scattering (b)

The PU technique introduced by H. J. Maris in the 1980s [22] derives from the well-known

optical pump and probe scheme. A femtosecond laser pulse (the pump beam) is absorbed at

the sample surface. This absorption generates, by thermal expansion, a longitudinal acoustic

wave that propagates through the system, perpendicularly to the free surface.

Partly reflected and transmitted at all the interfaces of the sample, the acoustic wave in

return is then probed at the free surface. It is detected from the reflectivity change of the

sample surface using a variable time-delayed laser pulse, commonly called the probe beam

(see fig. 2.16.). For a monolithic film onto a substrate, sharp reflectivity changes are

periodically detected, the so-called echoes, separated by a constant time delay (the time of

flight TOF) as in a standard sonar technique. These acoustical features superimposed on the

photo-thermal decaying background. The real and imaginary parts of the relative change in

the surface reflectivity r(t)/r0 induced by such acoustical strain may be measured by

interferometric or reflectivity measurements [22].

Consequently, if the thickness h of the layer is known, PU gives a direct access to the

longitudinal sound velocity.


Chapter 2

77

Additionally, in-plane acoustic waves could be generated in the GHz frequency range, if the

size of the excitation area is small enough [23] and detected using a classical interferometer.

One of the advantages of this technique is the large frequency domain of the generated

acoustic waves, mainly imposed by the optical properties of the materials.

By this way, a frequency range from 100 MHz up to 1 THz is routinely achieved,

overcoming by three orders of magnitude the bandwidth limitations associated to standard

piezoelectric sensors. Ultra-short acoustic waves have been used to investigate acoustic

properties in various systems, such as thin films, isolated dots or phononic crystals [24 - 26].

The Zr-Cu films metallic glasses, a standard picosecond ultrasonic approach is convenient to

extract the longitudinal sound velocity along the normal of the film-plane, VL.

In our experimental set-up, we use a mode-locked Ti: Sapphire laser source, operating at 800

nm with a repetition rate around 79.3 MHz. The pump beam is modulated at 1.8 MHz to

improve the signal-to-noise ratio. The longitudinal velocity is simply defined by VL =

2h/TOF, where TOF is the time delay between two consecutive echoes, measured with an

accuracy better than 1 ps; obviously, the accurate knowledge of the film thickness h is also

crucial.

Another advantage of the PU technique is its ability to measure directly the longitudinal

elastic constant along the growth axis, C33, from the relation 𝐶33 = 𝜌 𝑉𝐿2. Thus, considering

amorphous materials with isotropic elastic properties the result should be the same as for

the C11 = C22 = C33 elastic constants.

BLS has proved to be a very convenient technique for achieving a complete elastic

characterization of thin films and multilayered structures [28]. In a BLS experiment, a

monochromatic light beam is used to probe the thermally excited acoustic waves which are

naturally present in the investigated medium. The power spectrum of these excitations is

mapped out from frequency analysis of the light scattered within a large enough solid angle.

Because of the wave vector conservation rule in the phonon–photon interaction, the

wavelength of the revealed elastic waves is of the same order of magnitude as that of light

and then is much larger than the interatomic distances, so that the material can be described

as a continuous effective-medium.

Here, we used the backscattering interaction geometry (see fig. 1b), so that the modulus

value of the wave vector (Q) of the probed surface acoustic waves travelling parallel to the

film-plane is fixed experimentally to the value Q = 2qi sin (), where qi = 2/L denotes the

incident optical wave vector in air (L is the laser wavelength) and the incidence angle of

the light with respect to surface normal. The acoustical wavelength () is simply defined as

=2/Q and the velocity (V) of a surface acoustic wave is obtained from the frequency

measurements f thanks to the relation V = f * .

For a ratio h/ > 1 (practically 300 nm), the slower surface acoustic waves, namely the

Rayleigh surface wave, will be no longer influenced by the substrate, neither by the film’

thickness. Thus, the Rayleigh wave velocity will measure the property of the film only,

closely related to the transverse velocity, hence to its shear elastic constant G = C44.

The BLS spectra were obtained at room temperature in air, with typical acquisition times of



78

Typically 300 mW of a naturally p-polarized light was focused on the sample surface and

the backscattered light was analyzed by means of a Sandercok-type 3 + 3 pass tandem

Fabry–Perot interferometer.

For nearly opaque layers, such as the present Zr-Cu films with metallic character, the

scattering mechanism is restricted to the scattering of light by the dynamical corrugation of

the free surface by acoustic waves travelling parallel to the film plane [25]. Thus, we can

only observe the surface acoustic waves with a sagittal polarization (first component along

the Q direction and the second is vertical). For films with thicknesses above the acoustic

wavelength (~0.3 µm) and deposited on a substrate with higher acoustic phase velocity

(“slow” film on “fast” substrate), the inelastic scattered surface acoustic waves with a

sagittal polarization are: the Rayleigh wave (R) and the so-called Sezawa guided waves (S1–

Sn) [29] at higher frequencies. The frequency (velocity) and number of these Si surface

modes are dependent on both the thickness and the elastic constants of the film and of the

substrate, whereas the velocity of the Rayleigh surface wave, VR, is mainly dependent on

the transverse velocity VT = (C44/)1/2through the relation VR = VT, where is a slightly

varying function of the whole elastic constants of the film that remains close to 0.94.

ry measured by SEM

micrographs and XRR, respectively, with accuracy better than 2%. For seek of simplicity, all

films were considered as being an isotropic elastic effective medium that can be described

by two independent elastic constants: C11 and C44 = (C11 - C12)/2. In all cases, the C11

elastic constant is first measured by PU and then considered as a fixed known parameter in

the fitting procedure of our BLS spectra, leading to the C44 constant evaluation.

When the BLS signal/noise ratio was favorable, we can observe a singularity at higher

frequency (the scattered intensity by the ripple mechanism is null for the longitudinal

frequency threshold defined by Q//VL) that enables the measurements of VL [25]


Chapter 2

79

Chapter summary

In this chapter the experimental procedures used during the thesis was described. Three

main parts have been presented, dedicated to the synthesis of films, to their metallurgical

characterization and finally to their assessment in terms of functional performance

In the first part a short history and the basic concepts of the under-vacuum deposition

techniques are given with a special focus put on the PVD processes involved in this work.

Hard coatings are deposited using the arc-PVD in an industrial-size chamber, while metallic

glasses result from the deposition by magnetron sputtering using a laboratory reactor.

Coatings are examined at different levels according to their targeted properties. The primary

features of such materials that have to be controlled are the chemical composition (EDS,

GDOES), the nature of crystalline phases (XRD in different complementary configurations)

as well as their microstructure (FIB, SEM, TEM).

The microstructure of the coatings is explored through a multiscale approach starting from

the base metallographic level as the porosity rates to morphological view using advanced

microscopies to assess growth mechanisms of the as-deposited coatings as well as of the

oxide formed at high temperature (in situ ESEM)..

In the last section the functional behavior of the thin films is spotlighted. It is detailed how

films’ properties can be properly measured in relation with their specificity (hard or ductile)

and potential application. The thermal analysis, such as TGA or DSC, has permitted to

investigate the oxidation resistance of hard films and give some quantitative parameters

regarding thermal stability of metallic glass films.

Electrochemical measurements are used to understand the corrosion behavior of coatings

and estimate their protectiveness.

Finally, nano-indentation technique is presented as the suitable tool to characterize

mechanical properties of both types of film.

Elastic properties have been determined by Brillouin light scattering (BLS) and picosecond

ultrasonic (PU) techniques.



80

List of references

[1] P. Drude, Ann. Der Physik 36, p. 532, 1889.

[2] W. R. Grove, “On the Electro-Chemical Polarity of Gases,” Phil. Trans. R. Soc. Lond., vol.

142, pp. 87–101, 1852.

[3] A. W. Wright, A J. Sci. Arts 13, pp. 49–55, 1877.

[4] D. M. Mattox, R. E. Cuthrell, C. R. Peeples, and P. L. Dreike, “Technical note: Design and

performance of a movable post-cathode magnetron sputtering system for making PBFA II

accelerator ion sources,” Surface and Coatings Technology, vol. 33, pp. 425–432, 1987.

[5] P. D. Davidse, J. S. Logan, and F. S. Maddocks, U.S. Patent #3, 755, 123, 1973.

[6] J. Hohenstein, J. Vac. Sci. Technolo., 5 (2) 65, 1968.

[7] D. S. Campbell, The deposition of Thin films By Chemical Methods”, Chapter 5 from

Handbook of Thin Film Technology. L. Maissel and R. Glang, Mc Graw-Hill, 1977.

[8] S. Schiller, O. Heisig, and K. Goedick, Proceedings 7th Int’l Vacuum Congress 1545, Vienna,

Austria, no. 1977.

[9] R. F. Bunshah and D. M. Mattox, Physics Today, 1970.

[10] B. Rointan F., “Handbook of deposition technologies for films and coatings pag 35 chapter

1,” Novlel publications, 1994.

[11] R. L. Boxman, V. Zhitomirsky, B. Alterkop, E. Gidalevich, I. Beilis, M. Keidar, and S.

Goldsmith, “Recent progress in filtered vacuum arc deposition,” Surface and Coatings

Technology, vol. 86–87, Part 1, pp. 243–253, 1996.

[12] D. A. Karpov, “Cathodic arc sources and macroparticle filtering,” Surface and Coatings

Technology, vol. 96, no. 1, pp. 22–33, 1997.

[13] F. Sanchette, C. Ducros, T. Schmitt, P. Steyer, and A. Billard, “Nanostructured hard

coatings deposited by cathodic arc deposition: From concepts to applications,” Surface and

Coatings Technology, vol. 205, no. 23–24, pp. 5444–54532011.

[14] S. Audisio, “Revêtements et traitements de surface, fonctionnalités, durabilité, procédés”,

Chapter 14, p. 335.”

[15] T. J. A., “Am. Rev. Mater. Sci. 7 (1977) 239.”

[16] M. Ohring, “The materials science of thin films,” Academic Press, San Diego, California 92101-

4495, 1992.

[17] B. C .R. and et al., “Encyclopedia of Materials Characterization: Surfaces, Interfaces, Thin

Films”,” 1992.

[18] “Brucker manual tranning.” .

[19] L.I. Langford and A.J.C. Wilson, “Scherrer after six years: a survey and some new results in

the determination of crystallites size,” 1978.

[20] J. Creus, H. Mazille, and H. Idrissi, “Porosity evaluation of protective coatings onto steel,

through electrochemical techniques,” Surface and Coatings Technology, vol. 130, no. 2–3, pp.

224–232, 2000.

[21] L. Parratt, “Surface Studies of Solids by Total Reflection of X-Rays,” Phys. Rev., vol. 95, no.

2, pp. 359–369, j1954.

[22] C. Thomsen, J. Strait, Z. Vardeny, H. Maris, J. Tauc, and J. Hauser, “Coherent Phonon

Generation and Detection by Picosecond Light Pulses,” Physical Review Letters, vol. 53, no.

10, pp. 989–992, 1984.

[23] Y. Sugawara, O. B. Wright, O. Matsuda, and V. E. Gusev, “Spatiotemporal mapping of

surface acoustic waves in isotropic and anisotropic materials,” Ultrasonics, vol. 40, no. 1–8,

pp. 55–59, 2002.

[24] G. Abadias, P. Djemia, and L. Belliard, “Alloying effects on the structure and elastic

properties of hard coatings based on ternary transition metal (M = Ti, Zr or Ta) nitrides,”

Surface and Coatings Technology, vol. 257, pp. 129–137, 2014.


Chapter 2

81

[25] P. Djemia, F. Ganot, C. Dugautier, and M. Quilichini, “Brillouin scattering from the

icosahedral quasicrystal Al70.4Mn8.4Pd21.2,” Solid State Communications, vol. 106, no. 7, pp.

459–461, 1998.

[26] P. Djemia, M. Benhamida, K. Bouamama, L. Belliard, D. Faurie, and G. Abadias, “Structural

and elastic properties of ternary metal nitrides TixTa1 − xN alloys: First-principles

calculations versus experiments,” Surface and Coatings Technology, 215, pp.199–208, 2013.

[27] M. A. van Dijk, M. Lippitz, and M. Orrit, “Detection of acoustic oscillations of single gold

nanospheres by time-resolved interferometry,” Phys. Rev. Lett., vol. 95, no. 26, p. 267406,

2005.

[28] G. K. Horton and A. A. Maradudin, Dynamical properties of solids. Amsterdam; Oxford:

Elsevier., 1995.

[29] P. Djemia, C. Dugautier, T. Chauveau, E. Dogheche, M. I. D. Barros, and L. Vandenbulcke,

“Mechanical properties of diamond films: A comparative study of polycrystalline and

smooth fine-grained diamonds by Brillouin light scattering,” Journal of Applied Physics, vol.

90, no. 8, pp. 3771–3779, 2001.


Content

3.1. Morphological characterization of the CrN-based coatings ........................................................ 82 3.2. Chemical characterization................................................................................................................. 87 3.3. Mechanical characterization ............................................................................................................. 90 3.4. Structural characterization of as-deposited ceramic thin films ................................................... 91 3.5. Microstructural and crystallographic investigation: solid solution state ................................... 95 3.6. High temperature coatings durability: oxidation resistance ........................................................ 98

3.6.1. Real time and in situ monitoring .............................................................................................................. 98 3.6.2. Real time oxidation observation by Environmental-SEM ................................................................... 101 3.6.3. Kinetic approach of the oxidation .......................................................................................................... 104 3.6.4. Qualitative approach of oxidation behavior ......................................................................................... 111 3.6.5. Microstructural investigation of the CrAlYN coating oxidized at 950 °C ........................................ 119

Chapter summary ................................................................................................................................... 125 List of references ..................................................................................................................................... 126

Chapter 3

Influence of Al and/or Y incorporation to CrN

coatings on their high temperature oxidation resistance

Overview

In this chapter the ceramic coatings produced by arc-PVD method will be discussed with

emphasis on their oxidation resistance at high temperatures. These types of coatings are

mainly used for production of the machining tools during severe utilization conditions.

The microstructural, chemical, structural and thermal modifications which undergo under

such extreme conditions are analysed and discussed in the light of a multiscale approach.


Influence of Al and/or Y incorporation to CrN coatings on their high temperature oxidation resistance

82

3.1. Morphological characterization of the CrN-based coatings

The morphology of the deposited coatings is an important characteristic to be known, since

it provides direct information on the type of film’s growth. From the morphological

characterization made by SEM, some representative plan view and cross-section

micrographs have been chosen to be presented in figure 3.1.

SEM plan view SEM cross-section

Figure 3.1. Representative SEM plan-view micrographs for CrN (a, a1), CrAlN (b, b1) and

CrAlYN (c, c1)

From these micrographs, it seems quite clear, that all three coatings have a heterogeneous

surface with many growth defects of metallic macroparticles (i.e., droplets) and

open/superficial holes. The droplets number appears to be higher for the Al-containing

films, consequence of the different melting points for Al (660 °C) and Cr (1856 °C).

These ejected particles of about few microns in size are emitted from the cathode’s surface

due to either explosion of the matter, or to the high plasma pressures within the cathode

spot on the locally molten material [1, 2].


Chapter 3

83

The defects density and spreading are usually linked to the adopted process parameters

during the deposition but are unlikely to affect much the applications involving high

mechanical loads. The droplets, sometimes also called spits (inhomogeneities) are ejected

from the molten source and tend to point in the direction of the coating flux plume [2 - 4].

The films have a thickness of about 2 µm with a dense and compact growth through the

entire substrate’s thickness. Besides, adherence of the film seems very high, as reflected by a

sharp substrate/coating interface. The coatings’ cross-section displays the very common

PVD columnar-type morphology as already described in numerous studies [5 - 9].

According to the Thornton diagram (see figure 2.7. from chapter 2) the developed

morphologies can be indexed to the zone I [10].

The different melting points of Cr and Al elements induce formation of more depositions

defects of the Al-containg films (Figure 3.1.) and likely different open porosity rates.

This parameter, of great importance for physico-chemical reactivity, are extracted from

potentiodynamic curves (Figure 3.2.) according to the calculation procedure detailed in

chapter 2 [11].

Coating Porosity

ratio(%)

CrN 0.05

CrAlN 0.17

CrAlYN 0.70

Figure 3.2. Potentiodynamic curves of CrN-based and bare steel. On the right side are presented

the open porosity rates

It is essential to keep in mind that such parameter corresponds to the substrate area exposed

to the electrolyte, and therefore only accounts for through-coating porosities, susceptible to

influence mainly the chemical behavior. Values indicate a significant difference between the

films’ porosity (table insert in figure 3.2.) in relation to the defect density observed on the

surface by SEM in figure 3.1.

The porosity measurements showed that the Y-containing film presents the highest

defective surface. The EDS chemical maps of the elements could give more insights on the

porosities distribution as presented in figure 3.3.

-800 -600 -400 -200 0 200 400 60010

-9

10-8

10-7

10-6

10-5

10-4

10-3

10-2

10-1

100

Potential (V/SCE)

Cu

rre

nt

de

ns

ity

j (

A/c

m2)

bare steel

CrAlN

CrAlYN

CrN



84

Figure 3.3. Zone of the EDS chemical analysis of CrAlYN coating deposited on M2 steel in as-

deposited state. The selected zones for comparing the Fe contents are highlighted with red circles

and numbered

The chemical maps show, at the EDS scale, a uniform distribution of the films elements,

with droplets richer in Al (positions 4, 5, 6) and zones with different Fe enrichments in

relation to the local film’s thickness (positions 1 , 2, 3 containing respectively 57.8, 3.7 and

1.3 at.% Fe). Similar chemical inhomogeneities were also reported by Petrogalli et al. for CrN

deposited by CAE technique [12].

Munz et al. were on the basis of a model explaining these different populations of holes

(Figure 3.4.), related the droplets formation and expulsion for PVD (TiN, TiAlN) coatings

deposited by arc evaporation on M2 high speed steel.

Figure 3.4. Schematic representation of the self-expulsion mechanism of droplets for PVD

coatings obtained by arc evaporation [13].

A similar model was also proposed by Mege-Revil et al., able to explain, when oxidized, the

preferential ejection of certain droplets in function of their size and position into the coating

[14]. The model is detailed in the next figure 3.5.


Chapter 3

85

Figure 3.5. Model explaining the preferential oxidation near the deposition defects for a arc

TiSiN coating [14]

The different melting points of the cathode target material elements affect the number and

size of droplets incorporated into the growing coating as reported by Creasey et al. for

instance for TiAlN [15]. Thus the droplets are generated under an eruptive effect during the

cathodic arc-generated metal ion etching process as exemplified in figure 3.4. a.

The solidification process takes place directly on the substrate surface and some of the Al

and/or Cr melt species (liquid phase) are splashed towards the substrate as larger or smaller

droplets. The weak droplet-film bonding strength and/or the high compressive stress levels

(Figure 3.4. b.) of the growing coating should result in a reduced adhesion.

Some droplets could be just simply ejected from the growing film [16]. Panjan and

colleagues studied the growth defects in PVD coatings in numerous works [2, 4, 17, 18].

They highlighted also that the droplets generated during the etching step could serve as

seeds for the growth of nodular-like defects which perhaps occupy large zones from the

substrate to the top of the coating.

Besides, S. Harsha proposed that a coalescence effect during the initial stages of films

nucleation could lead as well to formation of defects with different surface energies [3].

Shortly after the film growth is initiated, the material is pushed away and disk-like growth

defects (empty holes) are generated (Figure 3.5.c).

Finally, the resulting empty holes can be refilled by continuous growing film. Nevertheless,

some droplets could leave the substrate and the coating in a late stage of film growth or

remained always connected to the surface/inside the coating [13].

These different steps of surface’s defects formation have also been observed in our case by

FIB cutting across negative (porosity) as well as positive (droplet) reliefs.



86

Figure 3.6. illustrates the coating before (a, d) and after FIB cut (b, c).

Figure 3.6. SEM micrographs from the FIB experiments showing the general topography (a, d),

and examples of holes and droplet (b, c) for CrAlYN coating deposited on M2 steel in as-

deposited state. The green lines indicate zones where the cuts were performed

From the figure showed above confirms perfectly the “droplet’s ejection model” proposed

by Munz. The two types of holes, emerging or not on the surface are also pointed out.

Regarding droplets, we can see their wide range of size: small droplets may only slightly

affect the film’s growth, while micrometer-sized droplets are badly inserted into the coating

forming large channels prone to favors oxidants’ penetration.

Based on the Munz’ droplet formation mechanism (Figure 3.4.) and taking into account

experimental results of figures 3.6., we consider that in some areas the droplets are ejected,

process that leave to empty or partially filled cavities.

Hence during the growth of CrAlYN several mechanisms that promote deposition defects

take place:

the craters or holes are found to be variable in depth depending on when the droplet

was ejected,

high amount of droplets with different sizes and nature.


Chapter 3

87

3.2. Chemical characterization

In order to determine the chemical composition of the prepared coatings, a study by energy

dispersive spectroscopy (EDS) studies was carried out. We will focus our investigations on

the chemistry of as-deposited films and of some deposition defects (e.g., droplet, holes).

The important fact that should be mentioned here is that results from EDS measurements

give rather a qualitative information, since no standards were applied for experiments.

For each sample a set of at least 6 measurements on different zones were performed to

ensure a proper reliability. In figure 3.7. an example of the typical EDS chemical analysis for

CrAlN coating is illustrated for different surface elements (general, droplet, hole).

Figure 3.7. EDS chemical profiles for the CrAlN coating showing the general composition and

specific details (droplet and hole). The determined elemental concentrations by EDS are

summarized in table 3.1. (Analyzed area comprised the whole image)

Table 3.1. Chemical composition of the deposited films obtained by EDS

Coating Chemical elements (at. %) Al/

(Al+Cr)

Al/(Al+Cr)

target Cr Al N Fe Y

CrN 73.20

±4.56 -

26.30

±4.85

0.50

±0.49 - - -

CrAlN 26.32

±0.41

37.24

±59

35.80

±0.62

0.64

±0.25 - 0.58 0.64

CrAlYN 26.42

±0.64

36.12

±0.52

35.50

±0.94

0.80

±0.31

1.16

±0.12 0.58 0.64

The CrN deposited by arc PVD exhibits an over-stoichiometric nature, while the

composition of Al-containg films is slightly more balanced regarding the Cr/N ratio.

Such a gap could nevertheless be an artifact due to the poor accuracy of EDS spectroscopy to

measure light elements due to mainly of a preferential absorption phenomena of x-ray low

energy. Ratios between metallic species are more accurate.



88

The Fe signal detected is from the substrate and most likely from the open porosities. The

ratio of metallic elements of about 0.58 is close to the one of the multicomponent target

(0.64). Although both ternary Cr-Al-N and Cr-Al-Y-N are richer in Al (than in Cr), we have

decided to called them CrAlN and CrAlYN for an easier comparison with the CrN coating,

used as the reference film.

For a more quantitative approach of the chemical composition, the Glow Discharge Optical

Emission Spectroscopy (GDOES) was employed. Moreover, this technique affords key-

information on the distribution of elements throughout the coating’s thickness.

The GDOES experiments have been performed by Grégory Marcos from Institut Jean

Lamour, Nancy.

In figures 3.8., 3.9. and 3.10. are illustrated the chemical profiles for all three coatings.

Figure 3.8. GDOES profile of as-deposited CrN coating on M2 steel substrate

Figure 3.9. GDOES profile of as-deposited CrAlN coating on M2 steel substrate

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

0

10

20

30

40

50

60

70

80

90

100

Co

mp

os

itio

n (

at.

%)

depth (m)

N

Cr

Fe

M2 Substrate

CrN

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

0

10

20

30

40

50

60

70

80

90

100

CrAlN

depth (m)

Co

mp

osit

ion

(at.

%)

NFe

Al

Cr

M2 Substrate


Chapter 3

89

Figure 3.10. GDOES profile of as-deposited CrAlYN coating on M2 steel substrate

For all the samples, the chemical profiles are homogeneous along the films’ thickness.

The interface corresponds to the zone where iron of the substrate. These results confirm the

film thicknesses observed by SEM cross-sectional imaging. All elements are accurately

measured thanks to the use of suitable standards, except for Y.

Therefore, yttrium content estimated at around 5 at.% is greatly overestimated. However,

what is important to point out is also its uniform distribution. Calculated by difference with

all other elements, a 1 at.% content is deduced, value consistent with the initial doping of the

ternary target. Using the ratios of the other elements, its real chemical content is found at 1

at. % along the whole film’s thickness, as previously counted by EDS. The average relative

atomic concentration of the chemical elements as well as the ratio Al/(Cr + Al) in the

coatings are listed in the Table 3.2.

Table 3.2. Chemical composition of the deposited films obtained by GDOES

Coating Chemical elements (at. %) Al/

(Al+Cr)

Al/(Al+Cr)

target Cr Al N Fe Y

CrN 47 - 50 - - -

CrAlN 22 33 44 - - 0.62 0.64

CrAlYN 21 35 40 - 1 0.62 0.64

The absolute elemental concentration (determined by GDOES) is close to the stoichiometry

of the deposition targets composition and is well corroborated with the EDS measurements.

This general characterization provided some initial results with respect to deposition

condition about the film’s chemical and morphologies. Since we are dealing with CrN-like

hard coatings we will present next the mechanical properties and further the oxidation

behavior at high temperatures.

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

0

10

20

30

40

50

60

70

80

90

100

depth (m)

Co

mp

osit

ion

(at.

%)

N

Y

Fe

Cr

Al

CrAlYN

M2 Substrate



90

3.3. Mechanical characterization

Mechanical properties of ceramic coatings are of interest as these properties strongly

influence friction and wear performance of interfaces. The commercial available tests, such

as nanoindentation, can be used to obtain the hardness and Young’s Modulus of elasticity.

In operation conditions, the wear performance of coated steel tools is primarily controlled

by the chemical and physical properties, as well as by the mechanical characteristics.

Recently, it was shown that Al additions into binary chromium nitride coatings improved

the tribo-mecanical properties [19 - 25] and incorporation of small amounts of yttrium retard

the diffusion processes resulting in higher thermal stability [26 - 28].

In this section we will present the mechanical and tribological behaviors of CrN and Al-

containing films. For the ceramic coatings obtained by arc-PVD method the mechanical

properties studied by nanoindentation are sensitive to surface state due to high amounts of

deposition defects and/or substrate effect [3, 5].

In order to ensure the reliably of the results at least 20 indentations were done with respect

to the 10 % rule of film’ thickness (Figure 3.11.).

Figure 3.11. Example of typical hardness (left) and Young’s Modulus (right) curves obtained

during nanoindentation (CrN coating). The average values were calculated from the highlighted

region (100 – 200 nm)

The nanoindendation curves show an important scattering in relation to the heterogeneous

and rough nature of the indented surface. Therefore, averaged values for hardness and

elastic modulus will rather be considered, after interpretation of graphs at indentation depth

between 100 and 200 nm. These results are summarized in table 3.13. together with other

parameters such as elastic strain to failure factor (H/E) and plastic deformation resistance

factor (H3/E2).

Table 3.3. Mechanical characteristics of CrN and Al-containing films in as-deposited state

Coating Hardness

(GPa)

Young’s Modulus

(GPa) H/E ratio H3/E2

CrN 27 ± 2 396 ±34 0.06 0.12

CrAlN 26 ± 5 514 ± 77 0.05 0.07

CrAlYN 22 ± 7 423 ±97 0.05 0.06


Chapter 3

91

The hardness results were measured to be in the range of 22 GPa for CrAlYN to 27 GPa for

CrN, while elastic modulus from 396 GPa for CrN and 514 GPa for CrAlN coating.

The Al and/or Y incorporation do not enhance the hardness as it could be expected. The

values are generally in the same range of those reported in literature for CrN [29 - 32] and

CrAlN [6, 19, 20 – 25, 33 – 38].

For CrAlYN few studies are presenting results on mechanical and tribological properties.

Yet compared with the available studies, our values resembles to ones already reported [25,

26, 28]. On the other hand, with the hardness and elastic modulus values some other

important parameters can be calculated, such as elastic strain to failure factor or plastic

deformation resistance factor. The higher values of (H/E) and (H3/E2) for the CrN film, could

imply an improved tribological properties. For Al-containing films these parameters

indicate similar tribological properties as CrN.

The mechanical properties stability is of a main importance when the ceramic materials are

used during industrial applications. Their structure contributes to this stability as will be

discussed in the following section.

3.4. Structural characterization of as-deposited ceramic thin films

The PVD deposition techniques, such as sputtering or arc evaporation, can contribute to the

formation of metastable structures with high compressive residual stress in the coatings

which could improve adherence and performance in operating conditions [39]. These

coatings are sensitive to structural changes, requiring a high potential energy to change to

an equilibrium state, so as they not breakdown during mechanical applications. The nitride

coatings (e.g. TiAlN, TiN) deposited by PVD have, usually, preferential growth orientations

and sometimes lattice parameter deviations due the existence of impurities into the

structure.

Hence a detailed structural characterization was further carried out with the aim to identify

the sample’s crystalline phases and to explore their structural parameters (including lattice

parameter or crystal domains).

In figure 3.12. is presented the X-ray diffraction patterns of as-deposited coatings deposited

on M2 steel substrate.

Figure 3.12. XRD patterns of as-deposited films on M2 steel substrate. The green color highlights

the AlN phase, while the CrN phase is presented with a red color with the theoretical positions

from ICDD database

35 40 45 50 55 60 65 70 75 80 85

CrAlYN

CrAlN

fcc- AlN (a = 4.04 Å)

fcc- CrN (a = 4.14 Å)

Substrate (Fe, Mo, C)

{20

0}

of

Fe

{21

1}

of

Fe

{22

2}

{31

1}

{22

0}

{11

0}

of

Fe

{20

0}

Inte

ns

ity

/a

rb.

un

its

2deg

CrN

{11

1}



92

Owing to the low thickness of films, we can also see the appearance of several intense peaks

belonging to the M2 substrate (e.g., 2θ = 45°, 65°, 82°). The existence of the substrate peaks

on the XRD diffractograms impedes a proper indexation of the crystalline phases, especially

for the 2 ~ 45 °where several peaks overlap. The initial interpretation of the XRD patterns

indicates a structure close to CrN and AlN phases, but without any additional details due to

the substrate’s peaks superimposition.

The structure is among the first important feature of the ceramic coatings that has to be

known before performing any other mechanical or property investigations.

For a better diffractograms’ interpretation, films were deposited on an amorphous silica

substrate and their XRD patterns are shown in figure 3.13. below.

Figure 3.13. XRD patterns of the CrN, CrAlN and CrAlYN deposited on SiO2 substrate in the as-

deposited state. The figure shows the theoretical positions of the AlN and CrN phases from the

ICDD database

Without any harmful influence of substrate, peaks’ discrimination becomes easier. The films

exhibit several intense peaks in the {111}, {200} and {220} growth directions.

The aluminum and/or yttrium incorporation make peaks broaden, implying a structural

refinement of the lattice and smaller crystal domains [7]. Regarding the latter films, peaks

are localized between the CrN and AlN unstrained phases, suggesting a possible complete

solid solution of chromium and/or aluminum in the f.c.c. lattice type [40].

In the International Centre for Diffraction Data (ICDD) database several AlN phases can be

found as presented in table 3.4. Hence it is challenging to use the right phase to proper

identify the structure without any additional information. Equilibrium phase of AlN is here

excluded because peaks positions are not adequate.

35 40 45 50 55 60 65

inte

ns

ity

/a

rb.

un

its

CrAlYN

CrAlN

CrN

2 /degree

{11

1}

{20

0}

{22

0}

fcc- CrN (a = 4.14 Å)

fcc- AlN (a = 4.04 Å)

fcc- AlN (a = 4.12 Å)


Chapter 3

93

Table 3.4. The f.c.c. potential phases with similar lattice parameters available in the ICDD

referenced database

Phase Strongest intensity Lattice

a /Å

CrN (76-2494) {200} 4.14

AlN (87-1053) {111} 4.31

AlN (25-1495) {420} 4.12

AlN (46-1200) {200} 4.04

For the CrN coating, the peaks have the same intensities as the unstrained CrN structure

with a preferential orientation along the {200} plan, while for the Al-containing films, the

crystal domains are strongly orientated along the {111} plan, suggesting a columnar effect as

noticed in refs. [41, 42]. Depending on the selected deposition parameters, the preferential

growth of arc-evaporated CrN can have different orientations, microstructure and phase

compositions (e.g., Cr2N, Cr, CrN) [43].

The common preferential orientation for CrN deposited by arc evaporation or other PVD

methods is often reported to be {200} [33, 44 - 48], although {111} are rarely identified [5].

In some cases the coatings revealed textures diverting from the classical {111} or {200} to

{211} or {122} preferred orientations [48]. For the CrN deposited in this work the main

crystallographic reflections are well indexed with the f.c.c.-CrN phase.

On the other hand, the CrN and AlN phases belong to the same space group (Fm3m) and

have similar lattice parameters as described in table 3.4., which makes difficult an accurate

indexation of the XRD patterns. Based on the XRD peaks from figure 4.13. it is possible to

attain other structural parameters that could help determining the structure of the Al-

containing films. For this approach, all the diffracted peaks have to be considered to achieve

a correct fitting of the experimental results. The evolution of the lattice parameters for the

coatings investigated are displayed in the figure 3.14.

Figure 3.14. Development of the lattice parameters for the as-deposited CrN, CrAlN and CrAlYN

coatings. The figure shows the theoretical positions of the f.c.c.-CrN (ICDD card no. 76-2494) and

the two f.c.c.-AlN phases, which lattice parameters are 4.12 or 4.04 Å (AlN1-ICDD card no. 25-

1495 and f.c.c.-AlN2 card no. 76-2494, the third variety of a=4.31 Å being not realistic)



94

As it is shown in the figure, the structural investigation of CrN film yielded a lattice

parameter close to 4.14 Å. The obtained value and the intensity ratio of the detected

reflections are identical to the stress-free CrN structure from ICDD database [47].

Incorporation of aluminum and yttrium induces distortions in the f.c.c. lattice.

The lattice parameters for the Al-containing are around 4.09 Å, which appears to be a value

between the one of AlN2 (4.04 Å) and CrN (4.14 Å) structures. The value of lattice

parameters (around 4.09 Å) recorded for Al-containing made us to believe that the films’

structure is a solid solution type of Cr(Al)N type with complex structural arrangements

between the AlN1 and CrN phases. Moreover, the AlN (4.04 Å) structure corresponds to an

unstable phase obtained under high pressure (temperature > 1400 °C and pressure > 16.5

GPa). Thus, this metastable structure can be formed due to the compressive stress level

during the arc deposition which produces higher lattice strains.

The lattice parameter is slightly different for CrAlYN film, which could be attributed

slightly different higher lattice strains due to incorporation of yttrium larger atoms.

The (Cr,Al)N intermediate nature of our coating confirmed other previous results of the

literature. Many reports have stated that when the Al content is below ~ 77 at.% a solid

solution with lattice distortions is generally formed [19, 31, 41 – 43, 49 - 53].

Hence it appears that the structure of the Al-containing coatings is in a metastable form,

where Al atoms could substitute the Cr atoms in the f.c.c.-CrN lattice as suggested by refs.

[49, 51, 52]. The Cr atoms could equally substitute the Al atoms in the f.c.c.-AlN lattice as

described by Reiter and coworkers [53].

For the Y-containing film it has been found that yttrium could replace as well the Al and Cr

atoms in the CrAlN lattice [8, 54]. Several different authors suggested that AlN forms solid

solutions with chromium when the deposition process is far from the equilibrium

conditions [51 - 53]. Based on the mixing enthalpies for AlYN and CrYN structures reported

by F. Robere et al. [54] the yttrium content must be limited in order to obtain single phase

films [9].

Moreover, the structural differences are directly observed on the evolution of the crystal

domains size as derived from the Scherrer’s formula [55] shown in figure 3.15.

Figure 3.15. Crystal domains’ size for the CrN, CrAlN and CrAlYN coatings in as-deposited state


Chapter 3

95

From the figure it can be seen that CrN coating exhibits crystal domain size in the range of

25 – 45 nm. The columnar structure becomes finer once Al and/or Y is added in the cubic

matrix. Moreover, the different values of crystal domains size are explained by a texture

effect. For instance, <111> texture has a consequence to give higher extracted from {111}

diffraction peak.

The Al and/or Y addition refine the structure as the average crystallize size decreased from

34 nm for CrN to 24 nm for CrAlN and 22 nm for Y-containing film. These values for CrN

are slightly lower compared with the ones reported for instance by Ernst et al. [5] for a

corresponding arced CrN (~ 40 nm) but is close to the sputtered CrN (maximum average 33

nm) reported by Lee et al. [56].

For the other two coatings, the crystal domains size are higher as compared with sputtered

CrAlN (maximum average 11 nm) reported by Chunyah and colleagues [57], but close (17 -

18 nm) to the ones found by others [38, 40] or Qi et al. for Y-containg film (27 nm for 1.5 at.%

Y content) [9].

The grain size tailoring and distribution is crucial for improving the orientated-properties,

such as mechanical or tribological properties of the ceramic coatings [19 – 25, 33].

Results show that the structural transformations would occur at higher temperatures when

Al and/or Y is incorporated in the fcc lattice, mainly due to a structural refinement [41, 42,

49, 53, 58, 59].

The XRD results from this section presented the main structural differences between the

coatings in their as-deposited state. For a detailed investigation, TEM investigations were

employed to directly observe the microstructure at a local scale.

3.5. Microstructural and crystallographic investigation: solid solution state

The objective of this part is to characterize the metastable structure resulting from the

association of Al and/or Y with CrN, previously established by XRD results. TEM was

mainly performed for ternary and quaternary coatings.

For this study, TEM thin foils were prepared by focus ion beam (FIB) using the cutting

procedure described in the experimental methods’ chapter.

Figure 3.16. depicts some representative examples of the developed film microstructure

obtained in High Angle Annular Dark Field (HAADF) and Bright Field (BF) modes.

First mode gives rather information on the chemistry of film, while BF in more adapted for

microstructural aspects.

Figure 3.16. FIB cross-sectional TEM micrographs showing a general view of the CrAlN coating

obtained in the HAADF mode (a) and the one of CrAlYN coating obtained in bright field (b)



96

From this figure it is clearly noticed the typical columnar microstructure along the growth

direction, as remarked as well from the SEM cross-sections (figure 3.1).

The samples display also incorporation of small droplets inside the film microstructure.

Some detailed zones showing the columnar-like microstructure are given in figure 3.17.

From the figures the well orientated columnar structure resembles of the one reported for

CrN [48] or CrAlN [60, 61] coatings. Here is noteworthy to mention some aspects related to

structural imperfections. At a nanometer scale, the crystalline solids can enclose large

numbers of various defects or imperfections [16].

This structural state favors the formation of lattice defects during deposition, considering

that Al and Cr atoms can substitute one each other at all proportions in the fcc lattice [2, 13,

17]. TEM observations indicate a complex columnar microstructure with imperfections that

are related to the arc-deposition process (inside and superficial droplets, continuity defects,

nanostratification features…), as exemplified in figure 3.17.

Figure 3.17. Bright and dark fields TEM micrographs detailing the inside deposition defects for

CrAlYN coatings deposited on M2 steel substrate

It is important to underline the very small dimension of semi-spherical droplets.

Such particular shape can be explained by the partially melt nature of droplet, arriving on

the growing surface and then embedded into the coating. It can be observed the “comet-

like” architecture induced by the droplet insertion, leading to imperfect embedding and to

potentially harmful defects for corrosion/oxidation aspects.

Furthermore, to better understand the microstructure a selected area was defined, where

SAED has been performed, as shown in figure 3.18.

Figure 3.18. A zone where SAED has been performed for the film presented in figure 3.16.b. The

zone selected represents the whole CrAlYN film. N is the normal to the film surface


Chapter 3

97

From SAED the diffraction rings can be assigned to {111}, {200}, {220}, {311}, {222}, {400},

{331}, {420} and {422} planes of a cubic Cr(Al)N phase. Due to high intensity of {111} spot

along the normal to the film surface, it can be observed a certain preferential orientation of

the crystals along the {111} plan growing direction, confirming XRD results from figures

3.13.

A general EDS analysis achieved with a probe size of 0.30 µm for CrAlN and 0.75 µm

diameter for CrAlYN coating have been taken in the middle of the films structure (Figure

3.19.) endorses a composition close to the deposition target’s one as highlighted from table

3.2.

Figure 3.19. Selected zone for EDS chemical analysis for the as-deposited CrAlYN (a) and CrAlN

coating (b)

Table 3.5. Local EDS analysis performed on the FIB cross sectional TEM for Al-containing

coatings in the as-deposited state

Coating Chemical elements (at.%)

N Al Cr Fe Y

CrAlN 47 32.3 20.5 0.2 -

CrAlYN 47 32 19 0.6 0.9

Based on the values from table, the coating chemical stoichiometry is Al32.3Cr20.5Y0.9Fe0.2N47

and Al32Cr19Y0.9Fe0.6Y0.9N47 with a Al/(Al+Cr) ratio of 0.62 which is quite close to the

deposition target ratio 0.64. The chemical analysis performed as this local microstructural

state resembles to the ones found by GDOES at a coarser level (tables 3.2).

Finally, the findings from the TEM investigations confirmed the columnar growth of a

Cr(Al)N solid solution structure with a preferential orientation along {111} as previously

noted as well from the XRD results (Figure 3.13.).

The compositions of the as-grown Al-containing films are similar to the deposition target,

denoting the formation of a dense and chemical quite homogeneous coating in the thickness.

In the following we will present the oxidation resistance of such coatings during exposure at

high temperatures.



98

3.6. High temperature coatings durability: oxidation resistance

The investigation of high temperature (HT) oxidation resistance takes many forms and

usually the main interests are focused on oxidation process, kinetics and involved

mechanisms. In the case of CrN-like ceramic coatings from this thesis, the emphasis is put

on kinetics of degradation with consequences on microstructural changes due to heating.

Such coatings are generally used in engineering applications to protect the tools’ surface,

delay the microstructural evolution and extend their lifetime. Therefore, this section will

deal with the rate measurements, an analysis of the reaction-products morphologies and an

estimation of structural modifications. The degradation of CrN-like films during oxidation

at HT has been explored adopting a multiscale approach based on a quantitative,

qualitative, and in situ characterization at different levels from SEM to TEM.

3.6.1. Real time and in situ monitoring

Structural transformations during continuous heating (HT- XRD)

In the previous sections related to the structure of as-deposited films, it was shown the

general differences between the coatings’ structure. The primary application of the ceramic

coatings is to protect the tools’ surface (e.g. cutting) against its oxidation during utilization

in industrial processes. In order to determine the phases´ transformations and oxidation

behavior of these materials, the coatings were XRD-characterized using an in situ heating

chamber from 600 °C up to 1020 °C every 30 °C.

Structural evolution of CrN coating during in situ HT-XRD

Based on in situ XRD investigation no structural modifications occurred into the cubic CrN

lattice up to 600 °C (Figure 3.20.).

Figure 3.20. CrN structure’s modification during in situ HT XRD. Film is deposited on SiO2

substrate (left all the structural evolution and right zoom for β-Cr2N)


Chapter 3

99

From 630 °C, small fractions of Cr2O3 (ICDD card 38-1479) were detected, together with a

weak amount of β-Cr2N phase (ICDD card 35-0803) formed due to nitrogen releases, causing

a less dense structure. Some authors reported that in the 600 – 700 °C temperature range is

possible to form already the eskolaite phase Cr2O3 [36, 43]. This oxide scale once formed will

grow continuously accelerated by the nitrogen release as the temperature increases [33].

It has to be pointed out that phase transformation from CrN to Cr2N phase can also involve

non-thermodynamic factors, such as stress relaxation-induced phase transformation [62].

From 700 °C, it is noticed an increase of the main peaks of Cr2O3 phase at 2 ~ 24° and 37 °,

while β-Cr2N phase disappeared at 750 °C, perhaps replaced by the Cr phase, detected from

930 °C.

Further increasing the temperature results in a slight increase in the chromium oxides peak

intensities and inversely a decrease of CrN peaks is observed. At the highest heating

temperature, the film still reveals small fractions of fcc-CrN and Cr phases, while the Cr2O3

becomes the dominant phase.

To summarize, CrN is easily oxidized from 630°C, temperature at which its transformation

into β-Cr2N initiates. Above this temperature, the film loses its stability, which demonstrates

a limited oxidation resistance at high temperatures.

Structural evolution of CrAlN coating during in situ HT-XRD

In the previous section we have seen how the CrN structure evolves and what phase

transformations may be involved during oxidation in air. It is now the moment to study the

effect of Al addition to the CrN lattice following the same procedure.

In figure 3.21. is illustrated the structural evolution of the CrAlN coating during continuous

heating.

Figure 3.21. CrAlN structure’s changes during in situ HT XRD. Film is deposited on SiO2




100

For the films containing aluminum, diffractograms are significantly modified.

At 630 °C, no traces of oxides are indeed detected as compared with CrN coating and it is

necessary to reach 930°C to observe the first tiny peaks. The Al incorporation delays also the

structural transformation into β-Cr2N phase, detected at 780 °C. This late structural

transformation highlights the beneficial effect of Al incorporation into the CrN lattice [42]

[19]. Many authors suggested that the Cr2N phase is the first reaction initiated during

heating in air due to a continuous release of nitrogen [36, 40].

When heated at higher temperatures, structure does not sustain such conditions and the

Cr2O3 oxide would appear at 930 - 960 °C. Finally at 1020 °C, the main contribution on the

XRD patterns is still composed by peaks of the Cr(Al)N solid solution. The presence of this

nitride phase should be attributable to the unaccomplished decomposition process over the

entire coating thickness, highlighting the difficulty for nitrogen to be released [64].

At this temperature, small Al2O3 oxide peaks could be as well observed. We consider that

such a protective oxide layer was probably already formed on the coating surface at lower

temperatures, but its presence would not be recorded owing to its amorphous structure [36].

Several authors reported for the CrAlN coatings in the same temperature (maximum 1000

°C) the existence of a mixed (Al,Cr)2O3 [53, 64, 65] layer, while others mentioned the

presence of some Al2O3 traces [66, 67].

According to previous studies the alumina could be formed as early as 650 – 700 °C and can

maintain its amorphous state for temperatures as high as 900 °C [43].

The presence of the mixed structure of fcc-Cr(Al)N and Cr2N phases shifts the oxidation

resistance to higher temperature, as compared to the CrN coating.

Structural evolution of CrAlYN coating during in situ HT-XRD

The CrAlYN coating undergoes similar structural transformations during heating in air as

compared to CrAlN as shown in figure 3.22.

Figure 3.22. CrAlN structure’s changes during in situ HT XRD. Film is deposited on SiO2



Chapter 3

101

Structural transformations really begin at 780 °C, with formation of both Cr2N and Cr2O3.

The latter phase is more clearly spotted at 930 °C. It is worth mentioning that peak at the 2

~ 25° is at an intermediate position between the peak of Cr2O3 and of Al2O3.

Such a result seems corroborate the mixed nature of the oxide: (Cr,Al)2O3. By comparison

with CrAlN, it appears that yttrium incorporation would earlier initiate the global oxidation

of Al2O3 and impulse the formation of the oxide layer on the films’ surface.

Even more, several authors showed that oxide scale formed on the films surface could

contain as well chromium, Y2O3 and/or Cr2O3 grains [40, 67].

Zhu et al. [7] stated that the Y incorporation into TiAlN coating promoted the formation of

Y2O3 and refined the structure. In our case, Y2O3 oxide was not found, probably due to its

low content in the structure.

Results reported here allow us establishing few remarks linked to the oxidation behavior of

CrN and Al-containing films during in situ HT XRD as will be developed in the following.

The Al and/or Y additions form metastable solid solution in the CrN lattice [19, 31, 41, 42,

50] and refines the structure, improving the structure stability and oxidation resistance.

It is well accepted that the structural transformations control the oxidation behavior and

long term stability for such coatings [49, 53, 58, 59].

A further finer chemical characterization could still endorse the nature of the oxides formed

after in situ XRD annealing.

3.6.2. Real time oxidation observation by Environmental-SEM

The XRD characterization gives only information on crystallized phases and when their

content exceeds at least some percent of the sample volume.

In order to put into the light the oxidation mechanism throughout the process, a real time in

situ approach is recommended, as already successfully tested in previous studies on TiN-

based coatings [14, 68].

By the means of SEM used in environmental oxidative conditions, we were interested to

further study the superficial evolution of films deposited on M2 steel during its oxidation.

This approach is complementary with the in situ XRD and we expected to observe the

oxides’ formation and morphology’s changes. Nevertheless, for security reasons, it was not

possible to exceed a temperature of 900°C.

Tests were then performed from room temperature up to 900 °C in the water vapor

atmosphere during continuous heating.

Sequences recorded during these experiments are presented in figure 3.23. A and B for CrN

and CrAlYN, respectively.



102

Figure 3.23. Sequences recorded during continuous heating by ESEM for CrN coating (A) and

CrAlYN (B). The yellow arrows highlight the filament-like features and the red circle a droplet

For the M2 substrate coated with CrN, the surface does not reveal any signs of oxidation

below 700 °C, temperature at which some filament-like features are spotted on the surface.

These fine batons are believed to be the Cr2O3 oxide as confirmed by in situ XRD results

where this phase is detected at 750 °C. Several studies reported the formation of chromium

oxide phase in the 600 – 700 °C temperature range [36, 43].

From 800 °C, the surface begins to be covered by large oxidized grains and the previous

features seem to be enclosed on the layer formed. At 900 °C, the surface is totally oxidized

and covered by a coarser rough grain-like morphology.

The Al-containing films behave in the similar way when heating until 900 °C and thus in

figure 3.23.B only the sequences of CrAlYN coating will be presented.

The film’s surface does not significantly change during continuous heating up to 900°C,

implying a better oxidation resistance.

It is finally required to reach 900 °C to observe the first surface modifications, limited

roughly to a surface’s roughness change: the film exhibits then a grain-like morphology, not

really different to the initial one, with much smaller oxides grains as compared to CrN

coating. Hence for Al-containg films the surface endured only minor oxidation as confirmed

by the post mortem SEM micrographs shown in figure 3.24.


Chapter 3

103

Figure 3.24. SEM micrographs of post ESEM-oxidation for CrN (a) and CrAlYN (b) at different

magnifications

Based on SEM micrographs the surface of CrN exhibit a grain-like well-crystallized

morphology with some large interconnected batons, endorsing the previous state from the

real time oxidation tests. The coating surface seems to be almost completely covered with

this type of features. Comparing with CrN coating, the Al-containing films seems to

undergo less modifications of the surface state. The SEM micrograph obtained after

oxidation inside ESEM illustrates the presence of flower-like morphology.

Based on the analysis of SEM micrographs from Figures 3.23. and 3.24., the obtained

morphologies imply the formation of different type of oxide layer for the Al-containing

films. It is well possible that a preferential formation of an amorphous Al2O3 layer occurs,

covering the surface, which can hinder the clear development of the oxidation products.

The films structure is furthermore analyzed by XRD as shown in figure 3.25.

Figure 3.25. XRD patterns of CrN coating (a) and CrAlYN (b) after oxidation performed inside

ESEM



104

The XRD patterns suggests the formations of oxide layer on the coating surface composed

by a Cr2O3 + Fe3O4 + Fe2O3 mixed scale for CrN and Cr2O3 + Fe3O4 for CrAlYN.

For CrN coating, the nitride Cr2N phase is still spotted but only as minor contribution, while

the initial CrN is fully transformed.

Regarding the Al-containing film the Cr(Al)N phase is still well defined after oxidation at

900 °C, denoting the better oxidation resistance together with a high structural stability.

To summarize, the high reactive atmosphere during ESEM experiments led to an almost

totally oxidation of CrN film at 900 °C. Secondly, the detection of some traces of Fe-based

oxides could suggest a diffusional process from the substrate which has to be further

investigated. At the final temperature step a complex oxide layer is indeed formed on the

surface which could block the inward diffusion of oxygen.

The XRD findings of Chawla et al. [69] for AlCrN and TiAlN coatings after oxidation at 900

°C revealed also the existence of Fe2O3 phase, while Polcar et al. [40] reported the possible

existence of the same phase after oxidation at 1300 °C for CrAlN.

The oxidation behavior during in situ HT XRD and ESEM real time observation considering

both general and local levels were the first attempts to evaluate the oxidation resistance of

the CrN-like coatings. Yet for a quantitative approach it is essential to have access to other

parameters such as activation energy or kinetic constants.

The thermal TGA analysis could provide such responses, as it will be presented in the

following section.

3.6.3. Kinetic approach of the oxidation

In this section the oxidation resistance of the coatings is studied in several modes with the

aim to better understand and explain the oxidation behavior.

Oxidation in dynamic mode

The oxidation test consists in heating the coated steel from 500 °C to 1100 °C with the

purpose of getting insights on the onset critical temperature of oxidation.

In figure 3.26. is represented the thermograms obtained during oxidation for all three

coatings.

a)

b)

Coating Tc

(°C)

CrN 677

CrAlN 656

CrAlYN 618

Figure 3.26. The TGA thermograms obtained in dynamic mode from 500 °C to 1100 °C (a) and

their interpretation in terms of oxidation temperatures (b)


Chapter 3

105

From the figure it can be seen that whatever the coating, there is any mass variation below

650 °C. The critical oxidation temperature is approximated based on a mathematical model

presented in ref. [70]. Based on this approach the critical oxidation temperatures are

presented in the table from figure 3.26.b. It was expected that the critical oxidation

temperatures to be higher for Al and/or Y containing films which is not observed here.

These differences could be explained by the presence of larger and deeper porosities

observed on the surfaces of coatings (see table for porosity values figure 3.2.) which induces

an early oxidation of the surface. Based on the results from the dynamic oxidation tests it

could be implied that there is no beneficial effect of Al and/or Y incorporation into CrN,

such measurement being mainly impacted by the continuity defects.

Moreover, if we consider further the standard energy of reaction of the oxides (see table 3.6.)

the thermodynamics could better clarify the evolution during the dynamic oxidation.

Table 3.6. Standard enthalpy of formation for different types of oxides

Compound ∆𝑯𝒇

𝟎

(kJ/mol)

Fe2O3 - 822

Fe3O4 - 1120.9

Cr2O3 - 1128.4

Al2O3 - 1669.8

Oxidation behavior in isothermal mode

In order to get more insights on the oxidation behavior of the CrN and Al-containing thin

films, isothermal experiments are performed from 750 to 950° every 50°, and the resulting

thermograms are shown in figure 3.27.

Based on in situ XRD results, we decided to choose several temperatures, which allow

investigating the oxidation resistance in a wide temperature range (750 – 950 °C).

We have deliberately oxidized the samples for 4 h to ensure severe conditions of oxidation.



106

Figure 3.27. Mass variation per unit surface as a function of time recorded every 50 °C in the 750

– 950 °C temperature range for all three coatings. Graphs are plotted using the same scale for a

comparison purpose

For all coatings the mass gain increases with the temperature and time. We can observe that

thermograms display the first two hours a strong tendency towards a parabolic evolution

(Δm = kp x t1/2), suggesting elemental processes linked to the diffusion.

The oxidation rate at 750 °C is quite negligible and some slight differences appear at 800 °C

with a somewhat higher oxidation for the quaternary coating. Up to 850°C, films’ mass

gains are below 2 mg/cm2 after 4 h oxidation time, which is consistent with the findings of

previous studies [6, 33, 43].

From 850 °C the CrN coating gain double in mass as compared with Al-containing films,

implying a poor oxidation resistance from this temperature. Consistent is then the fact that

the latter film has the highest mass increase in the 900 – 950 °C temperature domains,

denoting its limited protectiveness.

In opposite, addition of Al and/or Y affords a strong beneficial effect. It is important to

underline that such a high protective character in so severe conditions is furthermore

obtained with a heterogeneous and highly defective surface. Therefore, it can be supposed

that the greater extent of oxidation resistance at high temperatures given by Al and/or Y

could be the consequence of a chemistry effect and/or of a finer film’s microstructure.


Chapter 3

107

Considering a parabolic variation of mass gains with the temperature, kinetic parameters

can be calculated for each temperature.

The imposing effect of aluminum appears then at 850°C and a further improvement due to

yttrium is evidenced at 950°C (Figure 3.28.).

Figure 3.28. Parabolic kinetic constant of each sample at different isothermal oxidation

temperatures

Within the 750 – 800 °C temperature range, all coatings evidence a good protectiveness with

very low mass gains. The higher values of the kp for CrAlYN at 800 °C, are probably related

to the roughness and defective nature of its surface. The strong and preferential reactivity of

pores for arc coatings with high porosity rates (0.7% for CrAlYN here) has already been

attested for TiN-based coatings [14].

Figure 3.29. Sequence of the oxidation process of a TiN-based coated part a) after 120 min at 950

°C (circles highlight two droplets subject to the early oxidation), b) after 140 min at 1000 °C, c)

after 160 min at 1000 °C and d) after 180 min at 1000°C [14]

750 800 850 900 9500.0

5.0x10-5

1.0x10-4

1.5x10-4

2.0x10-4

2.5x10-4

3.0x10-4

Temperature /°C

kp / m

g/c

m2.s

1/2

CrN

CrAlN

CrAlYN



108

At higher temperatures, oxidation reactions begin to be affected by the film’s stability, itself

dependent on the nitrogen release [36, 47, 62].

From 850 °C, the kinetic parameters as plotted in figure 3.28. are lower for the Al-containing

thin films compared with the CrN.

For the CrN coating, from this temperature the decomposition reactions are accelerated due

to the almost fully nitrogen release and formation of the oxide scale as observed from the

mass gain evolution and XRD patterns (Figures 3.20. and 3.27.). The oxide scale formed by

CrN is reported to be less compact, which could produce short diffusion paths, allowing still

a faster diffusion of reactants.

This assumption is confirmed by the increase of the kinetic constant values as the isothermal

oxidation temperature increases (900 – 950 °C temperature range).

The better kinetic parameters noted at high temperatures for Al-containing films can be

explained by the strong Al-N covalent nature of bonds, hindering the nitrogen loss and

consecutive oxidation [65]. Another explanation may rely on the possible diffusion of

yttrium ions at the grain boundaries, thereby limiting even more the general diffusion

process for the CrAlYN coatings [71].

Besides, considering the diffusion mechanism, it is expected that the oxide scales formed for

Al-contain films are denser and more compact [43, 65, 67] than the one formed on CrN.

At the same time, considering a linear evolution of thermograms, we can plot the graph

according to the Arrhenius’ equation (Figure 3.30.) and calculated the activation energy

parameter (table 3.7.).

Figure 3.30. Arrhenius plot ln (kp) vs. 1/T for the CrN-based coatings based on [750 – 950 °C]

thermograms

The linear regression and the slope values suggest a similar transport mechanism by which

oxidation reactions occur. The calculated values are compared with the already reported

ones from literature in table 3.7.

8.0x10-4

8.4x10-4

8.8x10-4

9.2x10-4

9.6x10-4

1.0x10-3

-14

-13

-12

-11

-10

-9

-8

-7

CrN

CrAlN

CrAlYN

ln(k

p)

(mg

/c

m2.s

-1)

1/T (K-1)

CrN: y = -21660x + 9.30

CrAlN: y = -20992x + 8.80

CrAlYN: y = -17001x + 5.20


Chapter 3

109

Table 3.7. Calculated activation energies for our CrN-based coatings compared with values

reported in literature

Coating Activation energy (kJ/mol) References

CrN

176 ± 23 This study

186 Panjan et al.,Thin Solid Films 281, 298, (1996)

163 Otani and Hofmann, Thin Solid Films 287, (1996)

110 Chen and Lu, Thin Solid Films 515, 2179, (2006)

CrAlN

174 ±10 This study

179.4 Li et al., Oxid Met 68, 193 (2007)

130 Reiter et al., J. of Vacuum Sc.& Tech. A, 25, (2007)

CrAlYN 142 ± 11 This study

The values of activation energies obtained for CrN (176 ± 23 kJ/mol) and CrAlN (174 ± 10

kJ/mol) coated steel are in good agreement with the results from other studies.

It can be seen that even though CrAlYN exhibit lower mass gains and hence higher

oxidation resistance than CrAlN, which themselves resist better that CrN, it seems that

oxidation is easier to be activated for this coating.

This can be explained through two origins: a microstructural reason from a one hand in

relation with the porous nature of the surface, a chemical reactivity reason on the other

hand, Y being able to favor the preliminary formation of the oxide scale.

Such oxide layer affords to the underlying coating a protective character, it is not paradoxal

to have simultaneously a low activation energy combined with a high oxidation resistance.

The reported values are generally those of films obtained by PVD magnetron sputtering

method, which give rise to less surface droplets and open porosities.

The studies available in literature do not give any information on the coatings porosity rates,

parameter which can greatly influences the values obtained.

Based on the evolutions from figure 3.31., three temperature domains can be distinguished:

regime I for 750 C and 800 C, a transition regime between 850 and 900 C and regime II for

temperatures higher than 900 C as illustrated in the figure below.

Figure 3.31. Evolution of the final mass gain for the CrN and Al-containing films, after the 4-hour

oxidation



110

In regime I, oxidation conditions can be considered as quite mild and not discriminant with

a good protectiveness for all coatings. For us, morphological aspects of coatings play the

predominant role in this region. Above 850 C, the CrN coating starts to rapidly gain more

mass pointing out its oxidation resistance limit.

In the transition regime T, both Al and Y containing films have equivalent low mass gains

whereas the mass gain of CrN film at 900 C doubled.

The positive effect of addition of Al and/or Y into CrN film begins to be really and strongly

noticed mainly in regime II where the CrN coating has by far the highest mass gain. At the

highest oxidation temperature (950 C) Y induces a further oxidation protection reflected by

much lower mass gains than CrN and even than CrAlN.

The better oxidation resistance of CrAlN compared to CrN agrees with the previous reports

[65, 72 - 75]. The Al benefits to an easier/earlier formation of a dense (Cr, Al)2O3 protective

mixed oxide scale (either amorphous or crystalline) developed on the film surface, which

acts as an effective diffusion barrier slowing down the inward diffusion of the oxygen at

high temperatures [8, 36 – 38, 40, 43, 53, 64, 67].

In the case of CrAlYN coating, the yttrium endorses a faster formation of a dense Al2O3 layer

which restrains the diffusion of the oxygen and hence protects the coating from further

oxidation [76 - 79]. In the studies of Rovere et al. [8] and Qi et al. [9], it is stated that Y favors

the preferential oxidation of stable Al2O3 over Cr2O3 with an Al rich outer oxide layer on top

of a Cr-rich oxinitride layer. Some authors specified that Y atoms may diffuse out

concomitant with Cr and Al, forming a more efficient protective oxide layer [71].

Moreover, Zhu et al. [7] stated that Y incorporation promotes the formation of Y2O3 and

Al2O3 and hinders the grain boundary outward transport of Ti and Al cations for a TiAlN

coating.


Chapter 3

111

3.6.4. Qualitative approach of oxidation behavior

Films’ morphology after isothermal oxidation

The morphological characterization after isothermal oxidation was done by SEM to analyze

the evolution of surface morphology during isothermal oxidation.

The SEM micrographs for selected representative temperatures are presented in figure 3.32.

CrN CrAlN CrAlYN

Figure 3.32. Top view SEM micrographs of coatings deposited on M2 steel substrate after

thermogravimetric tests at several temperatures

In accordance with thermal changes recorded by TGA, morphology of the CrN and Al-

containing films is also modified. The surface morphologies of coatings oxidized for 4h at

750 °C (and 800°C) are very similar to that of the as-deposited coatings confirming that this

temperature corresponds to mild conditions.



112

Starting with 850 °C the surface is transformed into a rock-like morphology with different

size and shapes.

Some fine batons-like features are also spotted. It can be observed that by increasing the

oxidation temperature morphological changes are promoted from 850 °C for CrN and

delayed to 900 °C for Al-containing films.

With an increase of temperature, coarser oxide grains with disk-like shape and irregular

forms appeared, embedded in the fine-grained oxide matrix.

Interesting to note is that at 900 °C the surface seems to be formed by fine grain oxidized

features which further on, at 950 °C enlarge considerably and increase their number.

At these temperatures the morphology is composed by a mixture of fine granular and

faceted structures, uniformly distributed on the film surface.

The SEM micrograph after dynamic oxidation test, conducted up to 1200°C, reveals similar

well-crystallized baton-like features.

The SEM micrographs of CrN and Al-containing film show that under isothermal oxidation

at moderate temperatures (750 – 850 °C) the surfaces undergo only slight morphology’s

modifications, implying a similar oxidation resistance. Yet at higher temperatures (> 900 °C)

the CrN’s surface strongly transforms while other ternary and quaternary films better resist.

Finally, these results confirm the thermogravimetric trend as highlighted in figure 3.27.

Observation of oxidized films deposited on Si substrate and considered in cross section is

more adapted for a deeper mechanistic approach.

In figure 3.33. only some selected temperatures have been chosen to be presented since no

significant differences have been noticed during SEM investigations.

CrN CrAlN CrAlYN

Figure 3.33. SEM cross section micrographs of CrN and Al-containg films deposited on Si

substrate at selected temperatures On top surface thin oxide layer is formed (< 200 nm)


Chapter 3

113

The main remark is that for all coatings a superficial oxide layer has been formed on the

films’ surfaces from 750 °C. The thin and dense oxide overlayer is accompanied also by

microstructural modifications, especially for the CrN sample. For the latter coating,

columnar initial grains are transformed into a micro-granular structure, probably due to the

formation of Cr2N phase and subsequent film oxidation.

At 900 °C, the CrN is even more transformed due to the higher quantity of released

nitrogen. The film underwent here the final structural transformation as confirmed by the

appearance of pure Cr phase in the XRD pattern from figure 3.20.

At the final oxidation temperature, 950 °C, the morphology of CrN endures major changes

owing to the growth of Cr2O3 grains, resulting from a grained porous structure that could be

a further harmful factor leading to its inferior oxidation resistance compared with the other

Al-containg films.

Regarding the Al-containing films, the fine columnar-like growth is still well preserved

whatever the isothermal oxidation temperatures. For these films, the structure does not

evolve a lot as revealed as well by XRD spectra (Figures 3.21., 3.22.).

We decided to choose coatings oxidized at 950 °C and deposited on the M2 steel, as model

conditions for further analyzing the oxidation behavior.

3.6.4.1. Structural and chemical analysis after isothermal oxidation at 950 °C

Here we will examine the structural and chemical behavior of each coating after isothermal

annealing at the most extreme isothermal oxidation temperature 950 °C.

CrN

In figure below is presented the XRD and cross-section for CrN coating.

Figure 3.34. XRD patterns (a) and cross-section view of CrN after isothermal oxidation at 950 °C

for 4 h (b)

After annealing the reflections of substrate’s elements (indicated Fe) are also spotted.

The nitride initial phase seems still present but with a tiny contribution corresponding to

remaining traces of CrN and Cr2N.

From our interpretation, the isothermal annealing weakens the bonding between grains,

resulting in a granular structure and formations of pores in the films as illustrated in SEM

cross-section micrograph from 900 °C (Figure 3.33.).



114

Figure 3.35. Chemical EDS maps of CrN deposited on M2 steel substrate obtained after

isothermal oxidation at 950 °C for 4h

On its surface an oxide scale with a complex chemical composition is formed since several

peaks indexed as Fe2O3, Fe3O4, Cr2O3, CrO3 and Al2O3 oxides are detected in XRD pattern.

The chromium oxide appearance after annealing has been well documented as being an

exothermic reaction that normally takes place above 700 °C as reported elsewhere [33, 35 –

37, 43, 63, 80].

This complicated indexation infers the development of a complex oxide layer, where the

diffusion of Fe-based species could have an impact on the oxidation resistance.

The formation of the Fe-based oxides on the films surface could suggest an outward

diffusion of the substrate elements towards the film surface [34].

Secondly, the chemical data confirm the full transformation of nitrides into oxide phases

(Figure 3.36.). Elemental analyses attests indeed the absence of nitrogen in the outer layer,

but also show a substrate sublayer depleted in iron. The Fe-based oxides detected by XRD

are probably coming from this diffusion.

Figure 3.36. Chemical profiling analysis of the CrN after oxidation at 950 °C for 4 h (left EDS and

right GDOES)


Chapter 3

115

The EDS profile clearly confirms the formation of chromium oxide, as seen as well from the

GDOES profile. Secondly, it seems that iron tends to diffuse from the substrate to occupy the

empty place let by Cr (which at its turn diffused on the surface to form Cr2O3 phase) in the

CrN-lattice. The release of nitrogen during oxidation destabilized the structure and the non-

oxidized Cr species diffuse in the iron-based substrate.

This assumption is consistent since their intrinsic characteristics: atomic sizes (1.26 Å for Fe

and 1.28 Å for Cr) and bcc-lattice parameters (2.86 Å for Fe and 2.88 Å for Cr) are very close,

together with a Fe-Cr phase diagram showing an infinite mutual dissolution of both metals.

Hence, an alloy FeCr is probably formed at the interlayer level between film and substrate

as highlighted by the GDOES from the figure above.

CrAlN

For the ternary film the isothermal oxidation at 950 °C for 4 hours only slightly transform

the structure as the {111} diffraction plane of the nitride solid solution phase is still strongly

detected, while the contribution of oxide’s peaks is far lower (Figure 3.37.a).

Moreover, SEM cross-section micrographs from figure 3.33. (on Si substrate) certify the

better structural stability, with an obvious columnar microstructure and a thin and

superficial oxide scale as highlighted in figure 3.37. b.

Figure 3.37. XRD patterns (a) and cross-section view of CrAlN after isothermal oxidation at 950

°C for 4 h (b)

The Al incorporation into CrN lattice would endorse the thermal stability, thanks to

structural transformations made more difficult in the ternary (Cr, Al)N solid solution [19, 31,

41, 42, 50, 51]. Main oxide phases are Cr2O3 and Fe3O4, while the Al2O3, Fe3O4, CrO3 are as

well presented but at inferior quantities. Several peaks lay between the Al-rich and Cr-rich

oxides at 2θ ~ 24.7°, 33.9°, which again imply presence of a mixed oxide scale (Cr,Al2)O3.

Several authors confirmed the formation of a thick mixed oxide scale (of about 500 nm)

where both Al and Cr species [43, 64, 65, 67] and even Fe species [34, 39] were detected.



116

If higher affinity of aluminum to oxygen compared to chromium (more thermodynamically

stable) should lead to the formation of a dominant Al-oxide scales, experimental peaks

attributed to Al2O3 have lower intensities due to its probable still amorphous nature at 950

°C. The cross-section SEM micrograph from figure 3.37.b shows the formation of an oxide

on the surface of a thickness of about 300 nm. The features have different sizes, which could

correspond, as suggested by XRD, to different oxides.

The oxide nature and thickness after isothermal oxidation at 950 °C for 4h can be easily

observed from the EDS chemical maps shown in figure 3.38.

Figure 3.38. Chemical EDS maps of CrAlN deposited on M2 steel substrate obtained after


The nitrogen, aluminum and chromium are still detected on the whole film thickness, which

clearly implies the presence of the nitride-like solid solution.

On the film’s surface, the top layer is greatly enriched in oxygen, corresponding to the dense

protective oxide layer. This oxide scale blocks efficiently the inward diffusion of oxygen and

outward diffusion of the species such as nitrogen.

Considering that Cr possess a higher mobility that Al [9], Cr may out-diffuse to the top layer

to form Cr-rich oxide scale, but also migrate to the substrate to associate with iron as

suggested by the Cr richer zone measured by EDS at the substrate / coating interface for the

CrN coating (Figure 3.36.). as seen also for CrAlN in figure 3.38.


Chapter 3

117

CrAlYN

The XRD pattern of CrAlYN coating oxidized at 950 °C is not significantly different as the

one of CrAlN film. The same kind of oxide-based peaks appear in the XRD diffractogram

with the predominant nitride-based phase still intense as illustrated in 3.39.

Figure 3.39. XRD patterns (a) and cross-section view of CrAlYN after isothermal oxidation at 950

°C for 4 h (b)

No peaks belonging to Y2O3 are observed, most likely due to the low Y content (maximum

1.5 at.%). Such compound was determined by Qi and colleagues by XPS an Y2O3 on films

(1.7 at.% Y) oxidized 2 hours in air at 1100 °C [9].

However, for such high Y contents, this element affects the morphology of the (Al,Cr2)O3

scale, making it porous and no more protective [8]. Oxide scale for CrAlYN is still thinner

than for CrAlN (170 and 380 nm respectively), endorsing once more the further beneficial

effect given by yttrium.

These results are in good agreement with the profiles presented by other authors for CrAlN

[73, 74] and CrAlYN [8, 9, 71].

Figure 3.40. Chemical EDS maps of CrAlYN deposited on M2 steel substrate obtained after




118

From the chemical maps we can observe that oxygen signal is only concentrated on the top

surface of the coating, corresponding to the oxide scale. Metallic elements, namely Al and

Cr, are uniformly distributed through the entire thickness, which validates indeed that film

structure has not been transformed during isothermal oxidation at 950 °C.

Another important observation is the uniform distribution of Y throughout the whole film’s

thickness.

For complementary and deeper analysis, EDS and GDOES profiles are exposed figure 3.41.

Figure 3.41. Chemical profiling analysis of the CrAlYN after oxidation at 950 °C for 4 h (left EDS

and right GODES)

The oxide layer formed is a duplex (Al,Cr)2O3 with and Al-richer outer oxide layer and Cr-

richer inner layer has a lower thickness (~ 250 nm) compared with CrAlN coating as

estimated from EDS chemical maps (Figure 3.40.) and GDOES profile (Figure 3.41.).

We can see also that the diffusion mobility of iron versus Cr is confirmed near the

coating/substrate interface where Cr local enrichment is detected.

In opposite, Al levels were unchanged, attesting its more difficult diffusivity into the

coating. This differential diffusion mobility between Cr and Al influenced by yttrium

remains an open question.

Regarding the results from the previous sections, if we consider the in situ XRD structural

evolution, morphology changes, oxidation resistance or chemical profiling analysis we

could say that the CrAlYN is the best solution for industrial applications amongst the three

studied film.

In the following section we would like to concentrate our attention further on the

microstructural modifications endured by this coating during isothermal oxidation at 950°C.

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

0

10

20

30

40

50

60

70

80

90

100

depth (m)

CrAlYN

co

mp

osit

ion

/ a

t.%

M2 Substrate

FeNO

Y

CrAl


Chapter 3

119

3.6.5. Microstructural investigation of the CrAlYN coating oxidized at 950 °C

For a more detailed analysis of the oxide scale formed after 4 h oxidation at 950 °C, FIB/TEM

cross-sections were specifically prepared. Also, qualitative EDS analyses were acquired in

cross-section to elucidate the complex chemical composition of the oxide scale.

In figure 3.42. TEM cross-section image of the CrAlYN coating after oxidation at 950 °C from

the substrate until the oxide layer is shown.

Figure 3.42. FIB cross-section TEM showing a general view for the CrAlYN coating after

oxidation at 950 °C for 4 h deposited on M2 steel substrate

The figure highlights that the film microstructure seems unaffected by the oxidation since

the columnar-like growth is clearly observed. On the top of the film the oxide layer with a

thickness of about 240 nm matches well with the GDOES profile and SEM micrographs

shown before. An essential analysis is to control the oxygen content after isothermal

oxidation inside the film. Hence an EDS chemical analysis with a quite large probe’s size

(0.35 μm) was performed in the middle of the coating (Figure 3.43.).

Chemical

element

Content

(at.%)

Al 35

Cr 12

N 40

Y 1.1

Fe 6

O 6

Figure3.43. TEM micrograph in the middle of the CrAlYN film far away from the substrate and

surface. The circle represents the probe size (0.35 μm) used for the EDS analysis



120

The coating preserved its initial chemical chemistry after isothermal oxidation at 950 °C.

The ratio of the Al(Al+Cr) is 0.61 which is the same as the one measured in the as-deposited

state and really close to the target initial ratio of 0.64. The yttrium is also well quantified at

1.1 at. %, value really close to the one of the multicomponent target (1.5 at. %). The iron and

oxygen are present as well, but their values are not really significant.

The second observation is that the film still exhibits a complex polycrystalline structure with

numerous diffracted domains as shown in figure 3.44. (black circles) and similar with the

one from as-deposited presented previously in figure 3.18. The SAED is taken in the same

region as the EDS chemical analysis. In addition to nitride compound, some hypothetic

further red circles could be indexed as the Cr2O3 oxide (induced perhaps by a post

oxidation of thin foil after preparation).

Figure 3.44. A zone where SAED has been performed for the film after oxidation at 950 °C for 4h

illustrated previously in figure 3.49. The indexed black circles are related to f.c.c.-Cr(Al)N solid

solution and red circles to rhomb.-Cr2O3

These main spots are related to a complex structure indexed in the {111}, {220} plans of a fcc

Cr(Al)N structure. The microstructure of this film is rather complex, thus further detailed

analysis has been performed. The figure 3.45. illustrates a cross-sectional TEM image

obtained in HAADF mode.

Figure 3.45. General HAADF TEM image of CrAlYN coating after oxidation at 950 °C for 4h. The

highlighted zone will be analysed separately


Chapter 3

121

The figure shows evidence of a columnar structure with zone having different chemical

compositions since the contrast is different. With the aim of obtaining complementary

information on the chemical and microstructural changes due to oxidation at 950 °C, further

deeper investigations were carried out in the A region (Figure 3.46).

Figure 3.46. Bright field TEM image of the selected A zone from previous figure 3.45.

From the figure it can be seen many submicron contrasted grains, differentiated by their

own chemistry. EDS analysis results performed with a beam probe of 20 nm are gathered in

table 3.8.

Table 3.8. EDS chemical analysis from the zone A (Figure 3.45.) performed inside TEM on the

CrAlYN after oxidation at 950 °C for 4 h for different beam probe on figure 3.46.

Chemical

element

(at. %)

Position

1 2 3 4 5 6 7

Al 31 32 31 31 29 33 35

Cr 18 18 18 18 12 14 18

Y 0.6 0.8 0.6 0.9 0.9 0.8 1.0

Fe 0.4 2 0.4 2.1 15 8.1 1.0

N 50 45 50 48 39 41 45

O 0 2.2 0 0 4.1 3.1 0

It is important to underline that contrary to what was already published [76 - 79] the yttrium

contents measured within the grain, as well as at the grain boundaries, are similar which

discredits segregation in the intergranular space. Within grains (positions 1 – 3) the same

chemical composition as the one of as-deposited coating is preserved with a 0.63 ratio for

Al/(Al+Cr) contents. Moreover, we may notice that the values of Fe and O the are really low

excepted when considering the positions from 4 to 7 on both sides of the grain boundary,

the iron reaches a maximum content of 10-15 at.% at the intergranular region.



122

In order to assess the chemical evolution in zone A (Figure 3.45.) a line scan analysis was

carried out as shown in figure 3.47.

Figure 3.47. EDS chemical profile analysis in zone A on the CrAlYN oxidized at 950 °C for 4 h

It is clear that yttrium has a homogeneous chemical distribution along the whole film as

confirmed also by the EDS line scan and GDOES profile (Figure 3.41.). The Al and Cr have a

similar chemical evolution, while the iron was surprisingly detected as well.

The presence of iron at this location so far from the substrate could be only explained by a

diffusional mechanism which yields during the isothermal oxidation. The activation of the

dislocations sites during heating could be another explanation for the detection of iron so far

from the substrate. The high compressive stress from the as-deposited state could group

these dislocations and form paths for diffusion of the substrate species.

The exchange was more observable for the CrN coating after oxidation at 950 °C, where a

FeCr alloy was formed at the interface of the film-substrate as confirmed by the GDOES

chemical profile (Figure 3.36.).

Some authors proposed theoretical models adapted to nanostructured materials to explain

the mechanism of lattice expansion that leads to crystal lattice deviations from its perfect

structure [81 - 83]. They concluded that at the grain boundaries a certain amount of excess

volume in a form of vacancies could generate defects due to the disordered arrangement of

the atoms.

Besides, EDS chemical analyses of CrAlYN oxidized 4 hours at 950 °C, also revealed that

iron was detected as well among constitutive elements of the oxide layer. We did also fine

analyses of the outer oxide scale, as presented in figure 3.48.


Chapter 3

123

Chemical

element

Content

(at.%)

Al 18

Cr 14

N 0

Y 0.2

Fe 4.5

O 63

Figure 3.48. Top oxide layer of the CrAlYN film after oxidation at 950 °C for 4 h together with

the table showing the chemical contents. The circle represents the beam size (50 nm) used for the

EDS analysis

The high content of oxygen of about 63 at.% on the top of the CrAlYN after long oxidation at

950 °C confirms the formation an oxide layer composed mainly by Al, Cr and apparently in

this area of some Fe species. Not surprising is the absence of nitrogen, considering that the

analysis has been done on the oxide scale.

The yttrium content is measured about 0.2 at. %, which seems sufficient to improve

oxidation resistance and give rise to a well-adherent oxide scale.

In two other selected zones from the oxide formed on the CrAlYN surface, the ratio of

chemical elements are rather changed as presented in figure 3.49.

Chemical

element

(at. %)

Position

1 2

Al 9.5 29

Cr 20.5 5.4

Y 0 0.8

Fe 2.7 0.5

N 0 0.7

O 67 64

Figure 3.49. Top oxide layer of the CrAlYN film after oxidation at 950 °C for 4 h together with

the table showing the chemical contents. The circle represents the beam size (50 nm) used for the

EDS analysis

In this area, the EDS chemical analysis exhibits differences in terms of oxide composition.

From the table inserted in the figure, it can be seen that oxygen has a constant value around

65 at.%, while the contents of other metallic elements are evolving differently. Considering

the EDS findings, in position 1 a “rather chromium oxide-based crystal” is grown, while the

second point indicate a “rather aluminum oxide-based crystals.



124

At the same time, the GDOES profiles (Figure 3.56.) confirmed the formation of an Al2O3

oxide-rich scale with a thickness of about 0.3 μm and therefore, here, using this local

investigation, the EDS results revealed the same oxide scale chemistry. Hence based on the

findings from the TEM microstructural investigation of the CrAlYN coating after oxidation

at 950 °C for 4 h and corroborated with other results, we can state that the oxide scale is

composed of a mixed oxide characterized by a large dissolution domain.

The microstructural and chemical analysis of the CrAlYN coating after oxidation at 950 °C

for 4 h showed us that the f.c.c. solid solution structure is well protected against oxidation.

The Al and Cr-oxide layer formed on the film surface successfully blocks the inward

diffusion of oxygen. Among the all three coatings studied the CrAlYN had always the better

thermal stability and was less impacted by the temperature.


Chapter 3

125

Chapter summary

The multiscale approach proposed for studying the oxidation resistance of ceramic coatings

at high temperatures has allowed finding significant insights related to their structural

stability and oxidation resistance.

The arc-PVD coatings are generally well adherent with a highly columnar structure, good

mechanical properties and a final chemical metallic ratio close to the one of the initial

targets. The growth defects, such as droplets, open porosities have been identified as

playing an important role on the oxidation behavior at high temperatures.

The structure of the ceramic films has been studied using an in situ XRD approach to follow

their phase transformations, while the solid solution state was confirmed by TEM results.

The addition of Al and/or Y refined the structure decreasing crystal domains size from 35

nm for classical CrN to 22 – 24 nm for the Al-containing films. Y is found to be well-

dispersed throughout the film thickness.

The in situ real time approach has allowed following the surface oxidation using a global

method, by in situ XRD and by a local surface observation using environmental SEM tests

under vapor oxidative atmosphere. The high reactivity of CrN from 700°C has thus been

clearly established.

From thermogravimetric quantitative measurements, we can conclude that the oxidation

resistance for the three nitrides in the moderate temperature domain (750 – 800 °C) is mainly

influenced by their morphology, while for more severe environments (900 – 950 °C) the

effect of the coating’s chemistry becomes preponderant.

The CrN showed a limited protectiveness of the coated substrate during oxidation at 950 °C

since the film is fully oxidized as showed by the XRD, GDOES or EDS chemical maps.

The aluminum addition increases the structural stability due the different enthalpy of

formation between the CrN (ΔHf = -117 kJ/mol)) and AlN (ΔHf = -318 kJ/mol), which has a

direct effect on retarding the oxidation and inward diffusion of the oxygen.

The better oxidation resistance of the CrAlN compared with CrN can also be related to the

easier/earlier formation of a dense (Cr, Al)2O3 protective mixed oxide scale which may act

as the diffusion barrier layer (thickness of about 500 nm).

Regarding the CrAlYN coating, the addition of yttrium endorsed the preferential formation

of an Al-enriched (Al, Cr)2O3 oxide which may explain the more protective character of the

superficial scale. Another explanation could be based on the lattice distortion of the solid

solution structure (lattice parameter aCrAlYN > aCrAlN) which could make more difficult the

diffusion of the N and/or O.



126

List of references

[1] I. G. Brown, “Cathodic Arc Deposition of Films,” Annu. Rev. Mater. Sci., vol. 28, no. 1, pp.

243–269, 1998.

[2] P. Panjan, M. Čekada, M. Panjan, and D. Kek-Merl, “Growth defects in PVD hard

coatings,” Vacuum, vol. 84, no. 1, pp. 209–214, 2009.

[3] K. S. Sree Harsha, “Chapter 9 - Nucleation and Growth of Films,” in Principles of Vapor

Deposition of Thin Films, K. S. S. Harsha, Ed. Oxford: Elsevier, 2006, pp. 685–829.

[4] P. Panjan, M. Čekada, M. Panjan, D. Kek-Merl, F. Zupanič, L. Čurković, and S. Paskvale,

“Surface density of growth defects in different PVD hard coatings prepared by sputtering,”

Vacuum, vol. 86, no. 6, pp. 794–798, 2012.

[5] W. Ernst, J. Neidhardt, H. Willmann, B. Sartory, P. H. Mayrhofer, and C. Mitterer,

“Thermal decomposition routes of CrN hard coatings synthesized by reactive arc

evaporation and magnetron sputtering,” Thin Solid Films, vol. 517, no. 2, pp. 568–574, 2008.

[6] P. . Mayrhofer, H. Willmann, and C. Mitterer, “Oxidation kinetics of sputtered Cr–N hard

coatings,” Surf. Coat. Technol., vol. 146–147, pp. 222–2282001.

[7] L. Zhu, Y. Zhang, W. Ni, and Y. Liu, “The effect of yttrium on cathodic arc evaporated

Ti0.45Al0.55N coating,” Surf. Coat. Technol., vol. 214, pp. 53–58, 2013.

[8] F. Rovere, P. H. Mayrhofer, A. Reinholdt, J. Mayer, and J. M. Schneider, “The effect of

yttrium incorporation on the oxidation resistance of Cr–Al–N coatings,” Surf. Coat. Technol.,

vol. 202, no. 24, pp. 5870–5875, 2008.

[9] Z. B. Qi, Z. T. Wu, and Z. C. Wang, “Improved hardness and oxidation resistance for

CrAlN hard coatings with Y addition by magnetron co-sputtering,” Surf. Coat. Technol.

[10] T. J. A., “Am. Rev. Mater. Sci. 7 (1977) 239.”

[11] J. Creus, H. Mazille, and H. Idrissi, “Porosity evaluation of protective coatings onto steel,

through electrochemical techniques,” Surf. Coat. Technol., vol. 130, no. 2–3, pp. 224–232,

2000.

[12] C. Petrogalli, L. Montesano, M. Gelfi, G. M. La Vecchia, and L. Solazzi, “Tribological and

corrosion behavior of CrN coatings: Roles of substrate and deposition defects,” Surf. Coat.

Technol.

[13] W.-D. Münz, D. Lewis, S. Creasey, T. Hurkmans, T. Trinh, and W. Ijzendorn, “Defects in

TiN and TiAIN coatings grown by combined cathodic arc/unbalanced magnetron

technology,” Vacuum, vol. 46, no. 4, pp. 323–330, 1995.

[14] A. Mège-Revil, P. Steyer, G. Thollet, R. Chiriac, C. Sigala, J. C. Sanchéz-Lopéz, and C.

Esnouf, “Thermogravimetric and in situ SEM characterisation of the oxidation phenomena

of protective nanocomposite nitride films deposited on steel,” Surf. Coat. Technol., vol. 204,

no. 6–7, pp. 893–901, 2009.

[15] S. Creasey, D. B. Lewis, I. J. Smith, and W.-D. Münz, “SEM image analysis of droplet

formation during metal ion etching by a steered arc discharge,” Surf. Coat. Technol., vol. 97,

no. 1–3, pp. 163–175, 1997.

[16] H. Ljungcrantz, L. Hultman, J.-E. Sundgren, G. Håkansson, and L. Karlsson,

“Microstructural investigation of droplets in arc-evaporated TiN films,” Surf. Coat. Technol.,

vol. 63, no. 1–2, pp. 123–128, 1994.

[17] P. Panjan, P. Gselman, D. Kek-Merl, M. Čekada, M. Panjan, G. Dražić, T. Bončina, and F.

Zupanič, “Growth defect density in PVD hard coatings prepared by different deposition

techniques,” Surf. Coat. Technol., vol. 237, pp. 349–356, 2013.

[18] P. Panjan, D. Kek Merl, F. Zupanič, M. Čekada, and M. Panjan, “SEM study of defects in

PVD hard coatings using focused ion beam milling,” Surf. Coat. Technol., vol. 202, no. 11,

pp. 2302–2305, 2008.


Chapter 3

127

[19] J. E. Sanchéz, O. M. Sanchéz, L. Ipaz, W. Aperador, J. C. Caicedo, C. Amaya, M. A. H.

Landaverde, F. E. Beltran, J. Muñoz-Saldaña, and G. Zambrano, “Mechanical, tribological,

and electrochemical behavior of Cr1−xAlxN coatings deposited by r.f. reactive magnetron

co-sputtering method,” Appl. Surf. Sci., vol. 256, no. 8, pp. 2380–2387, 2010.

[20] E. Le Bourhis, P. Goudeau, M. H. Staia, E. Carrasquero, and E. S. Puchi-Cabrera,

“Mechanical properties of hard AlCrN-based coated substrates,” Surf. Coat. Technol., vol.

203, no. 19, pp. 2961–2968, 2009.

[21] K. Bobzin, N. Bagcivan, P. Immich, S. Bolz, R. Cremer, and T. Leyendecker, “Mechanical

properties and oxidation behaviour of (Al,Cr)N and (Al,Cr,Si)N coatings for cutting tools

deposited by HPPMS,” Thin Solid Films, vol. 517, no. 3, pp. 1251–1256, 2008.

[22] H. Hasegawa, K. Ohashi, S. Tsukamoto, T. Sato, and T. Suzuki, “Characterization of

quaternary (Cr,Al)N-based films synthesized by the cathodic arc method,” Surf. Coat.

Technol., vol. 202, no. 4–7, pp. 786–789, 2007.

[23] L. Aihua, D. Jianxin, C. Haibing, C. Yangyang, and Z. Jun, “Friction and wear properties of

TiN, TiAlN, AlTiN and CrAlN PVD nitride coatings,” Int. J. Refract. Met. Hard Mater., vol.

31, pp. 82–88, 2012.

[24] Y. Birol, “Sliding wear of CrN, AlCrN and AlTiN coated AISI H13 hot work tool steels in

aluminium extrusion,” Tribol. Int., vol. 57, no. 0, pp. 101–106, 2013.

[25] J. L. Mo and M. H. Zhu, “Sliding tribological behaviors of PVD CrN and AlCrN coatings

against Si3N4 ceramic and pure titanium,” Wear, vol. 267, no. 5–8, pp. 874–881, 2009.

[26] M. Brizuela, A. García-Luis, P. Corengia, D. González-Santamaría, R. Muñoz, and J.-J.

González, “Microstructural, Mechanical and Tribological Properties of CrA1YN Coatings

Deposited by Magnetron Sputtering,” Plasma Process. Polym., 6, no. S1, pp. S162–S167, 2009.

[27] F. Rovere and P. H. Mayrhofer, “Impact of yttrium on structure and mechanical properties

of Cr–Al–N thin films,” J. Vac. Sci. Technol. A, vol. 25, no. 5, pp. 1336–1340, 2007.

[28] T. Miyake, A. Kishimoto, and H. Hasegawa, “Tribological properties and oxidation

resistance of (Cr,Al,Y)N and (Cr,Al,Si)N films synthesized by radio-frequency magnetron

sputtering method,” Surf. Coat. Technol., vol. 205, Supplement 1, pp. S290–S294, 2010.

[29] Y. H. Cheng, T. Browne, and B. Heckerman, “Mechanical and tribological properties of

CrN coatings deposited by large area filtered cathodic arc,” Wear, vol. 271, no. 5–6, pp. 775–

782, 2011.

[30] J. J. Zhang, M. X. Wang, J. Yang, Q. X. Liu, and D. J. Li, “Enhancing mechanical and

tribological performance of multilayered CrN/ZrN coatings,” Surf. Coat. Technol., vol. 201,

no. 9–11, pp. 5186–5189, 2007.

[31] J. Lin, B. Mishra, J. J. Moore, and W. D. Sproul, “Microstructure, mechanical and

tribological properties of Cr1−xAlxN films deposited by pulsed-closed field unbalanced

magnetron sputtering (P-CFUBMS),” Surf. Coat. Technol., 201, no. 7, pp. 4329–4334, 2006.

[32] M. Brizuela, A. Garcia-Luis, I. Braceras, J. I. Oñate, J. C. Sánchez-López, D. Martínez-

Martínez, C. López-Cartes, and A. Fernández, “Magnetron sputtering of Cr(Al)N coatings:

Mechanical and tribological study,” Surf. Coat. Technol., 200, no. 1–4, pp. 192–197, 2005.

[33] Z. B. Qi, B. Liu, Z. T. Wu, F. P. Zhu, Z. C. Wang, and C. H. Wu, “A comparative study of

the oxidation behavior of Cr2N and CrN coatings,” Thin Solid Films., 544, 515-520, 2013

[34] D. . Lee, Y. . Lee, and S. . Kwon, “High temperature oxidation of a CrN coating deposited

on a steel substrate by ion plating,” Surf. Coat. Technol., vol. 141, no. 2–3, pp. 227–231, 2001.


reactive direct current magnetron sputtered CrAlN and CrN coatings,” Surf. Coat. Technol.,

vol. 201, no. 6, pp. 2193–2201, 2006.

[36] J. Lin, B. Mishra, J. J. Moore, and W. D. Sproul, “A study of the oxidation behavior of CrN

and CrAlN thin films in air using DSC and TGA analyses,” Surf. Coat. Technol., vol. 202, no.

14, pp. 3272–3283, 2008.



128

[37] Y. C. Chim, X. Z. Ding, X. T. Zeng, and S. Zhang, “Oxidation resistance of TiN, CrN, TiAlN

and CrAlN coatings deposited by lateral rotating cathode arc,” Thin Solid Films, vol. 517,

no. 17, pp. 4845–4849, 2009.

[38] Y. Lv, L. Ji, X. Liu, H. Li, H. Zhou, and J. Chen, “Influence of substrate bias voltage on

structure and properties of the CrAlN films deposited by unbalanced magnetron

sputtering,” Appl. Surf. Sci., vol. 258, no. 8, pp. 3864–3870, 2012.

[39] T. Vieira, J. Castanho, and C. Louro, “Chapter 16 - Hard Coatings Based on Metal Nitrides,

Metal Carbides and Nanocomposite Materials: PVD Process and Properties,” in Materials

Surface Processing by Directed Energy Techniques, Y. Pauleau, Ed. Oxford: Elsevier, 2006, pp.

537–572.

[40] T. Polcar and A. Cavaleiro, “High temperature properties of CrAlN, CrAlSiN and AlCrSiN

coatings – Structure and oxidation,” Mater. Chem. Phys., 129, no. 1–2, pp. 195–201, 2011.

[41] H. N. Shah and R. Jayaganthan, “Influence of Al Contents on the Microstructure,

Mechanical, and Wear properties of Magnetron Sputtered CrAlN Coatings,” J. Mater. Eng.

Perform., vol. 21, no. 9, pp. 2002–2009.

[42] G. S. Kim and S. Y. Lee, “Microstructure and mechanical properties of AlCrN films

deposited by CFUBMS,” Surf. Coat. Technol., vol. 201, no. 7, pp. 4361–4366, décembre 2006.

[43] J. Lin, N. Zhang, W. D. Sproul, and J. J. Moore, “A comparison of the oxidation behavior of

CrN films deposited using continuous dc, pulsed dc and modulated pulsed power

magnetron sputtering,” Surf. Coat. Technol., vol. 206, no. 14, pp. 3283–3290, Mar. 2012.

[44] C. Gautier and J. Machet, “Study of the growth mechanisms of chromium nitride films

deposited by vacuum ARC evaporation,” Thin Solid Films, vol. 295, no. 1–2, pp. 43–52, 1997.

[45] A. Ehrlich, M. Kühn, F. Richter, and W. Hoyer, “Complex characterisation of vacuum arc-

deposited chromium nitride thin films,” Surf. Coat. Technol., 76–77, pp. 280–286, 1995.

[46] P. H. Mayrhofer, G. Tischler, and C. Mitterer, “Microstructure and mechanical/thermal

properties of Cr–N coatings deposited by reactive unbalanced magnetron sputtering,” Surf.

Coat. Technol., vol. 142–144, pp. 78–84, 2001.

[47] P. H. Mayrhofer, F. Rovere, M. Moser, C. Strondl, and R. Tietema, “Thermally induced

transitions of CrN thin films,” Scr. Mater., vol. 57, no. 3, pp. 249–252, 2007.

[48] M. Odén, C. Ericsson, G. Håkansson, and H. Ljungcrantz, “Microstructure and mechanical

behavior of arc-evaporated Cr–N coatings,” Surf. Coat. Technol., vol. 114, no. 1, pp. 39–51,

avril 1999.

[49] H. Willmann, P. H. Mayrhofer, P. O. Å. Persson, A. E. Reiter, L. Hultman, and C. Mitterer,

“Thermal stability of Al–Cr–N hard coatings,” Scr. Mater., vol. 54, no. 11, pp. 1847–1851,

2006.

[50] O. M. S. Quintero, W. A. Chaparro, L. Ipaz, J. E. S. Barco, F. E. Beltrán, and G. Zambrano,

“Influence of the microstructure on the electrochemical properties of Al-Cr-N coatings

deposited by co-sputtering method from a Cr-Al binary target,” Mater. Res., vol. 16, no. 1,

pp. 204–214, 2013.

[51] A. Kimura, M. Kawate, H. Hasegawa, and T. Suzuki, “Anisotropic lattice expansion and

shrinkage of hexagonal TiAlN and CrAlN films,” in Surface & coatings technology, vol. 169–

70, pp. 367–370.

[52] X. Ding, X. T. Zeng, Y. C. Liu, F. Z. Fang, and G. C. Lim, “Cr1−xAlxN coatings deposited by

lateral rotating cathode arc for high speed machining applications,” Thin Solid Films, vol.

516, no. 8, pp. 1710–1715, 2008.

[53] A. E. Reiter, V. H. Derflinger, B. Hanselmann, T. Bachmann, and B. Sartory, “Investigation

of the properties of Al1−xCrxN coatings prepared by cathodic arc evaporation,” Surf. Coat.

Technol., vol. 200, no. 7, pp. 2114–2122, 2005.


Chapter 3

129

[54] F. Rovere, D. Music, J. M. Schneider, and P. H. Mayrhofer, “Experimental and

computational study on the effect of yttrium on the phase stability of sputtered Cr–Al–Y–N

hard coatings,” Acta Mater., vol. 58, no. 7, pp. 2708–2715, 2010.

[55] B. E. Warren, X-Ray Diffraction. Courier Dover Publications, 2012.

[56] J.-W. Lee, S.-K. Tien, and Y.-C. Kuo, “The effects of pulse frequency and substrate bias to

the mechanical properties of CrN coatings deposited by pulsed DC magnetron sputtering,”

Thin Solid Films, vol. 494, no. 1–2, pp. 161–167, 2006.

[57] Y. Chunyan, T. Linhai, W. Yinghui, W. Shebin, L. Tianbao, and X. Bingshe, “The effect of

substrate bias voltages on impact resistance of CrAlN coatings deposited by modified ion

beam enhanced magnetron sputtering,” Appl. Surf. Sci., vol. 255, no. 7, pp. 4033–4038, 2009.

[58] M. Bartosik, R. Daniel, C. Mitterer, and J. Keckes, “Thermally-induced formation of

hexagonal AlN in AlCrN hard coatings on sapphire: Orientation relationships and residual

stresses,” Surf. Coat. Technol., vol. 205, no. 5, pp. 1320–1323, 2010.

[59] Y. Birol and D. İsler, “Response to thermal cycling of CAPVD (Al,Cr)N-coated hot work

tool steel,” Surf. Coat. Technol., vol. 205, no. 2, pp. 275–280, 2010.

[60] Y. X. Wang, S. Zhang, J.-W. Lee, W. S. Lew, and B. Li, “Influence of bias voltage on the

hardness and toughness of CrAlN coatings via magnetron sputtering,” Surf. Coat. Technol.,

vol. 206, no. 24, pp. 5103–5107, 2012.

[61] J. Lin, M. Pinkas, W. D. Sproul, and J. J. Moore, “The phase and microstructure of CrAlN

films deposited by pulsed dc magnetron sputtering with synchronous and asynchronous

bipolar pulses,” Thin Solid Films, vol. 520, no. 1, pp. 166–173, 2011.

[62] C. Paternoster, A. Fabrizi, R. Cecchini, M. E. Mehtedi, and P. Choquet, “Thermal stability of

CrN x nanometric coatings deposited on stainless steel,” J. Mater. Sci., vol. 43, no. 10, pp.

3377–3384, 2008.

[63] A. H. Liu, J. X. Deng, H. B. Cui, J. Zhao, and X. Ai, “Oxidation Resistance of CrN and

CrAlN Coating Tools,” Adv. Mater. Res., vol. 189–193, pp. 137–141, 2011.

[64] A. E. Reiter, C. Mitterer, and B. Sartory, “Oxidation of arc-evaporated Al1−xCrxN

coatings,” J. Vac. Sci. Technol. A, vol. 25, no. 4, pp. 711–720, 2007.

[65] S. Khamseh, “Oxidation Resistance of CrAlN Films with Different Microstructures

Prepared by Pulsed DC Balanced Magnetron Sputtering System,” Mater. Trans., vol. 51, no.

2, p. 271, 2010.

[66] M. Kawate, A. Kimura Hashimoto, and T. Suzuki, “Oxidation resistance of Cr1−XAlXN

and Ti1−XAlXN films,” Surf. Coat. Technol., vol. 165, no. 2, pp. 163–167, 2003.

[67] M. Zhu, M. Li, and Y. Zhou, “Oxidation resistance of Cr1−xAlxN (0.18 ≤ x ≤ 0.47) coatings

on K38G superalloy at 1000–1100 °C in air,” Surf. Coat. Technol., vol. 201, no. 6, pp. 2878–

2886, 2006.

[68] P. Steyer, A. Mege, D. Pech, C. Mendibide, J. Fontaine, J.-F. Pierson, C. Esnouf, and P.

Goudeau, “Influence of the nanostructuration of PVD hard TiN-based films on the

durability of coated steel,” Surf. Coat. Technol., vol. 202, no. 11, pp. 2268–2277, 2008.

[69] V. Chawla, A. Chawla, B. S. Sidhu, S. Prakash, and D. Puri, “Oxidation Behavior of

Nanostructured TiAlN and AlCrN Thin Coatings on ASTM-SA213-T-22 Boiler Steel,” J.

Miner. Mater. Charact. Eng., vol. 09, no. 11, p. 1037, 2010.

[70] P. Steyer, D. Pilloud, J. F. Pierson, J.-P. Millet, M. Charnay, B. Stauder, and P. Jacquot,

“Oxidation resistance improvement of arc-evaporated TiN hard coatings by silicon

addition,” Surf. Coat. Technol., vol. 201, no. 7, pp. 4158–4162, 2006.

[71] T. C. Rojas, S. El Mrabet, S. Domínguez-Meister, M. Brizuela, A. García-Luis, and J. C.

Sánchez-López, “Chemical and microstructural characterization of (Y or Zr)-doped CrAlN

coatings,” Surf. Coat. Technol., vol. 211, pp. 104–110, 2012.



130

[72] R. E. Galindo, J. L. Endrino, R. Martínez, and J. M. Albella, “Improving the oxidation

resistance of AlCrN coatings by tailoring chromium out-diffusion,” Spectrochim. Acta Part B

At. Spectrosc., vol. 65, no. 11, pp. 950–958, 2010.

[73] O. Banakh, P. . Schmid, R. Sanjinés, and F. Lévy, “High-temperature oxidation resistance of

Cr1−xAlxN thin films deposited by reactive magnetron sputtering,” Surf. Coat. Technol., vol.

163–164, pp. 57–61, 2003.

[74] J. L. Endrino, G. S. Fox-Rabinovich, A. Reiter, S. V. Veldhuis, R. Escobar Galindo, J. M.

Albella, and J. F. Marco, “Oxidation tuning in AlCrN coatings,” Surf. Coat. Technol., vol.

201, no. 8, pp. 4505–4511, 2007.

[75] Y. Feng, L. Zhang, R. Ke, Q. Wan, Z. Wang, and Z. Lu, “Thermal stability and oxidation

behavior of AlTiN, AlCrN and AlCrSiWN coatings,” Int. J. Refract. Met. Hard Mater., vol. 43,

pp. 241–249, 2014.

[76] H. Scheerer and C. Berger, “Wear Mechanisms of (Cr,Al,Y)N PVD Coatings at Elevated

Temperatures,” Plasma Process. Polym., vol. 6, no. S1, pp. S157–S161, 2009.

[77] T. Dudziak, P. K. Datta, P. H. Mayrhofer, and F. Rovere, “High Temperature Oxidation

Resistance of CrAlYN-Coated Ti45Al8Nb,” Oxid. Met., vol. 75, no. 5–6, pp. 359–376, 2011.

[78] M. Moser, P. H. Mayrhofer, and H. Clemens, “Oxidation Resistance and Ductility of a

Coated γ-TiAl Based Alloy,” BHM Berg- Hüttenmänn. Monatshefte, vol. 153, no. 7, pp. 268–

272, 2008.

[79] S. Mato, G. Alcalá, M. Brizuela, R. Escobar Galindo, F. J. Pérez, and J. C. Sánchez-López,

“Long-term high temperature oxidation of CrAl(Y)N coatings in steam atmosphere,”

Corros. Sci., vol. 80, pp. 453–460, 2014.

[80] H.-Y. Chen and F.-H. Lu, “Oxidation behavior of chromium nitride films,” Thin Solid Films,

vol. 515, no. 4, pp. 2179–2184, 2006.

[81] W. Qin, T. Nagase, Y. Umakoshi, and J. A. Szpunar, “Lattice distortion and its effects on

physical properties of nanostructured materials,” J. Phys. Condens. Matter, vol. 19, no. 23, p.

236217, 2007.

[82] S. Ranganathan, R. Divakar, and V. S. Raghunathan, “Interface structures in

nanocrystalline materials,” Scr. Mater., vol. 44, no. 8–9, pp. 1169–1174, 2001.

[83] R. Banerjee, E. A. Sperling, G. B. Thompson, H. L. Fraser, S. Bose, and P. Ayyub, “Lattice

expansion in nanocrystalline niobium thin films,” Appl. Phys. Lett., vol. 82, no. 24, pp. 4250–

4252, 2003.


Content 4.1. Influence of the deposition parameters......................................................................................... 131

4.2. Chemical composition of the thin films ........................................................................................ 132

4.3. Structural analysis of prepared films ............................................................................................ 135

4.4. Morphology of as deposited films ................................................................................................. 139

4.5. Microstructural characterization .................................................................................................... 145

4.6. Phases’ formation and their thermal stability .............................................................................. 149

4.7. Thermal stability characterization ................................................................................................. 160

4.8. Oxidation effect during in situ XRD heating ................................................................................ 163

4.9. Towards a model of crystallization for Zr-Cu TFMGs ............................................................... 165

4.10. Mechanical properties ................................................................................................................... 169

4.11. How to improve durability of surface with Zr-Cu TFMGs? .................................................... 178

Chapter summary ................................................................................................................................... 182

List of references ..................................................................................................................................... 183

Chapter 4

Influence of copper additions on

amorphization of Zr-Cu thin films metallic glasses

Overview

This chapter addresses the influence of copper additions on the thermal structural stability of

Zr-Cu TFMGs. To better understand the metallic glass films behavior, a complex

multidisciplinary investigation has been performed.

It will be presented the relationship between the deposition parameters, the metallurgical

characteristics of the resulting films, in relation to their mechanical and electrochemical

properties.

It is finally demonstrated how their functional properties could propose them as new structural

materials for advanced engineering applications.


Chapter 4

131

4.1. Influence of the deposition parameters

4.1.1. Deposition rate

Generally, when thin films are considered, a prime organization to separate the prepared

samples within distinct groups is made based on the deposition rate. The metallic glass thin

films are produced with “basic” conditions (grown at grounded state and without further

external heating) by varying the deposition time and current intensities on the three metal

targets (see details in chapter 2).

In figure 4.1. is presented the variation of the deposition time as a function of the current

applied to the copper target. Taking into account the experience of the LERMPS laboratory

where coatings were deposited, deposition times were chosen to ensure a targeted films’

thickness of about 4 µm. Current value of the zirconium targets was kept constant at 1.2 A,

except for the lowest zirconium contents for which the intensity was slightly decreased at

about 1.0 A.

Figure 4.1. Evolution of the deposition time of Zr-Cu thin films as a function of current intensity

applied on the Cu target

The figure shows that by changing the current intensities applied on the metal targets and

the deposition periods, samples can be divided into two groups.

In the first zone, corresponding to samples deposited with a low intensity (<0.6 A), films

growth is slow and long deposition process is required (1.5 and 2.5 hours), whereas for high

intensity the expected film’ thickness is reached after about one hour.

A second parameter, namely the deposition rate is deduced from the effective measurement

of films’ thickness by cross-sectional SEM examination and the time used for the deposition

cycle (Table 4.1). From the table we may see the current intensities used for deposition, the

films thickness and calculated deposition rates.

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.60.5

1.0

1.5

2.0

2.5

3.0

Current intensity on Cu target /A

Zone II

Zone I

De

po

sit

ion

tim

e /

h


Influence of copper additions on amorphization of Zr-Cu thin films metallic glasses

132

Table 4.1. Details of the deposition parameters

Current

target

intensity (A)

Film

thickness (μm)

Deposition

rate

μm/h

Current

target

intensity (A)

Film

thickness

(μm)

Deposition

rate μm/h

0.10 6.0 3.0 0.34 5.8 3.6

0.15 4.5 2.0 0.35 4.4 3.0

0.17 4.6 2.2 0.46 4.0 3.1

0.20 6.2 2.5 0.58 4.4 3.2

0.25 5.8 2.4 1.00 5.2 8.0

0.27 5.8 2.4 1.50 4.2 4.3

0.28 5.7 2.3 1.50 3.2 3.2

0.45 5.8 4.0

As expected, deposition rates are lower for the samples in zone I, corresponding to low

current intensities, and increase with the current applied on the copper target.

This evolution suggests that copper contents would be different in the films matrix and

likely different properties and structure will be achieved.

Hence, a well correlation deposition conditions with copper content composition variation is

accomplished.

Regarding the deposition rate, in the first zone (current intensities < 0.6 A) the recorded

deposition rate is almost constant, around 3 µm/h, which denotes a high ability for such

metallic films to be deposited. For the highest copper contents samples (zone II) the

deposition rate augmented almost twice (~6µm/h) and again constant at 3 µm/h for last

sample.

4.2. Chemical composition of the thin films

In our setup, two different metallic targets were used to manage the flux of the two

sputtered metals. Therefore, composition of films can be varied, acting on the relative

current intensities applied to each metal target.

4.2.1. Current intensity on the Cu target

The control of the current intensities is always seen as a simple, but also powerful technique,

giving indications on the process characteristics and thin films growth.

In order to follow the concentration evolution as a function of the Cu target’s intensity

parameter, EDS spectroscopy measurements were carried out. During the deposition of the

Zr-Cu TFMGs two zirconium targets are used to balance the higher sputtering tendency of

Cu with respect to zirconium. The correlation between the chemical composition and

discharge current intensities is presented in figure 4.2.


Chapter 4

133

Figure 4.2. Variation of the atomic concentration of metallic element as a function of the current

intensity applied on Cu target. The oxygen has been detected as well with values around 1 at. %.

From this plot, a first conclusion to be drawn is that a continuous increase of the Cu current

intensity leads to a higher amount of copper introduced into the films’. The copper

enrichment of films can be tailored from about 10 to 90 at.% varying the intensity from 0.1 to

1.5 A. Upon the evolution of the composition, it is possible to discriminate four different

regimes. Zones I and IV exhibit the same type of evolution, namely at low and high current

intensities (<0.3 A, >1.2 A), with films quite poor and very rich in copper, respectively.

Table 4.2. presents the current applied on Cu target and the chemical composition of the

resulting films.

Table 4.2. Synthesis of the Zr-Cu thin film deposition

Sample

number

Current

target

intensity (A)

Cu

content

(at. %)

Sample

number

Current

target

intensity (A)

Cu

content

(at. %)

1 0.10 13.4 9 0.34 48.9

2 0.15 25.0 10 0.35 51.7

3 0.17 28.5 11 0.46 59.1

4 0.20 33.2 12 0.58 65.9

5 0.25 40.1 13 1.00 79.7

6 0.27 41.3 14 1.50 89.1

7 0.28 43.4 15 1.50 98.0

8 0.45 48.3

A comparison between these chemical results and those of both deposition rate and current

intensities shows the possibility to establish some preliminary correlations.

First of all, as shown previously, the deposition process ensures an easy-to-do method to

control the copper content into the matrix only by regulating the current targets’ intensities.

Deposition rates are relatively constant in regime II (intensities below 0.6 A) with an

obtained composition of roughly 50/50 at.%. Most of our films belong to this compositional

region, since results of the literature identifies this 1/1 ratio as the most suitable composition

giving rise to glass-like metals and interesting properties [1 - 3].

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8

0

10

20

30

40

50

60

70

80

90

100

at. % Cu

at. % Zr

IVIIIII

Current intensity on Cu target /A

Co

nc

en

tra

tio

n / a

t.% I



134

In this current intensities range, species are indeed mixed in a more disordered manner,

which could imply better local atomic arrangements and properties.

In order to investigate the whole Zr-Cu domain, some further films were also fabricated

with moderate conditions (regime III).

For the thin films the local ordering and chemistry are the main factors which may impact

the future behavior and orientated-properties [4 - 6]. For instance, a thin film with low free

volume and high packing density state could have higher resistance against crystallization

and enhanced mechanical properties [1, 2]. Thin films prepared in regimes I and IV with a

significant unbalanced Cu/Zr composition is expected to form crystalline nanophase

structures close to the pure metals one.

These initial results made possible to place the obtained compositions on the Zr-Cu phase

diagram to identify the structural domains as illustrated on figure 4.3. For this binary alloy

system the phase diagram is rather complex, with five eutectics and many intermetallic

compounds [7].

Figure 4.3. Zr-Cu thin film’ compositions placed on the Cu-Zr phase diagram [7]. Only the

bottom part of the diagram is presented for a better positioning of our samples. The numbers are

listed in table 4.2.

These initial findings presented mainly the deposition conditions, chemical compositions

and the evolution of thickness as a function of the deposition rate. The deposition

parameters used ensured variation of the chemical compositions as placed on the diagram

phase.

It is important to further determine the structural differences of these films, feature that will

be explored in the following section.


Chapter 4

135

4.3. Structural analysis of prepared films

Structural characterization of films is of prime importance to determine the influence of the

chemical composition on their structural arrangements. During PVD deposition, both the

thermodynamics (working pressure) and kinetics (energy of adatoms) contribute to the

films’ growth. The utilization of magnetron co-sputtering method from individual pure

metal targets ensure the mixing of the species on the film surface and not in the gas phase

[8].

Therefore, the codepositing from individual magnetron sources would imply the synthesis

of films with homogeneous chemical composition. In addition, the rapid deposition process

and the differences in atomic sizes between Zr (0.160 nm) and Cu (0.128 nm) together with

large negative heat of mixing (- 23 kJ/mol) propose the Zr-Cu binary system as an ideal

model for properly analyzing the glass forming ability (GFA) [9].

The complexity of this system should still enhance the probability of amorphization as

showed by the Zr-Cu phase diagram (Figure 4.3).

In figure 4.4. the XRD patterns for the as-deposited films are shown as a function of the Cu

content.

Figure 4.4. X-ray diffraction patterns for as-deposited Zr-Cu thin films deposited on SiO2

substrate

The Zr-richest film (Zr-13 at.% Cu) and Cu-richer film (Zr-98 at.% Cu) are clearly

crystallized and textured along the {111} orientated plans, belonging to fcc-Zr (ICDD card

no. 88-2329) and fcc-Cu (ICDD card 004-0836), respectively.

It is worth noting that fcc-Zr is not the usual/stable expected phase. We interpret its

presence by its nanometer size phase, confirmed by the peaks widths and crystallites

dimension calculated by Rietveld fitting (49 nm for fcc-Zr films and 30 nm for fcc-Cu film).



136

For the Zr-25 at.% Cu and Zr-28 at.% Cu films the structure is also predominantly

amorphous, even if some minor traces of crystal-like phases emerging from the amorphous

hump are detected.

Within the 33.3 – 89.1 at. % Cu range, they consist of only one broad diffraction maxima in

the 2θ range 35 – 45°, characteristic of amorphous structure with no noticeable Bragg

diffraction peaks. Moreover, diffractograms show, as well, that the changes in composition

induced by the variation of deposition parameters (current intensity applied on Cu target,

deposition time), result in different structural local atomic arrangements.

A general analysis of all samples’ evolution indicates a continuous shift of the amorphous

hump towards higher 2θ angles with increasing the Cu content as illustrated in figure 4.5.

Figure 4.5. Evolution of the hump positions and the equivalent lattice spacing deduced from the

interpretation of diffractograms of figure 3.5

The displacement of the amorphous hump can be explained also by the Vegard’s empirical

rule. The law holds that a linear relation exists, at constant temperature, between the crystal

lattice constant of an alloy and the concentrations of the constituent elements [10].

First explanation of this displacement could be related to the local structural organization.

By this means, it can be said that the systematic addition of copper in the matrix modify the

organization, which supposes a continuous movement/reorganization of the atoms at the

local range. The different atomic sizes and the mutual negative heat of mixing between Zr

and Cu (-23 kJ/mol) is expected to promote a random distribution of the atoms [11] and

different local structural arrangements. The local structure organization is believed to be one

of the key factors for this shift.

Secondly, a linear evolution is clearly seen for the equivalent lattice spacing calculated from

the Bragg’s relation considering the hump positions extracted from XRD patterns, which

resembles to the Vegard’s law involved in crystallized alloys. According to this empirical

law, the unit cell parameters should vary linearly with the increase of copper content [10].

Similar linear evolution has already been reported for films in the range of 30 to 70 at.% [9,

10, 12]. The extracted values of d fit quite well the linear progression from the crystallized

pure fcc-Zr (0.260 nm) down to the pure fcc-Cu (0.209 nm).


Chapter 4

137

Hence, shrinkage of the distances between the closest neighbors is established, leading to

more compact structures. The same linear variation has been as well reported by Mattern et

al. [12] in the 30 – 75 at.% Cu regime for Zr-Cu BMGs. Their calculations of the slope gives a

values of 0.69 x 10-4 nm/at.% Cu which is very close to the one obtained for our Zr-Cu

TFMGs (0.60 x 10-4 nm/at.% Cu).

Another possible explanation for this linear evolution may result from a change of the

interatomic interactions’ nature, which could modify the local distribution of atoms as

suggested by Ogata et al. [13]. They reported that for some BMGs the atomic configuration

tends to have directional bonding generated with the surrounding atoms. Due to their

highly negative heat of mixing, both Zr and Cu atoms have a strong tendency to attract each

other and most likely tight bond pairs should have been formed [8, 14]. These bonding’s

architecture may locally stress the atomic structure, creating defects. The strong mutual

affinity favors the formation of continuous solid solution (any demixion phenomena) as

required for the Vegard’s law application. This is the reason why Park et al. classified the Zr-

Cu binary system as a metal-metal type system with a topological short range ordering [1].

The same linear progression has already been reported in other studies, where it was

explained through theoretical calculations based on simulations [12, 13, 15].

Authors stated that it may result from a change of the interatomic interactions’ nature which

could modify the local distribution of atoms. It has been reported that for some BMGs the

atomic configuration tends to have directional bonding generated with the surrounding

atoms.

On the other hand, diffractograms from figure 4.6. can also be interpreted regarding the

mean dimension of diffracting domains, considering the full width at half maximum

intensity parameter (FWHM) and applying the Scherrer’s formula.

In figure 4.6. the evolution of the FWHM parameter and diffracting domains size for the

compositions of amorphous films (33.3 – 89.1 at.% Cu domain) are summarized.

Figure 4.6. Evolution of the full width at half maximum parameter, and the corresponding

equivalent diffracting domains’ size for the amorphous compositions as a function of the

chemical composition



138

The progression of FWHM may give further insights about the continuous modification of

the distances between the first atoms’ neighbors. Values are passing through a maximum at

about 6.5° (1 nm-sized diffracted domain) with a bell-shaped dependence versus the

composition fitted with a polynomial law. Such result is independent on the fitting’s nature

adopted by the Voigt function or the Scherrer’s formula.

From the interpretation of the figure, three different regions can be distinguished.

The central part, ranging roughly from 45 to 78 at.% Cu where we found a plateau,

corresponds to the regime B, while films below 45 and above 75 at.% Cu define regimes A

and C, respectively. The trend denotes that by varying the copper contents in the matrix the

local structural ordering is affected, illustrated by different types of arrangement. The

smallest organized domains are associated to the wide central compositional region.

At the same time, when the quantities of available atoms are almost equal the mixing takes

place in a more complex manner, ensuring a higher disorder in the matrix. The samples

located in regime B reveal higher amorphization state compared to the other Cu-poorer or

richer thin films. Variation of the FWHM as a function of the copper enrichment displays a

bell-shape curve characterized by a maximum at 60 at.% Cu and ~6.5°.

Mattern et al. [12] found almost the same variation for which their bell-like curve showed a

maximum centered at 55 at.% Cu and ~6.6°.

Moreover, it has to be considered that X-ray diffraction is a bulk probe technique and the

whole illuminated sample volume contributes to the scattering. Therefore, the diffracting

domains contain information on the average structure in the glass, even if, locally organized

structures could exist. It is often argued that metallic glasses-based materials have a strong

icosahedral ordering [1, 13, 15 - 17] which is incompatible with any crystalline structure

However, owing to the small dimension of such nuclei, available information mainly result

only from simulations calculations.

With the results from the structural investigation, it can be said that the deposition

parameters (i.e., deposition rate, thickness) highlighted in the previous section are related

with different chemical compositions and structural states.

Considering these last results, a question arises related to the films morphologies and the

covering rate of the substrates, feature that will be explored in the next section.


Chapter 4

139

4.4. Morphology of as-deposited films

During sputtering of species the increased nucleation density and the mobility of adatoms

on the surface under bombardment may be imperative in disrupting the developed

microstructure/morphology. The film density is of great importance in thin films properties

such as: better corrosion resistance, lower electrical resistivity, higher hardness or increased

index of refraction. Similarly, most of vapor-deposition being non-equilibrium methods,

these processes influence the films’ microstructure and morphology [18, 19].

A microstructural analysis of surfaces was carried out to further examine the influence of

the deposition parameters in both plan and cross-section views:

Top-surface morphology

In figure 4.7. SEM micrographs of some selected thin films belonging to the four zones of

figure 4.2. The SEM micrographs reveal that as-deposited films exhibit a dense and smooth

morphology, whatever their copper contents. Such type of homogeneous surface

morphology is typically expected with the magnetron sputtering process.

SEM plan view

Figure 4.7. General top surface reveled by SEM for selected representative Zrx-Cu1-x thin films

deposited on Si substrate: 13.4 at.% Cu (a), 40.1 at.% Cu (b), 59.0 at.% Cu (b), 89.1 at.% Cu (c)



140

From the figure it can be seen that only small micrometer-sized defects are detected on

films’ surface. The surface topography quality is more linked to the deposition method than

to the composition, that’s why only some representative samples were selected.

Regarding coating’s growth there have been reported three well-known types of

morphologies referred to as: island growth (Volmer-Weber’s model), layer by layer growth,

or their combinations (Stanski–Krastanov’s model) [20].

The growth type adopted depends on the way of how the atoms are attracted to the

substrate. The Zr-Cu TFMGs would be deposited according to a Stanski–Krastanov’s model.

In their architecture mechanism, the initial layer growing on the substrate tries to cover the

substrate as closely as possible, straining itself until the strain energy increases as the layer

thickness increases [20]. In the final stage the surface become preferable energetically and

some island like-defects are formed on the films’ surface. Some examples that show this

type of growth are the metal films and semiconductor films [21]. Considering these aspects

and since the Zr-Cu thin films are regarded as metallic glass-like, the films are most likely

growing in a comparable mode.

In order to confirm such assumption, the surface needs to be examined at a deeper scale

using high magnification SEM micrographs and AFM observations as presented in figure

4.7. a and b. Micrographs show a very fine gathered grained morphology, leading to a quite

coarse substructure of some hundreds nanometers.

at.%

Cu SEM plan view AFM plan view RMS

13.4

7.13

nm

40.1

2.51

nm

Figure 4.7.a. Fine top surface examination by SEM and AFM for Zr-13.4 at.% Cu and Zr-40.1 thin

films deposited on Si substrate


Chapter 4

141

From this figure it is possible to observe that the films grow with a compactly grainy-like

microstructure. Surface roughness of films is crucial for die-forming, thermoplastic forming

or for other applications such as MEMS and nano-devices [22, 23].

at.%

Cu SEM plan view AFM plan view RMS

59.1

2.60

nm

89.1

2.80

nm

Figure 4.7.b. Fine top surface examination by SEM and AFM for Zr-59.1 at.% Cu and Zr-89.1 at.%

thin films deposited on Si substrate

The investigation of the surface topography reveals that whatever the copper content in the

matrix, films always exhibit fine and smooth surface topographies. The typical micro-

granular morphology with homogeneous grains varying from about 3 nm to 7 nm, the latter

value corresponding to the crystalline Zr-13.4 at.%.

Surface’s morphology characteristics are independent of the chemical composition, as

reflected by the preserved grainy-like and disordered structure.

It is therefore clear that the surface roughness of the as-deposited films is strongly related to

the sputtering parameters [24, 25].



142

Cross-sectional examination

In order to further study the influence of the deposition parameters on the films’ growth, the

obtained thin films were examined in fractured cross-section.

Figure 4.8. displays further SEM micrographs of the morphologies of fractured coatings

deposited onto silicon, underlying the effect of Cu addition on the films’ microstructure.

The cross-section analysis is an important result for the thin films, which has to be known

since it gives initially the thickness of the as-deposited films and then significant details on

the growth type or adhesion.

The PVD magnetron sputtering technique could give rise to different morphologies and

films growth. In the figure below only the most representative compositions are selected to

be shown, since the films’ morphology is quite similar among them.

SEM cross-section view

Figure 4.8. SEM cross-sections micrographs of Zr-Cu thin films containing: 13.4 at.% Cu (a), 40.1

at.% Cu (b), c. 48.3 at.% Cu (c), 59.1 at.% Cu (d), 65.9 at.% Cu (e) and 89.1 at.% Cu (f)


Chapter 4

143

The first remark to be highlighted is that all films have a relatively homogeneous thickness.

The film containing 13.4 at. % Cu reveals a characteristic fine columnar-like growth (Figure

4.8.a.) linked to its crystalline structure. Among the selected compositions presented in this

section, TFMGs develop shear vein striations in river-like array.

For the films deposited with a current below 0.6 A and belonging to compositional regions I

and II (Figure 4.2.) the SEM cross-section micrographs shows partial vein-like patterns.

These features are randomly distributed through the entire film’s thickness with different

orientations.

For Zr-40.1 at.% Cu TFMG (Figure 4.8.b.), the structures descend from the top film’s surface

down to the substrate, whereas for 48.3 at.% Cu TFMG (Figure 4.8.c.) the vein-like patterns

seem to arise from the substrate to the top surface.

Addition of more Cu into the matrix appears to impact the formation of these features and

to reduce their size as it is the case for the Zr-59.1 at.% Cu TFMG. Vein patterns morphology

is changing from the third compositional regime (Regime C, Figure 4.6.), where two types of

features, primary and secondary, can be distinguished.

The primary features are thicker and somehow parallel to the substrate direction, while the

second population of features is finer and grouped in between the main bigger former ones.

The changes in dimension of these features attest that from this composition (59 at.% Cu) a

gradual transition from the vein-like features to finer and probably featureless

characteristics is promoted.

Actually this is confirmed by the Zr-89.1 at.% Cu TFMG from regime IV for which

apparently a featureless morphology is detected on the cross-section, as illustrated in Figure

4.8.f. Nevertheless, adopting a deeper examination allowed us to realize that the apparently

featureless area was in fact also composed of a very fine strain striations network, as shown

in figure 4.9.

Figure 4.9. Nano-scale striation patterns (d) and very fine superficial vein-like network on the

fractured surface of the Zr-89.1 at.% Cu thin film

From figure 4.9. it can be observed many fine nanoscale waves separated by a less than 30

nm “wavelength”. On the top part there is a thin outer layer with superficial different

morphology and slightly more pronounced vein striations.



144

The origin of these patterns with local traces of shear vein striations could be explained by

specific concentration sites where the plastic flow could be absorbed preferentially during

mechanical deformation (i.e., fracture).

It is remarkable stating that the films’ morphology is changing from fine columnar-type for

lower copper containing thin films (related to its crystalline state) to vein and apparently

featureless-like structure for higher containing films. This general trend again could indicate

a composition effect associated to the plastic deformation mechanism.

The particular vein morphologies often reported on BMGs, are a further clue to identify the

metallic glass-like character of our Zr-Cu thin films. Such type of structures were especially

reported for instance by Nagendra et al. for La-based bulk metallic glass [26], P. Murali and

U. Ramamurty for Zr-based BMG (Vitreloy-1) [27] or by P.J. Hsieh et al. for Mg-based BMG

[28].

In those cases, vein-like associated morphologies may reflect differences in the intrinsic

mechanical behavior of the layer. Indeed, in the II and III regimes (Figure 4.2.) the fracture

characteristics are well-ordered in dense and smooth networks throughout the whole film’s

thickness.

The fine aspect of those patterns infers that beside its metallic glass-like character, the

plasticity of Zr-Cu TFMGs is a composition depended parameter. When the fracture surface

morphology exhibits a characteristic river pattern such as the one shown in figure 4.8.f. or

4.8.a., large plastic strains would be involved due to inhomogeneous material’s flow,

leading to a quite ductile behavior of the material.

In the literature, the presence of smooth and fine river-like features for metallic glass-like

materials is associated to localize plastic flow, which occurs in the shear bands indicating

low fracture toughness [29 - 31]. The initiation of periodic striations is supposed to be the

result of self-assembled cavities when their size reaches the critical value influenced by the

elastic wave. This could be an indication of activation of plastic flow and the deformation in

the compression process [29].

The nature of the structural features that control the plastic zone size, as well as the

connection between fracture and fundamental deformation processes, remain still to be

clarified by future research.

In addition, during a typical fracture, the formation of such features would reduce the stress

concentration and affect further the crack propagation by increasing the fracture toughness

[32, 33]. For BMGs materials, the formation of shear bands during the plastic deformation is

normally well-known and numerous models are available [27, 31, 34 - 37].

Some authors propose a deformation map for metallic glasses with focus on the

homogeneous and inhomogeneous flows, similar to the classical deformation mechanism

maps for polycrystalline materials [38].

The formation and localization of shear bands is frequently recognized as a direct

consequence of strain softening. The increment of the strain applied to a local volume

element softens that element allowing a continuous localized deformation even at high rates

[31]. For the metallic glasses, some potential reasons for strain softening and localization

includes the local production of free volume due to flow dilatation, structural local

evolution and heat generation. The dominant factor is thus believed to be a local change in

the state of the glass.


Chapter 4

145

In what concerns the TFMGs, only few reports [24, 25, 39, 40] are talking over the SEM

fractured morphologies phenomena in the light of deposition parameters and composition

effect. What can be pointed out is that SEM cross-section micrographs of Zr-Cu TFMGs

(Figures 4.12. - 4.13.), clearly illustrate the shear striations related to the fracture behavior

(low toughness) and metallic glass character.

Besides, results obtained in the Cu-rich side indicate that shear steps are separated into

many smaller shear surfaces and areas with viscous fracture type and vein-like patterns.

The spacing of these smaller shear planes have a size of several tens of nanometers. Intrinsic

fracture mechanical behavior of films can be influenced by its deposition step.

The vein-like features observed on SEM fractured micrographs are likely related to the

classical ductile fracture mechanism involved for the bulk metallic glasses. We can also

affirm that Cu enrichment of film promotes a reduction of the striation features, which leads

to smoother morphologies associated perhaps to a more brittle character.

These so-called general characterizations when dealing with thin films, such as

understanding the role of deposition parameters, crystallographic arrangements, or

thickness measurements allowed us attained a general idea of the as-deposited Zr-Cu thin

films.

The next section will be devoted to a microstructural investigation by TEM that would give

further insights on the structural differences between the samples.

4.5. Microstructural characterization

Results presented in the previous sections provided a general overview on the prepared Zr-

Cu thin films in terms of deposition conditions, surface morphology and structure.

When dealing with metallic glass-like thin films, it is important to verify at which scale the

material is indeed amorphous. In this section, the results of the microstructural

characterization are therefore discussed. The deep scale analysis is performed on several

compositions in order to explain the influence of the deposition parameters and copper

enrichments into the amorphous matrix. The selected compositions (40.1, 48.3, 59.1, 89.1 at.

%) will allow understanding the microstructure evolution as a function of copper content in

the matrix.

Figure 4.10 – 4.13. illustrate examples of nominal high-resolution TEM (HRTEM)

micrographs for the selected representative compositions. In general, the structure of all

deposited thin films consists mainly in a dominant fully amorphous matrix. For certain

compositions, some nanodomains in the diffuse matrix have been observed which resemble

to the behavior reported for Kondoh et al. [41]. The fine nanocrystallised phases, depending

on the composition, result in some Zr-Cu intermetallic compounds or pure copper.

For example, in the case of the Zr-40.1 at. % Cu film, the overall structure can be considered

as amorphous (consistent with XRD results), since only some weak traces of lattice planes as

shown in figure 4.10 are extracted from the Fourier transform (FT) pattern (indexed in the

figure with 1, 2, 3, 4).



146

Figure 4.10. HRTEM plane view micrographs of Zr–40.1 at.% Cu thin film (a) and Fourier

transform (FT) pattern. Spots are indexed with positions 1 – 4

The nanometer-sized crystals generate visible spots (indexed with positions 1-4) into the

Fourier transform pattern of corresponding micrograph. The extracted corresponding d

spacing value found (2.12 Å), could be attributed to Cu10Zr7 (600) and/or Cu51Zr14 (410)

intermetallic compounds, or to Cu (111).

Regarding the Cu-richer compositions (Zr-59.1 at. % Cu) the films have a different and more

complex microstructural nature. In figure 4.11.a. is highlighted the presence of only one

nanoparticle, while the Fourier transform pattern (Figure 4.11.a1.) reveals the presence of

two potential compounds.

Figure 4.11. HRTEM plane view micrographs of Zr–59.1 at.% Cu thin film with an inserted EDS

spectrum on 2 nm highlighted circle (a) and indexed Fourier transform pattern acquired on

region delimited by dotted line square of previous micrograph; non-underlined numbers are

corresponding to Cu indexation ([1 –10] zone axis), while underlined numbers are attributed to

the Cu10Zr7 ([− 742] zone axis (a1))


Chapter 4

147

The crystals from figure 4.11. are in a crystallographic relationship with a common

constitutive (111) Cu plan parallel to (215) of Cu10Zr7 phase. The indexation, based on dhkl

list and angles between spots, allowed us to identify a pure Cu nanocrystal in [1-10] zone

axis and Cu10Zr7 intermetallic nanocrystal in [-742] zone axis, respectively.

Besides, this result could suggest that a demixion phase could be promoted during the

nucleation/growth period. Further EDS analysis might confirm this assumption when

different electron probe’s sizes are used. For a 2 nm diameter beam localized on the

nanoparticle part (highlighted by the disk in figure 4.11.a.) the chemical composition

obtained is close to pure Cu (Zr4-Cu96 at. %), whereas with a coarser beam (80 nm

diameter) the obtained composition is slightly different as compared to the chemical

composition of film (Zr36-Cu64 at. %).

Concerning the samples with higher Cu contents, the HRTEM for the selected Zr-89.1.0 at.%

Cu film evidences formation of pure Cu phase as illustrated in figure 4.12.

Figure 4.12. HRTEM plane view micrographs of Zr–89.1 at.% Cu thin film (a) and the indexed

Fourier transform pattern (a1). The nanoparticle is composed of two twinned pure copper

crystals

From the indexation of the crystal using the Fourier transform, a pure copper phase is found

to be embedded in the amorphous matrix, which is in good agreement with the phase

detected by XRD at a global scale.

The lattice distance associated to the ring (2.12 Å) is at the same position as the hump

detected by XRD, and could be attributed to the most important {111} diffraction line of

copper (2.087 Å).

These TEM investigations gave us indications on how the fine structure of thin films may

evolve with their different chemical compositions.

Finally, we have furthermore concentrated the TEM investigations on the 48.3 at.% Cu

containing film from regime B (Figure 4.6.) since this compositional region is intended to

have the most interesting properties as mentioned in literature.

HRTEM image corresponding to a general zone and the TEM diffraction associated is

presented in next figure.



148

Figure 4.13. HRTEM plan view micrograph of Zr-48.3 at.% film with an inserted SAED (analysed

zone 400 nm in diameter)

For this sample no lattice plane traces were observed, which was also confirmed by Fourier

transform and SAED where no spots are seen. The diffuse superimposed halo suggests the

fully amorphous matrix. Besides, width of the diffuse ring is quite consistent with the

amorphous hump detected in XRD and confirmed that, whatever the adopted investigation

scale, the film containing 48.3 at. % Cu is fully amorphous by nature.

Despite intensive study performed on this sample, no hint of any crystallization is observed

in HRTEM imaging, Fourier Transform pattern or selected area diffraction.

Lastly, some preliminary conclusions can be drawn from this section. Although most of

films were identified as amorphous by X-ray diffraction, a deeper investigation at finer scale

by TEM reveals that they may be composed of nanometer-size crystallized domains.

It is worth mentioning that when comparing the amorphous vs. crystalline domain, for all

the compositions studied by TEM the overall structure was constantly amorphous with

some minor crystal-like domains embedded on different local zones.

Hence, the Zr-Cu thin films can be considered as fully XRD-amorphous at a global scale and

in the best case nanocrystalline-type (for some compositions) when studied at a local level

by TEM. The film containing 48.3 at.% Cu appeared to be fully amorphous whatever the

considered investigation scale.

In this context, their amorphous structure will tend to crystallize when heated, as will be

showed in the following section.


Chapter 4

149

4.6. Phases’ formation and their thermal stability

In this part, the focus is on the effect of Cu additions on the phase formation and thermal

stability, as well as on the thermal parameters of Zr-Cu TFMGs.

The way how metallic glass-like thin films may evolve with the temperature is thus

discussed and one example of the crystallization behavior from the three regimes mentioned

in figure 4.6. will be explained in details.

In this section for a simplification purpose, we will refer to the temperatures in K and not in

°C, as referred in the hard coatings chapter, since this is more common for the metallic glass

community.

4.6.1. Thermal structural evolution of the glassy structure

The ability of the PVD co-sputtering process to form Zr-Cu amorphous thin films has been

investigated in a quite wide compositional range as presented in figure 4.7.

As highlighted by many studies the crystallization of metallic glasses is an interface-

controlled three-dimensional isotropic growth process with an early-stage saturation of

nucleation.

Most studies demonstrate that these materials crystallize through a multistage manner [42 –

44] influenced, especially, by the short and long range diffusion leading to phase

separations [45 - 47].

Considering the characteristics of the hump interpreted in the figures 4.5. and 4.6. it seems

that films of regime B could present an enhanced behavior in terms of resistance against

crystallization and promising properties (i.e., mechanical).

In this regard, the crystallization behavior and structural stability were experimentally

pursued by means of HT-XRD in situ experiments, carried out using a continuous heating

from the room temperature up to 873 K.

Many earlier studies on Zr-Cu alloys [9, 14, 42, 48] have inferred that various properties of

amorphous alloys are directly related to their structural stability, evaluated by their

crystallization temperature. In fact the amorphous alloys having low crystallization

temperatures have inferior structural stability, and therefore structural changes can easily

undergo when an external energy is applied (i.e. thermal, mechanical).

The findings from previous reported studies on Zr-Cu system in a bulk or thin film form [2,

11, 41, 49 - 52] clearly showed the composition dependence of this system with similar

tendencies for crystallization [2], microhardness [52], reflectivity and sheet resistance [49].

Among them, the study of K. Park and coworkers on the atomic packing density and its

influence on the properties of Cu-Zr amorphous alloys is of prime importance. For their

investigated compositions the maximum level of atomic packing is situated between 55 and

75 at.% Cu domain which corresponds well to the regime B for Zr-Cu thin films

investigated.



150

In our case, it is noteworthy to mention that an originality of our work relies on the

important number of investigated samples, covering the whole Zr-Cu diagram and giving a

sound idea of the characteristic of the film, as a function of its chemical composition.

As a result, the intermetallic phases’ transition was perceived, from crystalline fcc-Zr, to

some intermetallic phases, like CuZr2, Zr10Cu7 or Zr51Cu14 to finally pure fcc-Cu for the

highest copper contents.

In figures 4.14., 4.16., and 4.19. will be illustrated the structural thermal evolution of glassy

metallic films belonging to each selected compositions from the three regimes. The top part

of figures presents a 3D-representation (axes: x=Bragg’s angle, y=temperature, z=intensity of

diffractograms, while on the bottom part is displayed some selected diffractograms

emphasizing the formation of metallic and intermetallic phases.

As a general behavior, the disordered structure of Zr-Cu TFMGs is stable up to about ~573

K. The well-known affinity of Zr for oxygen combined with long annealing times (~7 h)

favored the formation of ZrO2 oxides at early stages of heating (ΔG = -1037 kJ/mol at

standard conditions [53]). These oxides formed on the films’ surface are due to the poor

vacuum in the XRD oven (~ 102 Pa) leading to a loss of Zr atoms available in the structure.

The initial formed oxide is the tetragonal ZrO2 (ICDD card no. 50-1089), followed by the

monoclinic variety (ICDD card no. 37-1484).

It can be assumed that the monoclinic phase would be present on the top surface, while the

more surprising tetragonal ZrO2 would result from an internal oxidation (or oxygen inward

diffusion), presenting a nanosized grains submitted to high internal stresses.

When increasing the temperature within the supercooled region, no sharp peaks were found

and the broad amorphous hump remained unchanged. Depending on the initial

compositions, the structure is generally changing its state above 573 K when some

intermetallic compounds precipitate from the amorphous matrix.

We will therefore begin to present the results concerning the structural stability as a

function of temperature for each sample, beginning with the sample having lower copper

contents.

Evolution of the structure in regime A (Cu content < 45 at.% Cu)

Considering previous XRD results (Figure 4.4.) the as-deposited films containing 25.0 and

28.5 at.% Cu presented some hints of crystallization emerging from the hump, whereas from

a 33.3 at.% Cu content, the coating was fully XRD-amorphous.

This composition limit is consistent with recent results of Coddet et al. [3], suggesting a

threshold crystallization value of about 29 at.% Cu.

For the low Cu contents domain, some Zr-based nuclei are likely to be dispersed into the

amorphous matrix during the deposition. As a consequence, the fcc-Zr nanophase detected

for several compositions (25, 28.5 at. % Cu) from this regime does not correspond to the

usual expected stable phase in this alloy.

We suppose that during deposition a rather metastable nano-sized phase is formed. By

increasing the copper contents some amounts of the fcc-Zr nano phase are perhaps still

present in the matrix, but not in an enough quantity to grow crystal-like phases.


Chapter 4

151

The presence of this nanoscale precipitates (nuclei) could be responsible for the initial

crystallization, having most likely a role of germination-activated sites. Hence from this

regime we have chosen to present the thermal evolution of the Zr-40.1 at.% Cu film as


Figure 4.14. XRD patterns describing the thermal structural evolution of the Zr-40.1 at.% Cu

TFMG obtained using the in situ high temperature mode, a: global 3D view, b: summary of the

main structural transformations

We may see a 3D representation of the structural evolution for Zr – 40.1 at.% Cu film, while

in part b only the main structural events are shown. Using a more detailed characterization,

it can be seen that the film remains X-ray amorphous up to 603 K, where some minor traces

of body-centered tetragonal CuZr2 are detected. This bct-CuZr2 compound agrees to the

thermodynamically-expected phase and is the first phase to be formed.

Later on at 633 K, the metallic glass-like film is completely crystallized into this latter phase

(ICDD card no. 65-7783). At 693 K a new fcc-Zr2Cu phase appears (ICDD card no. 74-7476),

corresponding to an ordered arrangement of the tetragonal form and characterized by a cell

volume 16 times bigger. During annealing, initial nano-grains of bct-CuZr2 phase coarsened

accompanied by the further precipitation of fcc-Zr2Cu. Similar crystallization sequence has

already been reported by Nagase and Umakoshi [54] by an electronic irradiation of a

Zr66.7Cu33.3 metallic glass.



152

The final crystallization stage is reached when some traces of Cu10Zr7 phase (ICDD card no.

42-1187) are measured above 843 K.

In addition to the qualitative approach aimed to determine the thermal evolution of phases,

a more quantitative evaluation of the relative amount of each amorphous and crystallized

phase has been determined through the CaINne© software using the atoms’ positions

resulted from the Inorganic Crystal Structure Database (ICSD) for the studied structures

(CuZr2-ICSD 151846, Cu10Zr7-ICSD 164881, Cu51Zr14-ICSD 629471, Cu-ICSD 52256).

The proportion of amorphous matrix is estimated by the ratio of the amorphous hump’s

area over the one of the whole diffractograms using the TOPAS structural analysis software

as displayed in table 4.18.

These data allowed us to deduce the theoretical XRD profiles and to calculate the relative

volume ratios using the formula:

Cjv =

1

1+∑ViVj

∗xixj

j

(i ≠ j) [Eq. 4.1.]

where, Vi is the cell volume of phase i and xi is the cell number fraction of each phase

obtained from simulations.

Figure 4.15. presents a comparison between the simulated and experimental diffractograms,

which appears to be a good correlation.

Figure 4.15. Example of comparison between the simulated and experimental XRD pattern of Zr-

40.1. at% Cu obtained at 813 K during in situ XRD experiment


Chapter 4

153

From the simulated XRD pattern it is possible to deduce the relative amount of each phase

as is shown in table 4.3.

Table 4.3. Phases’ ratios and relative amount of amorphous phase in the matrix for the Zr-40.1 at.

% Cu TFMG

Zr-40.1 at. % Cu TFMG

Temperature

(K)

Phase (volume ratio %) Degree of

amorphous (%) bct-CuZr2 fcc-Zr2Cu Cu10Zr7

573 - - - 100

603 100 0 0 85

693 97 3 0 0

723 90 10 0 0

813 85 15 0 0

873 65 22 13 0

The same evolution presented in figure 4.14. it can be observed as well here. First remark is

that at 573 K the film is 100 % fully amorphous, as detected by XRD.

From 603 K the degree of amorphization decrease at 85 % since some traces of bct-CuZr2 are

recorded into XRD pattern. These traces represent 100 % of bct-CuZr2 in terms of phase

volume ratio.

From 693 K, due to already initiation of crystallization the second fcc-Zr2Cu is developed

into the matrix with a ratio of 3 %.

The further increase of the temperature modifies the phase’s volume rations between the

two intermetallic compounds and finally at 873 K, the third Cu10Zr7 represents 13 % from

the phase volume ratio of the matrix.

The crystallization behavior for this Zr-40.1 at. % film as observed from the in situ HT XRD

patterns (Figure 4.14.) is a multistage process with the formation of unstable initial phase

that further endure changes.



154

Evolution of the structure in regime B (45 at.% Cu < content < 75 at.% Cu)

In the second regime (represented by the Zr-59.1 at. % Cu, Figure 4.16.), the crystallization

temperatures are higher (663 K) compared with the other compositions, suggesting a better

structural stability and denser atomic packing state.

TFMGs from this regime endure similar structural transformations, with the amorphous

matrix transforming into different intermetallic compounds as shown in the next figure.




From the figure, it can be seen that first small peaks of both Cu10Zr7 and Cu51Zr14

intermetallics appeared from 663 K.

Based on the phase diagram, the most probable phase should have been the Cu10Zr7 phase,

which has priority during crystallization since it is more thermodynamically favored. At one

hundred degree higher, the thermal energy becomes high enough to allow the precipitation

of the second intermetallic compound: Cu51Zr14.

Within the [663-783 K] temperature range, there is a competition between the continuous

crystallization-coarsening of Cu10Zr7, and the crystallization of Cu51Zr14. The final stable

crystallization phases in the regime B upon heating at 873 K were both Cu10Zr7 and Cu51Zr14

intermetallic compounds. Similar multi-stage process of crystallization was confirmed

previously by the results proposed by Cui et al. for a similar Cu64Zr36 alloy’s composition

[42].


Chapter 4

155

The same estimation as previously showed related to the volume ratio for Zr-59.1 at.% Cu

TFMG is presented in table 4.4.


% Cu TFMG

Zr-59.1 at. %Cu TFMG

Temperature

(K)

Phase volume ratio (%) Degree of

amorphous (%) Cu10Zr7 Cu51Zr14

633 - - 100

663 100 0 76

783 37 63 0

843 12 88 0

873 4 96 0

The results presented in the table shows that at 663 K the Cu10Zr7 intermetallic compound

occupies 24% of the diffracted crystallized volume as single phase. This means that the

matrix is still mainly amorphous, since the degree of amorphous ratio is quite high (76 %).

From 783 K, the second intermetallic phase, namely Cu51Zr14 have grown and the volume of

the phase’s ratio is 88 % and 12 % for the Cu10Zr7 initial phase.

Interesting is that both phases are spotted in the XRD pattern in the fully crystallized matrix

(degree of amorphous state is at 0 %) as rather mixed compound. At the final temperature

(873 K), the initial Cu10Zr7 intermetallic phase is detected only as 4 % of the volume, while

the Cu51Zr14 represents 96 % of the total phase volume.

This behavior suggests that the initial Cu10Zr7 detected phase serve as nucleation site for the

formation of the more stable Cu51Zr14 phase.

On the other hand we wanted to investigate further the microstructure of this film at a

temperature close to 633 K, where the XRD showed 100 % amorphous structure.

Hence, we performed ex-situ annealing under vacuum for 30 minutes and 1 h at a

temperature equivalent to the in situ XRD (above 610 K).

The HRTEM images showing the microstructure after annealing are presented in figure 4.17.

Figure 4.17. HRTEM on Zr-59 at.% Cu TFMG after annealing above 610 K for 30 minutes (a) and

1 hour (b)



156

The findings imply that germination takes place with an increasing speed as a function of

time. By means of TEM we were capable to observe the nanocrystals with a size below 20

nm, which due the low volume fraction and their small size in XRD pattern seems to be

similar of that of 100 % amorphous structure.

Moreover, if we consider further the SAED pattern showed in figure 4.18., performed on the

sample annealed for 1h, we may conclude that the simulated XRD using CaRIne software

confirm that these nanocrystals are the ones who will form the Cu10Zr7 phase at 663 K, as

confirmed by the in situ XRD patterns (Figure 4.16.) and volume ratio phase fractions (Table

4.4.).

Figure 4.18. SAED TEM of the Zr-59.1 at.% Cu TFMG annealed above 610 K for 1h and the insert

simulated XRD profile corresponding to the Cu10Zr7 phase

Evolution of the structure in regime C (Cu content > 70 at.% Cu)

In this regime corresponding to the richest-Cu thin films presented in figure 4.19., pure

copper first crystallizes during annealing, even if the most probable intermetallic phase for

this composition domain should be Cu51Zr14.


Chapter 4

157




From this graph is seen that unlike the other Zr-rich films, no traces of zirconium oxides are

seen in the XRD patterns below 543 K. Starting from 603 K, crystallization process is

somewhat accelerated and signs of Cu51Zr14 phase were detected.

The multistage crystallization process is confirmed as well for this regime for which an early

devitrification with the formation of fcc-Cu is followed by a gradual decomposition into

Cu51Zr14 intermetallic.

From 753 K, film is fully crystallized with a further surface oxidation of zirconium.

Therefore, at the film top-surface, Zr is preferentially consumed by oxygen, inducing Zr

depletion (Cu enrichment) on subsurface, hindering the intermetallic compound

precipitation. A further heating up to 873 K makes the fcc-Cu and Cu51Zr14 the dominant

metallic phases recorded from XRD patterns.

The local structural arrangement is strongly transformed compared to the films of the other

regimes. Since the copper concentration is very high in the matrix, other types of metallic

bonds should be present. Cu and Zr atoms interact differently, changing the atomic packing

density and forming preferentially Cu-Cu bonding’s type. These main networks rule the

local structure and may manage the crystallization process [8, 15] as proven by molecular

dynamics simulations [1, 13, 15 - 17].

Their conclusions state that, for high copper contents, some preferential atomic Cu

icosahedral-type clusters may be distributed in the amorphous matrix. During the early

stage of annealing, these Cu clusters may act as germination sites for the growing of Cu-rich

phases. The high amount of Cu-rich species induces a change of the local atoms’ ordering,

susceptible to influence the film’s structural stability. Such an assumption could be

confirmed by the presence of fcc-Cu phase (ICDD card no. 70-3038) for the Zr-79.7 at.% Cu

TFMG from 513 K (see peaks at 43 and 50° in Figure 4.19).



158

The evolution of the phase volume ratio for this composition with respect of degree of

amorphous is highlighted in following table.


% Cu TFMG

Zr-79.7 at. %Cu TFMG

Temperature

(K)

Phase volume ratio (%) Degree of

amorphous (%) Cu Cu51Zr14

483 - - 100

513 100 0 74

603 20 80 68

753 30 70 0

873 55 45 0

The relative degree of amorphization coefficient decreases at 74 % once the pure copper

begin to crystallize at 513 K, occupying 100% of the diffracted crystallized domain.

From 603 K where the remaining amorphous structure is at 68%, the second Cu51Zr14 phase

is formed which occupies 80% of the total crystallized domain. At the final step (873 K) the

crystallized phases (Cu and Cu51Zr14) are almost 1/1 ratio in the matrix.

Based on the volume ratios reported in tables 4.3. – 4.5., the structural evolution of the

selected films was as well explained by means of theoretical models.

It is thus clearly confirmed that Zr-Cu thin film metallic glasses are governed by a multi-

stage crystallization process controlled by the initial compositions, kinetic and

thermodynamic factors.

In many of the primary studies it has been indirectly suggested that various mechanical and

physical properties of amorphous alloys are related to their structural stability [42, 45, 48, 56,

57] the degrees of which can be assessed by their crystallization temperature.

The thermal evolution for the selected compositions presented above permitted the

structural and crystallization differences during heating with an example from each

compositional regime (Figure 4.6.).

Furthermore, figure 4.20. gathers all the transformation temperatures deduced from the in

situ XRD patterns (only metallic phases are reported). The evolution again demonstrates

that crystallization is a dependent concentration process influenced by the atomic packing

density and local structural ordering.


Chapter 4

159

The same three regimes as those presented in figure 4.6. for FWHM’s parameter can be as

well distinguished once more.

Figure 4.20. Crystallization temperatures and crystallized phases deduced from in situ HT-XRD

These findings indicate that crystallization temperature gradually increases by increasing

the Cu content until it progresses to a maximum in the regime B. The obtained values

confirmed the general tendency reported in the earlier studies, validating a maximal

structural stability at around 55 at.% Cu. The evolution is comparable to the results from

other studies on Zr-Cu alloys [1, 2, 42, 48]. It is remarkable that the obtained results point

out that the region where the stability is at its highest level would be strongly enlarged as

compared with [1] to corresponding [45, 75] at.% Cu chemical composition range.

The crystallization values are overall higher than for the other two regimes indicating a

more effective stabilization of the supercooled liquid against crystallization. This evolution

is the same as the one already presented in figure 4.6.

The FWHM parameter from the XRD as-deposited patterns is therefore a reliable parameter

that could give direct information about the local structural state and final properties.



160

4.7. Thermal stability characterization

4.7.1. Differential scanning calorimetry

It is recognized that metallic glass-based materials are thermally stable against

crystallization due to their strong metallic bonds and network structures consisting of

bonded clusters [42, 45, 48, 56, 57]. The structural changes and crystallization behavior of the

Zr-Cu thin films strongly depends on the composition, diffusion process, activation energy

and thermodynamics.

The atomic configuration of these complicated structures usually corresponds to high

degrees of dense random packed structures, which may lead to high viscous materials

within their super-cooled liquid states and slow crystallization [42]. The investigation of

phase formation and crystallization kinetics is hereafter an important approach to evaluate

the thermal stability of metallic glass thin films.

In figure 4.21. is presented an example of a typical DSC curve, while Table 4.6. gathers

thermal parameters measured for the selected films, and compared to their equivalent

BMGs.

Figure 4.21. Example of a typical DSC curve. The Tg, Tx and Tp temperatures obtained at a

heating rate of 20K/min are highlighted. First Tp corresponds to Cu10Zr7 and the second peak (>

700 K) to Cu51Zr14 phase

The thermal behavior recorded during constant-rate heating by DSC give the most

important parameters linked to the glassy-like behavior as Tg on the one hand, and Tx on the

other hand, both gathered into the T (Tx-Tg) supercooled region.

The exothermic peak Tp corresponds to the transformation into the high temperature phase

with respect to initial composition as shown in figure 4.21.

The thermal parameters obtained by DSC are summarized in the next table.


Chapter 4

161

Table 4.6. Thermal parameters of Zr-Cu thin film metallic glasses

Alloys Tg(K)

(±3)

Tx (K)

(±3)

Tp (K)

(±3)

ΔT (K)

(±3)

Zr-33.3 at.% Cu 627 668 679 41

Zr-40.1 at.% Cu 619 684 693 56

Zr-41.3 at.% Cu 623 689 695 66

Zr-48.3 at.% Cu 625 693 731 68

Zr-48.9 at.% Cu 645 716 727 71

Zr-59.1 at.% Cu 672 724 742 52

Zr-79.7 at.% Cu 723 771 782 48

From the table we may see that the Tx and Tg monotonically increases as more copper is

added into the amorphous matrix. The composition dependence of the glass transition

temperature Tg and the crystallization temperature Tx are shown in figure 4.22. and 4.23.,

respectively.

Figure 4.22. Glass transition temperature Tg for the Zr-Cu TFMGs compared with the values

from literature for Zr-Cu BMGs [12]

The results for the Zr-Cu TFMGs are in good agreement with the ones from the literature

even if the recorded values are slightly lower at the higher copper contents.

Regarding the Tx evolution the value are more closer and almost on the same linear fitting

curve as shown in the figure below.

Figure 4.23. The onset crystallization temperature Tx for the Zr-Cu TFMGs compared with the

values from literature for Zr-Cu BMGs [12]

30 40 50 60 70 80580

600

620

640

660

680

700

720

740

760

Tg, TFMG

Tg, BMG

Cu content /at.%

Te

mp

era

ture

/K

30 40 50 60 70 80640

660

680

700

720

740

760

780

800

Tx, TFMG

Tx, BMG

Te

mp

era

ture

/K

Cu content /at.%



162

As highlighted from the table 4.6. and figures 4.22. – 4.23. the thermal parameters increases

continuously as a function of the copper content, which is in accordance with previous data

reported for Cu-Zr system [2, 9, 12, 48, 58].

The supercooled region as a function of the copper contents has also an interesting

evolution, as shown in figure 4.24. It appears the extended region pass through a maximum

in the same region as the crystallization temperature (Figure 4.20.) and FWHM’s parameter

(Figure 4.6.)

Figure 4.24. Evolution of the supercooled region (ΔT) as a function of Cu content

These results infer again better local structural stability in the regime B. In other words, it

means that the glassy structure can sustain higher outside energy (i.e., mechanical force,

heating) without changing its structural state. The films from these regions have the wider

supercooled region, which attests that could have improved properties as compared with

the films from the other regions.

It is important to highlight as well that the thermal stability of the Zr-Cu TFMGs seems to be

higher compared with other similar binary TFMGs. For instance, sputtered Mg-Cu TFMGs

crystallizes at temperatures of about two hundred degrees lower (Tg = 428 K, Tx = 460 K) [23].

4.7.2. Comparison between DSC and XRD results

The Tg, Tx and Tp parameters obtained from DSC experiments differ only of few degrees by

comparison with their BMGs counterparts (Table 2.6 Figures 4.23 – 4.24.), which seems to

confirm the metallic glass-like state of the Zr-Cu thin films. Thermal parameters values

increase continuously with the copper content for the presented compositions.

However, the crystallization temperature (Tp) deduced from DSC appear to be in

contradiction with the crystallization events detected by in situ HT-XRD. These differences

depend mainly on the basis on heating rates 30 K/min between each scan for XRD and 20

K/min for DSC experiments.

The crystallization of the Cu-Zr-based metallic glasses is generally a heating dependent

process influenced especially by the heating rate [59 - 63].


Chapter 4

163

4.8. Oxidation effect during in situ XRD heating

As described in the XRD patterns from section 4.6. the Zr-oxides are always first to be

formed due to well-known affinity of Zr for oxygen combined with long annealing times (~7

h) favored the formation of ZrO2 oxides at early stages of heating (ΔG = -1037 kJ/mol at

standard conditions [53]). These oxides formed on the films’ surface are due to the poor

vacuum in the XRD oven (~ 102 Pa) leading to a reduction of Zr atoms available in the

structure.

Such phenomena is particularly obvious for films of regime B and C, where the expected

phases (Cu10Zr7 and Cu51Zr14 respectively) are not formed and moreover by increasing the

temperature the crystallized phases are towards lower zirconium contents from the phase

diagram (Figure 4.3.), namely Cu51Zr14 and pure Cu.

Regarding the oxides’ nature, the two ZrO2 varieties, monoclinic (stable phase) as well as

tetragonal type as previously presented were obtained. Apart from regime C where Zr

amount is limited, we may assume that zirconia is formed on the top surface, preserving the

“coating’s bulk” where the transformations into intermetallics may occur.

In a further attempt to analyze the oxides influence on the film’s surface we have performed

XRD experiments in grazing incidence mode at several angles to confirm the oxide gradient

through the entire film thickness, as it can be found in figure 4.25.

Figure 4.25. XRD patterns for Zr-59 at.% Cu TFMGs obtained after crystallization and oxidation

at several incidence angles (a) and the zoom for region surrounded, representing the main m-

ZrO2 oxides peaks (a). The measurements have been performed at several incident angles: 2°, 4°,

6.5°, 9.5° and 14°. Distances in µm give the analysed depth for a residual reflected intensity if

10%

The XRD patterns of Zr-59 at.% Cu recorded after crystallization shows the diffracted

domains are changing as the angle of incidence is different. It appears that the intensity of

the monoclinic ZrO2 peaks decreases as the diffracted film’s thickness increases.

For instance, at an angle of incident of 2° the diffracting domain is limited only in the first

microns, meaning that only the top part is analyzed. As the incident angles increase up to

14° the diffracting domain is enlarged and finally all the film thickness is scanned as the

intensities of the zirconia peaks decreases (Figure 4.25. b).

25 30 35 40 45 50 55 60 65

a

5 m at 14°

4 m at 9.5°

3 m at 6.5°

2 m at 4°

2deg

Inte

ns

ity

/a

rb.

un

its

1 m at 2°

h-Cu51Zr14

bco-Cu10Zr7m-ZrO2

t-ZrO2

Inte

ns

ity

/a

rb.

un

its

2deg

5 m at 14°

4 m at 9.5°

3 m at 6.5°

2 m at 4°

1 m at 2°

top film's

surface{1

11}

{-111

}

b

25 26 27 28 29 30 31 32



164

Based on these findings we tested a model that approximate the oxide gradient formed on

the film surfaces. The resulting simulations are reported in figure 4.26. below.

Figure 4.26. Example of simulation model based on the m-ZrO2 oxide intensities acquired from

grazing incident angles

Based on the exponential constant (μm) showed inserted into the figure, we may say

that this evolution endorses the formation of a bilayer film in first 1.3 μm obviously richer in

m-ZrO2. It looks like the oxygen diffuses into the film and that zirconium out diffused on

the surface, forming the oxide phase.

This diffusional process will decrease the Zr-atoms quantity available in the matrix from the

as-deposited state, which will affect the crystallization by favoring the phase’s less-rich in Zr

as is not expected from the phase diagram considering the initial composition (Figure 4.3.).

Once the film is oxidized the initial composition is slightly changed and thus different

compounds are formed as previously showed in figure 4.20.

Furthermore, based in the in situ XRD structural evolutions from the figures 4.14., 4.16., and

4.19. we would like to propose some crystallization models which are valid for all the films

with respected to the compositional regimes (A, B, C) from the figures 4.6. or 4.20.

These three regimes have been always found to be similar whatever the parameters studied

in terms of crystallization resistance or extent of the supercooled region.

The crystallization of metallic glass materials is an important and actively studied

parameter, since closely related to the stability of the metallic glass.


Chapter 4

165

4.9. Towards a model of crystallization for Zr-Cu TFMGs

The in situ XRD patterns showed us that for amorphous materials the final crystalline

phases depends a lot of the initial composition and other factors, such as initial oxidation.

The devitrification of amorphous Zr-Cu TFMGs occurs through different pathways

governed by the thermodynamics (i.e., standard enthalpies of formation) and the diffusion

processes.

In the previous figures we revealed different film’s microstructures and crystallization

temperatures, so that three different regimes have been distinguished.

For each regime, we would like to propose a simplified model of crystallization involving

the different intermetallic compounds identified when heated, as highlighted in 4.20.

Regime A

For regime A (low Cu content), some Zr-based nuclei could be dispersed into the

amorphous matrix after the deposition, explaining the unexpected presence of the fcc

variety of zirconium.

Figure 4.27. Schematic representation of crystallization mechanisms for Zr-40.1 at.% Cu TFMG

corresponding to regime A

The presence of this nanoscale precipitates could be responsible for the initial premature

crystallization, playing most likely a role of germination-activated sites.

The bct-CuZr2 compound agrees to the thermodynamically-expected phase and is the first

phase to be formed (Figures 4.3. and 4.14.).

At higher temperatures, the intermetallic undergoes an allotropic transformation into the

fcc-Zr2Cu variety, characterized by an enlarged unit cell. Such structural modification has

already been attested by Nagase and Umakoshi [54].

The final crystallization stage is reached when some traces of Cu10Zr7 phase are detected

above 843 K.



166

Regime C

In regime C corresponding to the richest-Cu films, copper first crystallizes during annealing

even if the most probable intermetallic phase in this composition domain should have been

Cu51Zr14. This may be due to the initial presence of tiny Cu nuclei resulting from the

deposition step, as identified by HRTEM (Figure 4.12.)

During the early stage of annealing, these nuclei act as germination sites for the growing of

Cu nanodomains. Thermodynamics imposes afterwards a preferential formation of the

Cu51Zr14 phase to the detriment of copper. From 753 K, film is fully crystallized with a

further surface oxidation of zirconium. Therefore, at the film top-surface, Zr is preferentially

consumed by oxygen, inducing Zr depletion (Cu enrichment) on subsurface hindering the

intermetallic compound precipitation as illustrated in figure 4.28.

Figure 4.28. Schematic representation of crystallization mechanisms Zr-79.7 at.% Cu TFMG

corresponding to regime C

Regime B

For regime B, we found an extended compositional zone (45 – 75 at.% Cu), where the

resistance against crystallization is the highest, implying the best stability for these TFMGs.

Several authors studied this matter by molecular dynamics (MDs) for understanding the

atomic packing density and free volume in the Zr-Cu system [13, 15 - 17].

The atomic packing is linked to the geometry and interconnections of coordination

polyhedra, defined by the Voronoi’s indexes [12]. The ideal one is the icosahedral-type

noted as <0,0,12,0> that would be the most efficient atomic cluster governing the GFA and

packing state. Some authors as Peng et al. [15] and Park et al. [1] mentioned that the fraction

of the main considered polyhedra (<0,1,10,4>, <0,2,8,1>, <0,2,8,2>) exhibit a bell-shape

evolution with a maximum at about ~ 60 at.% Cu, which is perfectly confirmed by our

results from Figures 4.6. and figure 4.20.

These polyhedra correspond mainly to Zr-centered clusters, known for their preponderant

role to slow down atomic dynamics, to give rise to dense packing states, and to increase the

glass structure’s stability [12].

The presence of such clusters basically determines the stability of the entire system by

developing strings with other polyhedra to densely fill the remaining space (free volume).

This large chemical stability range corroborates previous results of Park et al. [1].


Chapter 4

167

In figure 4.29. we present an example of such icosahedra configuration for a Cu60Zr40 BMG.

Figure 4.29. Snapshot of the three-dimensional configuration for Cu60Zr40 MG (a) and packing

of 48 perfect icosahedra in the configuration, in each of which the center atom forms twelve 555

pairs with its neighbors (b). The blue and yellow spheres denote Zr and Cu, respectively [64]

From the above analysis it is fairly to assume that the Zr-Cu TFMGs crystallization is

hindered in this region where minimal free volume is expected at about ~60 at.% Cu.

Another complementary viewpoint to explain the high stability of regime B’s films is related

to the presence of a mixture of the two intermetallic phases, creating a more complex and

stable structural state. It becomes more difficult therefore to destabilize such disordered

compacted-like structure for obtaining a crystallized state.

Yamaguchi et al. [65] carried out thermodynamic measurements for Cu-Zr and found that

the standard enthalpy of formation for the two intermetallics have similar negative values.

Based on the Zr-Cu phase diagram (Figure 4.3.) the Cu51Zr14 intermetallic is expected to be

first formed considering that it is the most stable phase. Besides, usually the

thermodynamics gives information related only to the final/stable phase and not about the

first one to be formed. In figure 4.30. is presented the crystallization model proposed.

Figure 4.30. Schematic representation of crystallization mechanisms Zr-59.1 at.% Cu TFMG

corresponding to regime B

Based on our results from figure 4.16. and table 4.4. the Cu10Zr7 phase seems to play a role of

germination phase for the growth of the stable Cu51Zr14 compound.

In general, the metastable phases nucleate faster because they require weaker activation

energy in comparison with the stable phases.



168

Considering all these aspects, we believe that during initial crystallization, the kinetics

controls the preliminary formation of a Cu10Zr7 metastable phase and not directly the

thermodynamics, due to an easier nucleation allowed by icosahedral clusters.

The calculations of activation energy of crystallization for Zr-Cu amorphous alloys [66]

show a low value for the Cu composition close to 60%, which corresponds exactly to the

Cu10Zr7 chemical composition (Zr-59 at.% Cu film).

Within the [663-783 K] temperature range, there is a competition between the continuous

crystallization-coarsening of Cu10Zr7, and the crystallization of Cu51Zr14. All these thermal

induced transformation are gathered in figure 4.16 for the Zr-59.1 at.% Cu TFMG.

To summarize, the three models presented here showed that the films had different

crystallization resistance during heating, which is a direct relation of the local structural

state [1, 15, 16].

For the films from regime B (45 - 75 at.% Cu) where the disordered is supposed to be the

highest, the formation of crystal-like phase took place more difficult.

In the other two domains (A, C) the mixing of the Zr and Cr atoms formed indeed

amorphous structure as well, but their structural stability was found to be less stable when

heated. Such models could serve as general examples-approach for some other binary

TFMGs with similar phase diagrams such as Zr-Ni for instance.

On the other hand, the industry requests are often related to more classical and easier

accessible information as hardness of the thin films proposed for instance.

In what follows, we will present two main functional issues of the Zr-Cu TFMGs, like

mechanical properties or corrosion resistance.


Chapter 4

169

4.10. Mechanical properties

Recently, the researchers have become more and more interested by the mechanical

properties of smaller volumes, such as thin films, since this property has demonstrated to be

of crucial importance for the in-service behavior [5, 31, 67, 68].

The influence of the copper additions in the matrix and its effect on composition and

structural features of the films can also be correlated with the evolution of the mechanical

properties. Such data can be obtained by conventional nano-indentation techniques, but also

by advanced photoacoustic techniques such as Brillouin light scattering and picosecond

ultrasonics method. First technique give information on the mechanical characteristics of

films, while the two latter methods indicate rather the elastic characteristics.

4.10.1. Mechanical characteristics

The nanoindendation is classically the most on-hand experimental techniques used for

mechanical characterization of the thin films and coatings. The values for hardness and

Young’s modulus using the CSM thin films method are shown in figure 4.31.

Figure 4.31. Evolution of hardness and Young’s modulus (indentation test modulus EIT) as a

function of the Cu concentration. The values of metals are taken from ref. [69]

Hardness presents a bell-shape evolution with values ranging from 5–7 GPa for moderate

Cu contents. It can be pointed out that these values are far harder than those of pure metals

and present a maximum plateau around 60 – 80 at.% Cu. The Young's modulus (EIT) is less

sensitive to Cu with a continuous increase from 95 to 120 GPa

The hardness for some films is slightly lower than the values previously reported for the

corresponding Zr-based BMGs (~6.2 GPa) but the evolution are in line with the data

available in the literature for BMGs [3, 11, 70 - 72] or other binary TFMGs [22, 40].

On the other hand, the metallic glass films are supposed to have high strengh limits which

implies interesting inter-related elastic properties.

From these H and E values, we can also express the elastic strain to failure factor (H/E) (this)

as shown in figure 4.32., where the three regimes obtained previously are clearly detected.

The values of H/E reach again a maximum in the compositional same regime centered

around a 60 at% Cu enrichment, as it was already observed for the evolution of FWHM’s

parameter or crystallization temperatures.



170

Figure 4.32. Evolution of H/E ratio as a function of Cu content deduced for Zr-Cu thin films

This evolution indicates that the maximum atomic structural packing for Zr-Cu TFMGS is

therefore confirmed by this analysis. Many studies have indirectly proposed that the

mechanical and physical behaviors of amorphous Zr-Cu alloys are related to their local

atomic structural state [1, 73 - 75]. The atomic packing density for Zr-Cu can be assessed by

the crystallization temperature and H/E ratio.

The films with lower crystallization temperatures are regarded to be structurally less stable

and hence they can go through structural changes under applications of thermal or

mechanical energy. The general tendency not only confirms that the results are reliable, but

also validates the maximum structural stability and better mechanical properties attained

for the compositions in regime B.

Finally, the mechanical properties are compared in figure 4.33. with similar BMGs or

TFMGs from literature and a law as H = 0.074 E has been found. It appears that mechanical

characteristics of our films are of high quality, similar or even better than those of published

comparable coatings.

Figure 4.33. Example of mechanical properties for other Zr-based BMGs or TFMGs adapted from

references [5, 6, 22, 40, 76 - 78]


Chapter 4

171

4.10.2. Elastic characteristics

The general Young’s modulus determined by nano-indentation gives only partial

information related to thin films real elastic behavior.

The photo-acoustic technics have been developed around 20 years ago could to assess

complementary results on elastic characteristics. Their basic principles are more developed

in chapter 2 and in related papers. Measurements have been carried in close collaboration

with Pr. Philippe Djemia from the Laboratoire des Sciences des Procédés et des Matériaux,

Université Paris 13, France. The key-parameters (raw data) to measure experimentally are

VT and VL, , respectively the transverse and longitudinal velocities of the waves detected in

a Brillouin light scattering or picosecond ultrasonics experiment..

To reach elastic parameters, it is first necessary to measure the density of films (ρ). The mass

density was measured at the P’ laboratory (Poitiers, France) by Pr. Gregory ABADIAS using

the X-ray reflectivity (XRR) as shown in figure 4.34. ρ is determined by the best fit of the

experimental curve by an optical model from Parratt [79]. It acceptably reproduces both the

peak positions and the decay of the intensity at very low angles.

Figure 4.34. Representative XRR scans of Zr-Cu films (~ 5000 nm) deposited on (001) Si

substrates. Solid lines correspond to the best-fit simulations to experimental data (symbols).

Vertical dashed lines indicate the position of the critical angle c

Using this description, the mass density ρ of the individual layers were determined for all

samples and presented figure 4.35.a. The mass density parameter could give important

insights related to glassy structure such as the atomic volume calculated via the following

formula (Figure 4.35.b):

Va =M̅

ρNA=

(1−x)MZr+xMCu

ρNA [Eq. 4.2.]



172

Figure 4.35. Experimental data for Zr-Cu films: (a) mass density ρ deduced from simulations of

the XRR scans; (b) average atomic volume Va deduced from ρ using Eq. (4.5.). The line is linear

extrapolation of the data obtained in the range 30 < Cu < 80 and the dashed line is the Vegard’s

rule from respective bulk polycrystalline values (cross symbol)

The mass density increases according to a linear law as more copper is added in the matrix,

evolution that is necesseraly related to the different local atomic arrangements (state). In

fact, the atomic volume which discussion is more physcially appropriate, is decreasing

linearly as a function of the Cu concentration, resulting to densier films, as expected from

Vegard’s rule using the polycrystalline Zr and Cu element parameters. The linear variation

implies that alloys behave as ideal solid solutions with a simple average of the atomic

volume of each pure elements Zr and Cu. This evolution his also consistent with the

continuous shift of hump positions with the addition of copper in the matrix (Figure 4.4.).

In both cases the same trend implies that copper additions continuously move/reorganize

the atoms at the local range. Hence, the shrinkage of the distances between the closest

neighbors as more copper is added increases the mass density leading finally to a more

compact structures.

Moreover, one can notice that the extraoplation to x=0 leads to an estimate value of the

“hypothetical” amorphous Zr Va(a-Zr)=22.52 10-3 nm3, that is slightly below (3%) the

polycrystalline one for hcp-Zr Va(α-Zr)= 23.23 10-3 nm3. On the contrary, the extrapolation

for x=1 leads to a value very closed to the cfc polycristalline Cu Va(Cu)= 11.78 10-3 nm3. In

case of Cu BMGs alloys, I. Bakonyi evidenced an ideal solution, too, but with hcp-Zr Va(α-

Zr) bound [80].

A deeper and fully investigation of elastic properties by means of brillouin light scattering

(BLS) and picosecond ultrasonics (PU) was further performed on several compositions. The

experimental details and how these techniques are used can be found in refs. [81, 82].

Once the mass density is derminated from by fitting the XRR scans using an optical model

[79], the elastic constants from PU and BLS measurements of sound velocities, VL=(C11/)1/2

and VT=(C44/)1/2, respectively, can be extracted. C11 is the longitudinal elastic constant and

C44 = G = (C11 - C12)/2 the shear elastic constant/modulus.


Chapter 4

173

These two independent elastic constants fully characterize the elastic properties of our

amorphous and highly disordered nanocrystallyne coatings with isotropic elastic properties

through the following formula:

shear modulus 44

1211

2C

CCG

, [Eq. 4.3.]

Young’s modulus 4411

2

441111

CC

2CCC

E , [Eq. 4.4.]

bulk modulus 44113

4CCB , [Eq. 4.5.]

Poisson ratio 4411

4411

CC 2

2CC

, [Eq. 4.6.]

Figure 4.35.a. and 4.35.b show an example of a PU spectrum (Zr 57 at.% Cu) and of a BLS

spectrum (Zr-48a t.%Cu), respectively In case of the PU experiment, after the thermal decay,

the first echo is clearly seen after a time of flight TOF=2015 ps and for a distance d=2 x 4200

nm. Thus, we find VL=d/TOF=4169 m/s and C11=(VL)2= 134 GPa. In the case of the BLS

experiment, the main feature is associated to the Rayleigh surface wave (R). We find from a

best fit of the spectrum C44 = 26 GPa and VT=1903 m/s. We can notice in this particular case,

that the signal/noise ratio is very favourable, enabling the observation of the longitudinal

threshold singularity [82] linked to VL.

Figure 4.35. A typical PU spectrum from Zr- 57 at.% Cu (black line). The time of flight (TOF) is

obtained between the coincidence (pump and probe arriving at the same time) and the first echo

(a). A typical BLS spectrum from Zr- 48 at.% Cu (black line) for an angle of incidence of light =

65° and a best fit considering the ripple mechanism (red line). R denotes the Rayleigh surface

wave and LT the longitudinal threshold frequency (b)



174

For all samples, the C11 elastic constant is always first measured by PU and then considered

as a fixed known parameter in the fitting of the BLS spectra, leading finally to the evaluation

of the second C44 parameter. Evolution of VL and VT as a function of Cu concentration is


Figure 4.36. Evolution of the sound velocities VL and VT as a function of the Cu concentration

The dashed lines are the Vegard’s rule, calculated from the pure polycrystalline element

properties (cross symbols)

Both sound velocities seem to follow the same trend with values well fitted in the

amorphous region on a linear law as previously spotted for the mass density. The cristalline

to amorphous transition in the Zr-rich part is clearly evidenced by the catastrophic decrease

of the VL (Cu > 33 and VT (Cu > 13). This confirms actually the different structural state,

crytalline for Zr-13.4 at.% Cu, nanocrystalline (25 at.% Cu and 28.5 at.% Cu) as previously

spotted in the XRD patterns from figure 4.4.

From these velocity values we can extract C11 and C44 parameters, presented respectively in

figure 4.37. a. and 4.37.b as a function of the Cu concentration.

Figure 4.37. Evolution of the a) C11 and b) C44 elastic constants as a function of the Cu

concentration. The dashed lines are the Vegard’s rule calculated from the pure polycrystalline

element properties (cross symbols)


Chapter 4

175

The C11 constant seems to be much more sensitive to the presence of nanocrystalline phases

which could explain the apparent scattering of the data in figure 4.37.a, as relative

uncertainty in the thickness measurments does not exceed a few percents. The scattering in

C44 data is always less pronounced as it does not depend on the thickness of these thick

coatings.

The C44 elastic constant increases linearly from 22 GPa to 35 GPa in the amorphous region

and leads to an extroplated value for a “hypothetical” amorphous zirconium a-Zr of 17 GPa

(well below the polycrystalline value for hcp α-Zr) and for a-Cu 37 GPa (lower than cfc-Cu).

Higher value of shear modulus G=C44 should result in a decrease of the ductility of the

materials. In the case of the second constant, namely C11, we can obtain an extroplated value

for a “hypothetical” amorphous zirconium a-Zr of 84 GPa (well below the polycrystalline

value for hcp α-Zr) and for a-Cu 184 GPa (very closed to the cfc-Cu). The Young’s modulus

E can be calculated from the two independent elastic constants C11 and C44. (Eq. 4.4).

Figure 4.38. compares the Young’s modulus obtained from nanoindendation technique (EIT)

with the one derived from the combination of these two non-conventional photo-acoustic

techniques.

Figure 4.38. Young’s modulus E and nano-indentation test modulus EIT as a function of the Cu

concentration. The dashed lines are the Vegard’s rule, calculated from the pure polycrystalline

element properties (cross symbols)

We observe a similar behavior of the Young’s modulus as a function of the Cu concentration

whatever the employed method. The values gradually increases as more copper is

incorporated from about 70 GPa to 100 GPa for E(BLS+PU) in the amorphous region and leads

to an extrapolated value for a “hypothetical” amorphous zirconium a-Zr of 40 GPa, well

below the polycrystalline value for hcp α-Zr but in good agreement with the one

extrapolated from results on BMGs from Ristic et al. [83].



176

For nanoindentation, EIT varies from 100 GPa to 120 GPa with only a slight deviation of EIT

from linear Vegard’s rule [10]. The difference (10-20%) between the values calculated by

BLS+PU and nanoindentation may be only partly related to a substrate effect (ESi = 183 GPa)

not completely withdraw using nanoindentation technique even for these thick coatings

(~5000 nm) and in a less extent to the selected strain rate. Certainly, other effects contribute

to these differences, may be the volume that is probed by those techniques is different and

can lead to a quantitative different elastic response.

The elastic moduli give also a global macroscopic view of films stiffness, since its values

reflect both the interatomic bonding energies and local connectivity. The evolution of this

parameter indicates that the change of Cu content leads to microstructural change and

furthermore increases the elastic properties of the TFMGs. Several authors reported that this

phenomenon could be due to the intrinsic structural characteristics of the glassy films

consisting of packed atomic clusters. Finally, this kind of local structures allows efficient

packing of atom clusters that show evidently different properties [1, 84 - 86].

4.10.3. Mechanical behavior

Once we have measured the mechanical and elastic parameters, we can now identify some

correlations in-between and give some prediction on the plastic behavior.

In figure 4.39., the G shear modulus (C44 elastic constant) increases linearly as a function of

Young’s modulus E.

Figure 4.39. Shear modulus G as a function of the Young modulus E of ZrCu thin films metallic

glasses

These two elastic moduli are strongly inter-related through the best fit equation

G = 0.3648 E. The ratio G/E is slightly lower than 3/8 experimentally found for

polycrystalline materials [87]. It is also lower than 0.39 previously theoretically estimated for

fcc polycrystalline metals but very close to 0.367 found for bcc [88] providing us with a good

confidence in our results.

Furthermore, Pugh initially proposed an other exploitation of these results based on the G/B

ratio, developed initially for polycrystalline metals. It introduced a critical ratio (G/B)crit

allowing to define some limit domains where the material behaves as a ductile (at low G/B

values) or brittle (at higher G/B values) material [89 - 91].


Chapter 4

177

Later, Pettifor introduced a criterion linking the brittle-to-ductile transition in

polycrystalline cubic materials to the Cauchy pressure (C12−C44) [92]. Based on his findings,

positive Cauchy pressures correspond to a metallic character and to a ductile behaviour of

the material, while negative values suggest covalent bonding, and brittle behaviour.

More recently, these concepts have been extended to metallic and ceramic glasses, for which

a reliable set of data is now available. The brittle-to-ductile threshold occurs at (G/B)crit =0.41

(or equivalently =0.32) [93 - 95]. This lower critical value reflects the overall difficulty for

an amorphous material to deform plastically. The plastic flow in metallic glasses occurs very

locally in shear bands [93, 96] compared with polycrystalline metals, where flow is

dislocation mediated and delocalized by associated work hardening.

In figure 4.40.a. we ploted G/B as a function of Cu concentration while in figure 4.40.b., we

used the (G/B) and (C12 −C44) coordinates to plot the elastic properties of the whole set of

Zr-Cu alloys studied here.

Figure 4.40. Pugh criterion G/B ratios as a function of Cu concentration (a), G/B ratios as a function of

the Cauchy pressure (C12-C44) (b)

Whatever the considered formalism, G/B keeps for all films a small value below 0.41 (0.175-

0.32) as it is usual in amorphous alloys and (C12 −C44) > 0. The Cauchy pressure remains as

high as ∼35 - 90 GPa, suggesting an intrinsic stronger resistance of metallic glasses to shape

changes, as compared to crystalline materials. We may observe also that the values are

systematically situated in the region (C12 – C44) > 0, which implies a metallic character for all

films.

The ratio G/B reaches a value as low as 0.175 and a high positive Cauchy pressure (C12-C44) =

80 GPa that should result to the highest ductility behaviour of Zr – 33 at.% Cu film.

Such high ductility is for this coating is associated to a strong liquid-like character. Another

equivalent picture of the ductility of films can be obtained from the values of Poisson ratio

as illustrated in figure 4.41.



178

Figure 4.41. Poisson ratio as a function of Cu concentration

The results of the Poisson ratio indicate that can be viewed as a useful index to predict the

fragility and distinct structural differences among these Zr-Cu TFMGs with similar

structural characteristics consisting of efficiently packed clusters. Similar results have also

been reported in CuZr-based BMGs-forming system [97] and many studies reported that

high Poission’s ratio in metallic glasses could be correlated with a good plasticity [93, 98, 99]

Apart a nearly constant evolution at about 0.38 for most of the copper contents, some points

indicate high variation such as 30, 40 and 85 at. % Cu in particular, where high values are

observed. The highest deformability of the film containing 33 at.% Cu is to underlined.

Lastly, these findings point out as well that plasticity of Zr-Cu TFMGs are susceptive to

their composition since minor content deviation can change the fragility and/or Poisson

ratio, thus altering the microstructure. The results indicate that the composition sensitivity

to plasticity can be clearly reflected by the Poisson’s ratio.

The investigation of elastic properties by means of BLS and PU could be essentially used to

propose models able to predict the plastic behavior of metallic glass thin films since such

materials are supposed to be structurally and chemically homogeneous.

In terms of main features that may explain these tendencies it is important to state that, first

of all, the evolution of the mechanical properties and elastic constants are well correlated

with the structural evolution. Therefore, the mechanical properties are likely affected by the

atomic bonding and packing state [3, 5]. By enriching Zr–TFMG matrix with Cu, the elastic

parameters follow a linear evolution due to a change of the atomic packing through the

modification of bonding length and/or the shrinkage of thermal stability.

To conclude, whatever the measurement technique adopted, the high mechanical

characteristics of films have been proven, far higher than those of pure metals. Further

complementary advanced spectroscopic measurements have given access to elastic

characteristics, following a Vegard’s law in the composition range where films are

amorphous but with hypothetical amorphous bounds elements instead of the

polycrystalline ones. Moreover, these findings have pointed out as well that plasticity of Zr-

Cu TFMGs are dependent on their composition since minor content deviation can change

the Poisson’s ratio, thus altering the microstructure. The results indicate that the

composition sensitivity to plasticity can be clearly reflected by the Poisson’s ratio.

Interpretation of elastic constant confirmed the metallic glass character of all films.


Chapter 4

179

4.11. How to improve durability of surface with Zr-Cu TFMGs?

The metallic glass-based materials may have a high potential for applications where the

materials undergo corrosion attacks or oxidation. To date, several studies accredited a

superior corrosion and wear resistances of bulk metallic glasses based on late transition

metals, such as Zr [100, 101], Cu [102], Ni [103] or Ti [104 - 106] owing to their structural and

chemical homogeneity.

In the present section, we intend to present the corrosion behavior of Zr-Cu TFMGs

according to their electrochemical response with respect to structural state. Since the

corrosion occurs via oxide-reduction reactions, the electrochemical techniques are ideal for

studying the corrosion processes.

The corrosion resistance of materials is thus investigated by immersion of the film surface

into a sulfate solution in a controlled manner in order to measure the current flows as a

function of potential.

The corrosion behavior is often described by the current density, measured through

potentiodynamic curves as shown in figure 4.44. for the Zr-Cu deposited into inert

substrates.

Figure 4.44. Potentiodynamic polarization curves of Zr-Cu thin films tested in aerated Na2SO4

solution

Compared to alloys films, the almost pure copper coating shows an outstanding behavior,

which has to be correlated to its particular chemical homogeneity. For other binary films, the

incorporation of copper leads to an increase of the anodic potential. With low amount of

copper, the zirconium can form a passive film on the surface, but the copper does not

behave as a passive metal in such conditions. Therefore, the surface is less protective and the

corrosion is faster when more copper is added. This will shift the potential of the films to the

more negative direction, as actually observed for the graph above.

From the Tafel’s plot it can be observed a similar cathodic part, suggesting the same

reactivity for the cathodic reaction occurring at the surface. Considering again the anodic

part, a plateau attributable to a passive character of the surface is noticed for all the samples,

except the 98 at.% Cu. For moderate alloyed films (13.4 and 33.3 at.% Cu) an activation peak

of reactivity directly rises as indicated by the abrupt increase of the current density just after

the corrosion potential.



180

Copper cannot efficiently passivate the surface in Na2SO4 solution so that the passivation of

films is caused mainly by the presence of Zr, which has a high susceptibility to passivation

[107, 108]. Besides, it may be observed as well that the critical passivation current increases,

whereas the potential shifts to more negative direction with increasing of Cu contents, as

illustrated in table 4.7.

Table 4.7. Evolution of corrosion parameters as a function of Cu content

Cu content

(at.%)

Ecorr

(mV/SCE)

icorr

(µA/cm2)

Current

density

(µA/cm2)

at 0.37 V

13.4 58.29 0.3 ~1.5

33.3 -105.3 1.4 ~19

40.1 -110.7 1.6 ~23

48.3 -219.3 5.0 ~8.6

98.0 -45 10.0 ~10000

In the table is seen that the 13.4 at. % Cu film has the lowest corrosion rate and most positive

corrosion potential, which imply the best corrosion behavior in Na2SO4 solution among the

Zr-Cu thin films. This result is primarily attributed to the substantial ability of zirconium as

typical valve-metal, to form a highly passive film on the thin film surface [109].

It is well recognized that Zr-based oxides formed on the top-surface are usually stable and

denser that leads to the formation of a protective film associated with a high corrosion

resistance [32, 33]. The copper at this level does not influence at all the chemical or structural

state, as the film is crystalline into pure Zr-fcc phase.

From the table above we may see as well that as more copper is added into Zr-Cu matrix the

Icorr increases from 0.3 to 10 µA/cm2 for films containing 13.4 and 98 at.% Cu, respectively.

The corrosion potential changes from positive range for Zr-13.4 at. % Cu to more negative

values, which imply a modified interface as Cu content increases. The evolution of the

current density and thus the corrosion behavior for the films can be assessed by comparing

the current density values at a same corrosion potential from the anodic part.

The last column from the table 4.36., presents the values recorded at a potential of 0.37 V.

This potential has been chosen because it is situated on the anodic domain for all the thin

films, which could give further insights on the passivation ability.

The main remark is that for all samples the same tendency is observed as for the general

evolution of the icorr, namely as more copper is added into the matrix the icorr increases from

1.5 µA/cm2 for the crystalline Zr- 13.4 at. % to 1 x 104 µA/cm2 for the crystallized Zr-98 at. %

Cu. However, comparison of this two extreme films is not really appropriated between they

are strictly different in terms of structure on the one hand, of chemistry on the other hand.

Nevertheless, what is sure is that a too high copper level affects the passive-ability of the

film. A compromise has then to be found to combine corrosion protectiveness and

interesting mechanical properties.


Chapter 4

181

Besides, the electrochemical impedance spectroscopy (EIS) could give more insights on the

electrochemical dynamics and changes undergoing at the thin films’ surface. The EIS spectra

for Zr-Cu thin films recorded after 3 hours immersion time in Na2SO4 solution are displayed

in figure 4.37.

Figure 4.48. Nyquist representation for the Zr-Cu thin films

The Zr-13.4 at.% Cu reveals the most capacitive behavior which conducts to the best

resistance thanks to its high quality passive film. In opposite, the Zr-98 at.% Cu film does

not form a protective barrier and indicates a small circle associated with a poor corrosion

resistance.

For the amorphous compositions (33.3, 40.1 and 48.3 at. % Cu), the Nyquist diagrams

indicate that the films are characterized by not well separated capacitive loops with similar

evolution. Considering the high values of charge transfer resistances, we can conclude that

corrosion protectiveness is important but difficult to be discriminated between the different

films’ compositions. This issue has to be further investigated and a deeper study is required

to better understand the influence of the copper contents on the corrosion resistance.

In this perspective, our complementary study on other Zr-Me (Me = Ni, Co) TFMGs

submitted for publication presents a more detailed approach with respect to corrosion

resistance of such type of thin films where the Zr is the main constituent element.

The results are not presented since is not the purpose of the current project.

The main conclusion from the corrosion tests is that the copper addition at an extreme level

may have a harmful influence on the corrosion resistance of Zr-Cu thin films.

In opposite, for low additions, Zr affords to the film its high passive-ability. For moderate

Cu contents, corrosion resistance remains very high, such characteristic being attributable to

the formation of a passive protective film.



182

Chapter summary

In this chapter it has been presented the metallic glass Zr-Cu thin films deposited by

magnetron sputtering in a wide compositional range scanning the whole phase diagram.

The adopted deposition setup involved pure metal targets (Cu and Zr), which allowed a fine

control of the films’ chemical composition.

The samples investigated have been divided in three categories for an easier comparison

between their behaviors and properties as regime A (< 45 at. % Cu), regime B (45 at.% <

content < 75 at.% Cu) and regime C (Cu content > 75 at.% Cu).

The films have an XRD-amorphous structure in the 33.3 – 89.1 at.% Cu composition range,

with a fine grained morphology. The metallic glass films exhibit typical vein-like features

when fractured and studied by SEM in cross-section which are among the characteristics of

classical metallic glasses materials. The presence of such features on the final fracture

implies a high plastic deformation domain.

A deeper TEM-scale analysis showed that the films’ fine structure mainly consists in

organized nanodomains embedded into the glassy matrix.

Thermal parameters have been attained by means of DSC and in situ XRD experiments.

An important section is devoted to the crystallization resistance and thermal stability of

these thin film metallic glasses, where a crystallization mechanism model is further

presented.

Among the compositions studied, in the 45 – 75 a t. % Cu (corresponding to regime B) the

films have the highest crystallization resistance and good thermal stability. The

crystallization process is governed through a multistage process controlled by the time and

the temperature. For all three regimes, a model of devitrification of films with the

temperature is proposed

Mechanical properties of TFMGs have been found to be around 95 – 120 GPa for the

Young’s Modulus and 5 – 7 GPa for hardness. The elastic properties have confirmed their

ductile state and composition dependence.

Whatever the investigation scale or technique employed, films belonging to the 45 – 75 at.

% cu composition ranges were systematically identified as the best.

In the final part of the chapter, the corrosion resistance of these materials is tested and

compared with their copper very rich (98 at.% Cu) or very poor (13.4 at.% Cu) crystallized

counterparts.

It has finally been proved that metallic glass films could be a suitable candidate for

corrosion protection thanks to their high chemical and structural homogeneity.


Chapter 4

183

List of references

[1] K.-W. Park, J. Jang, M. Wakeda, Y. Shibutani, and J.-C. Lee, “Atomic packing density and

its influence on the properties of Cu–Zr amorphous alloys,” Scripta Materialia, vol. 57, no.

9, pp. 805–808, 2007.

[2] T. Minemura, J. J. van den Broek, and J. L. C. Daams, “Formation and thermal stability of

amorphous Cu‐Zr thin films deposited by coevaporation,” Journal of Applied Physics, vol.

63, no. 9, pp. 4426–4430, 1988.

[3] P. Coddet, F. Sanchette, J. C. Rousset, O. Rapaud, and C. Coddet, “On the elastic modulus

and hardness of co-sputtered Zr–Cu–(N) thin metal glass films,” Surface and Coatings

Technology, vol. 206, no. 17, pp. 3567–3571, 2012.



Alloys Compd. 615, S348-S351, 2014


metallic glass with high content of immiscible tantalum,” Surface and Coatings Technology,

vol. 205, no. 2, pp. 587–590, 2010.

[6] F. Zeng, Y. Gao, L. Li, D. M. Li, and F. Pan, “Elastic modulus and hardness of Cu–Ta

amorphous films,” Journal of Alloys and Compounds, vol. 389, no. 1–2, pp. 75–79, Mar. 2005.

[7] H. Okamoto, “Cu-Zr (Copper-Zirconium),” J Phs Eqil and Diff, 29, no. 2, pp. 204–204, 2008.

[8] G. A. Almyras, G. M. Matenoglou, P. Komninou, C. Kosmidis, P. Patsalas, and G. A.

Evangelakis, “On the deposition mechanisms and the formation of glassy Cu–Zr thin

films,” Journal of Applied Physics, vol. 107, no. 8, pp. 084313–084313–6, 2010.

[9] B. F. Lu, J. F. Li, L. T. Kong, and Y. H. Zhou, “Correlation between mechanical behavior

and glass forming ability of Zr–Cu metallic glasses,” Intermetallics, 19, 7, 1032–1035, 2011.

[10] A. Denton and N. Ashcroft, “Vegard’s law,” Physical Review A, vol. 43, no. 6, pp. 3161–

3164, 1991.

[11] C. J. Chen, J. C. Huang, H. S. Chou, Y. H. Lai, L. W. Chang, X. H. Du, J. P. Chu, and T. G.

Nieh, “On the amorphous and nanocrystalline Zr–Cu and Zr–Ti co-sputtered thin films,”

Journal of Alloys and Compounds, vol. 483, no. 1–2, pp. 337–340, 2009.

[12] N. Mattern, A. Schöps, U. Kühn, J. Acker, O. Khvostikova, and J. Eckert, “Structural

behavior of CuxZr100−x metallic glass (x = 35−70),” Journal of Non-Crystalline Solids, vol.

354, no. 10–11, pp. 1054–1060, 2008.

[13] S. Ogata, F. Shimizu, J. Li, M. Wakeda, and Y. Shibutani, “Atomistic simulation of shear

localization in Cu–Zr bulk metallic glass,” Intermetallics, 14, no. 8–9, pp. 1033–1037, 2006.

[14] G. A. Almyras, C. E. Lekka, N. Mattern, and G. A. Evangelakis, “On the microstructure of

the Cu65Zr35 and Cu35Zr65 metallic glasses,” Scripta Materialia, 62, no. 1, pp. 33–36, 2010.

[15] H. L. Peng, M. Z. Li, W. H. Wang, C.-Z. Wang, and K. M. Ho, “Effect of local structures

and atomic packing on glass forming ability in CuxZr100−x metallic glasses,” Applied

Physics Letters, vol. 96, no. 2, pp. 021901–021901–3, 2010.

[16] I. Kaban, P. Jóvári, V. Kokotin, O. Shuleshova, B. Beuneu, K. Saksl, N. Mattern, J. Eckert,

and A. L. Greer, “Local atomic arrangements and their topology in Ni–Zr and Cu–Zr

glassy and crystalline alloys,” Acta Materialia, vol. 61, no. 7, pp. 2509–2520, 2013.

[17] T. Takeuchi, S. Nakano, M. Hasegawa, K. Soda, H. Sato, U. Mizutani, K. Itoh, and T.

Fukunaga, “Local atomic arrangements and electronic structure of Zr-Ni-Al bulk metallic

glass: –Analysis by use of the relevant crystals–,” Materials Science and Engineering: A, vol.

449–451, pp. 599–604, 2007.

[18] B. Rointan F., “Handbook of deposition technologies for films and coatings pag 35

chapter 1,” Novlel publications, 1994.

[19] R. F. Bunshah, Handbook of Hard Coatings: Deposition Technolgies, Properties and Applications,

1 edition. Park Ridge, N.J. : Norwich, N.Y: William Andrew, 2002.



184

[20] K. S. Sree Harsha, “Chapter 9 - Nucleation and Growth of Films,” in Principles of Vapor

Deposition of Thin Films, K. S. S. Harsha, Ed. Oxford: Elsevier, 2006, pp. 685–829.

[21] B. C .R. and et al., “Encyclopedia of Materials Characterization: Surfaces, Interfaces, Thin

Films”,” 1992.

[22] J. P. Chu, J. S. C. Jang, J. C. Huang, H. S. Chou, Y. Yang, J. C. Ye, Y. C. Wang, J. W. Lee, F.

X. Liu, P. K. Liaw, Y. C. Chen, C. M. Lee, C. L. Li, and C. Rullyani, “Thin film metallic

glasses: Unique properties and potential applications,” Thin Solid Films, vol. 520, no. 16,

pp. 5097–5122, 2012.

[23] J. C. Huang, J. P. Chu, and J. S. C. Jang, “Recent progress in metallic glasses in Taiwan,”


[24] Y. H. Liu, T. Fujita, A. Hirata, S. Li, H. W. Liu, W. Zhang, A. Inoue, and M. W. Chen,

“Deposition of multicomponent metallic glass films by single-target magnetron

sputtering,” Intermetallics, vol. 21, no. 1, pp. 105–114, 2012.

[25] J.-H. Chu, H.-W. Chen, Y.-C. Chan, J.-G. Duh, J.-W. Lee, and J. S.-C. Jang, “Modification of

structure and property in Zr-based thin film metallic glass via processing temperature

control,” Thin Solid Films., 561, 38-42, 2014

[26] N. Nagendra, U. Ramamurty, T. T. Goh, and Y. Li, “Effect of crystallinity on the impact

toughness of a La-based bulk metallic glass,” Acta Materialia, 48, 10, pp. 2603–2615, 2000.

[27] P. Murali and U. Ramamurty, “Embrittlement of a bulk metallic glass due to sub-Tg

annealing,” Acta Materialia, vol. 53, no. 5, pp. 1467–1478, 2005.

[28] P. J. Hsieh, S. C. Lin, H. C. Su, and J. S. C. Jang, “Glass forming ability and mechanical

properties characterization on Mg58Cu31Y11−xGdx bulk metallic glasses,” Journal of

Alloys and Compounds, vol. 483, no. 1–2, pp. 40–43, 2009.

[29] H. Peng, S. S. Li, Y. P. Qi, and T. Y. Huang, “Mg–Ni–Gd–Ag bulk metallic glass with

improved glass-forming ability and mechanical properties,” Intermetallics, vol. 19, no. 7,

pp. 829–832, 2011.

[30] J. Eckert, J. Das, S. Pauly, and C. Duhamel, “Mechanical properties of bulk metallic glasses

and composites,” Journal of Materials Research, vol. 22, no. 02, pp. 285–301, 2007.

[31] C. A. Schuh, T. C. Hufnagel, and U. Ramamurty, “Mechanical behavior of amorphous

alloys,” Acta Materialia, vol. 55, no. 12, pp. 4067–4109, 2007.

[32] A. L. Greer, “Metallic glasses: Damage tolerance at a price,” Nat Mater, 10, 2, 88–89, 2011.

[33] W. Ma, H. Kou, J. Li, H. Chang, and L. Zhou, “Effect of strain rate on compressive

behavior of Ti-based bulk metallic glass at room temperature,” Journal of Alloys and

Compounds, vol. 472, no. 1–2, pp. 214–218, 2009.

[34] W. J. Wright, T. C. Hufnagel, and W. D. Nix, “Free volume coalescence and void

formation in shear bands in metallic glass,” Journal of Applied Physics, 93, 3, 1432, 2003.

[35] C. E. Packard and C. A. Schuh, “Initiation of shear bands near a stress concentration in

metallic glass,” Acta Materialia, vol. 55, no. 16, pp. 5348–5358, 2007.

[36] Y. Ritter and K. Albe, “Thermal annealing of shear bands in deformed metallic glasses:

Recovery mechanisms in Cu64Zr36 studied by molecular dynamics simulations,” Acta

Materialia, vol. 59, no. 18, pp. 7082–7094, 2011.

[37] G. S. Yu, J. G. Lin, M. Mo, X. F. Wang, F. H. Wang, and C. E. Wen, “Effect of relaxation on

pressure sensitivity index in a Zr-based metallic glass,” Materials Science and Engineering:

A, vol. 460–461, pp. 58–62, 2007.

[38] C. A. Schuh, A. C. Lund, and T. G. Nieh, “New regime of homogeneous flow in the

deformation map of metallic glasses: elevated temperature nanoindentation experiments

and mechanistic modeling,” Acta Materialia, vol. 52, no. 20, pp. 5879–5891, 2004.

[39] T.-P. Hsiao, Y.-C. Yang, J.-W. Lee, C.-L. Li, and J. P. Chu, “The microstructure and

mechanical properties of nitrogen and boron contained ZrCuAlNi thin film metallic glass

composites,” Surface and Coatings Technology. 237, 276-283, 2013


Chapter 4

185




[41] K. Kondoh, J. Fujita, J. Umeda, and T. Serikawa, “Estimation of Compositions of Zr-Cu

Binary Sputtered Film and Its Characterization,” Research Letters in Materials Science, vol.

2008, pp. 1–5, 2008.

[42] X. Cui, F. Q. Zu, W. J. Zhang, Z. Z. Wang, and X. Y. Li, “Phase competition of Cu64Zr36

and its effect on glass forming ability of the alloy,” Crystal Research and Technology, vol. 48,

no. 1, pp. 11–15, 2013.

[43] K. S. Lee, Y. M. Jo, and Y.-S. Lee, “Crystallization and high-temperature deformation

behavior of Cu49Zr45Al6 bulk metallic glass within supercooled liquid region,” Journal of

Non-Crystalline Solids, vol. 376, pp. 145–151, 2013.

[44] S. Zhang, T. Ichitsubo, Y. Yokoyama, T. Yamamoto, E. Matsubara, and A. Inoue,

“Crystallization Behavior and Structural Stability of Zr50Cu40Al10 Bulk Metallic Glass,”

Materials transactions, vol. 50, no. 6, pp. 1340–1345.

[45] Y. T. Cheng, T. H. Hung, J. C. Huang, P. J. Hsieh, and J. S. C. Jang, “Thermal stability and

crystallization kinetics of Mg–Cu–Y–B quaternary alloys,” Materials Science and

Engineering: A, vol. 449–451, pp. 501–505, 2007.

[46] J. S. C. Jang, Y. W. Chen, L. J. Chang, H. Z. Cheng, C. C. Huang, and C. Y. Tsau,

“Crystallization and fracture behavior of the Zr65−xAl7.5Cu17.5Ni10Six bulk amorphous

alloys,” Materials Chemistry and Physics, vol. 89, no. 1, pp. 122–129, 2005.

[47] K. B. Surreddi, S. Scudino, M. Sakaliyska, K. G. Prashanth, D. J. Sordelet, and J. Eckert,

“Crystallization behavior and consolidation of gas-atomized Al84Gd6Ni7Co3 glassy

powder,” Journal of Alloys and Compounds, vol. 491, no. 1–2, pp. 137–142, 2010.

[48] D. Xu, B. Lohwongwatana, G. Duan, W. L. Johnson, and C. Garland, “Bulk metallic glass

formation in binary Cu-rich alloy series – Cu100−xZrx (x=34, 36, 38.2, 40 at.%) and

mechanical properties of bulk Cu64Zr36 glass,” Acta Materialia, vol. 52, no. 9, pp. 2621–

2624, 2004.

[49]

magnetron co-sputtering,” Thin Solid Films, vol. 275, no. 1–2, pp. 164–167, 1996.

[50] J. Qin, X. Yong, Z. Xin-Yu, L. Gong, L. Li-Xin, X. Zhe, M. Ming-Zhen, and L. Ri-Ping, “Zr-

Cu Amorphous Films Prepared by Magnetron Co-sputtering Deposition of Pure Zr and

Cu,” Chinese Physics Letters, vol. 26, no. 8, p. 086109, 2009.

[51] J.-H. Shin, D.-I. Kim, K. Cho, H. Suematsu, K. H. Kim, and R. Nowak, “Plasticity-induced

nanocrystallization in a Zr65Cu35 thin film metallic glass,” Journal of Non-Crystalline

Solids, vol. 362, pp. 65–68, 2013.

[52] J. Musil and P. Zeman, “Structure and microhardness of magnetron sputtered ZrCu and

ZrCu-N films,” Vacuum, vol. 52, no. 3, pp. 269–275, 1999.

[53] W. M. Haynes, CRC Handbook of Chemistry and Physics, 94th Edition, 94 edition. Boca Raton

(Fla.); London; New York: CRC Press, 2013.

[54] Y. U. T. Nagase, “Phase stability of amorphous and crystalline phases in melt-spun

Zr66.7Cu33.3 alloy under electron irradiation,” Scripta Materialia, vol. 48, no. 9, pp. 1237–

1242, 2003.

[55] L. Xie, P. Brault, A.-L. Thomann, and L. Bedra, “Molecular dynamic simulation of binary

ZrxCu100−x metallic glass thin film growth,” Applied Surface Science, vol. 274, pp. 164–170,

2013.

[56] Y.-L. Gao, J. Shen, J.-F. Sun, G. Wang, D.-W. Xing, H.-Z. Xian, and B.-D. Zhou,

“Crystallization behavior of ZrAlNiCu bulk metallic glass with wide supercooled liquid

region,” Materials Letters, vol. 57, no. 13–14, pp. 1894–1898, 2003.



186

[57] D. Okai, Y. Shimizu, N. Hirano, T. Fukami, T. Yamasaki, and A. Inoue, “Isothermal

crystallization in supercooled liquid state for Ca50Mg22.5Cu27.5 metallic glass,” Journal of

Alloys and Compounds, vol. 504, Supplement 1, pp. S247–S250, 2010.

[58] D. Wang, Y. Li, B. B. Sun, M. L. Sui, K. Lu, and E. Ma, “Bulk metallic glass formation in

the binary Cu–Zr system,” Applied Physics Letters, vol. 84, no. 20, p. 4029, 2004.

[59] J. C. Qiao and J. M. Pelletier, “Crystallization kinetics in Cu46Zr45Al7Y2 bulk metallic

glass by differential scanning calorimetry (DSC),” Journal of Non-Crystalline Solids, vol.

357, no. 14, pp. 2590–2594, 2011.

[60] J. C. Qiao and J. M. Pelletier, “Enthalpy relaxation in Cu46Zr45Al7Y2 and

Zr55Cu30Ni5Al10 bulk metallic glasses by differential scanning calorimetry (DSC),”


[61] Y. X. Zhuang, P. F. Xing, H. Y. Shi, J. Chen, P. W. Wang, and J. C. He, “On the heating rate

dependence of crystallization temperatures of metallic glasses,” Journal of Applied Physics,

vol. 108, no. 3, p. 033515, 2010.

[62] J. Cui, J. S. Li, J. Wang, H. C. Kou, J. C. Qiao, S. Gravier, and J. J. Blandin, “Crystallization

kinetics of Cu38Zr46Ag8Al8 bulk metallic glass in different heating conditions,” Journal of

Non-Crystalline Solids, vol. 404, pp. 7–12, 2014.

[63] D. Kawase, A. P. Tsai, A. Inoue, and T. Masumoto, “Crystallization on supercooled liquid

in metallic Zr‐Cu‐Al glasses,” Applied Physics Letters, vol. 62, no. 2, pp. 137–139, Jan. 1993.

[64] Y. L. Sun and J. Shen, “Icosahedral ordering in Cu60Zr40 metallic glass: Molecular

dynamics simulations,” Journal of Non-Crystalline Solids, vol. 355, no. 31–33, pp. 1557–1560,

2009.

[65] K. Yamaguchi, Y.-C. Song, T. Yoshida, and K. Itagaki, “Thermodynamic investigation of

the Cu–Zr system,” Journal of Alloys and Compounds, vol. 452, no. 1, pp. 73–79, Mar. 2008.

[66] A. I. Zaitsev, N. E. Zaitseva, Y. P. Alekseeva, E. M. Kuril’chenko, and S. F. Dunaev,

“Thermodynamic Properties of Melts and Phase Equilibria in the Copper–Zirconium

System,” Inorganic Materials, vol. 39, no. 8, pp. 816–825, 2003.

[67] Y. Liu, S. Hata, K. Wada, and A. Shimokohbe, “Thermal, Mechanical and Electrical

Properties of Pd-Based Thin-Film Metallic Glass,” Japanese Journal of Applied Physics, vol.

40, no. Part 1, No. 9A, pp. 5382–5388, 2001.

[68] D. H. Bae, H. K. Lim, S. H. Kim, D. H. Kim, and W. T. Kim, “Mechanical behavior of a

bulk Cu–Ti–Zr–Ni–Si–Sn metallic glass forming nano-crystal aggregate bands during

deformation in the supercooled liquid region,” Acta Materialia, vol. 50, no. 7, pp. 1749–

1759, 2002.

[69] K. Edalati and Z. Horita, “Universal Plot for Hardness Variation in Pure Metals Processed

by High-Pressure Torsion,” Materials Transactions, vol. 51, no. 5, pp. 1051–1054, 2010.

[70] J.-J. Pang, M.-J. Tan, K. M. Liew, and C. Shearwood, “Nanoindentation study of size effect

and loading rate effect on mechanical properties of a thin film metallic glass

Cu49.3Zr50.7,” Physica B: Condensed Matter, vol. 407, no. 3, pp. 340–346, 2012.

[71] Z. T. Wang, K. Y. Zeng, and Y. Li, “The correlation between glass formation and hardness

of the amorphous phase,” Scripta Materialia, vol. 65, no. 9, pp. 747–750, Nov. 2011.

[72] J. P. Chu, J. C. Huang, J. S. C. Jang, Y. C. Wang, and P. K. Liaw, “Thin film metallic

glasses: Preparations, properties, and applications,” JOM, vol. 62, no. 4, pp. 19–24, 2010.

[73] T. L. Cheung and C. H. Shek, “Thermal and mechanical properties of Cu–Zr–Al bulk

metallic glasses,” Journal of Alloys and Compounds, vol. 434–435, pp. 71–74, 2007.

[74] W. Zhang, C. Qin, X. Zhang, and A. Inoue, “Effects of additional noble elements on the

thermal stability and mechanical properties of Cu–Zr–Al bulk glassy alloys,” Materials

Science and Engineering: A, vol. 449–451, pp. 631–635, 2007.


Chapter 4

187

[75] T. Abe, M. Shimono, M. Ode, and H. Onodera, “Estimation of the glass forming ability of

the Ni–Zr and the Cu–Zr alloys,” Journal of Alloys and Compounds, vol. 434–435, pp. 152–

155, 2007.

[76] J. C. Ye, J. P. Chu, Y. C. Chen, Q. Wang, and Y. Yang, “Hardness, yield strength, and

plastic flow in thin film metallic-glass,” Journal of Applied Physics, vol. 112, no. 5, pp.

053516–053516–9, 2012.

[77] H. S. Chou, J. C. Huang, L. W. Chang, and T. G. Nieh, “Structural relaxation and

nanoindentation response in Zr–Cu–Ti amorphous thin films,” Applied Physics Letters, vol.

93, no. 19, pp. 191901–191901–3, 2008.

[78] J. C. Ye, J. Lu, Y. Yang, and P. K. Liaw, “Extraction of bulk metallic-glass yield strengths

using tapered micropillars in micro-compression experiments,” Intermetallics, vol. 18, no.

3, pp. 385–393, 2010.

[79] L. Parratt, “Surface Studies of Solids by Total Reflection of X-Rays,” Phys. Rev., vol. 95, no.

2, pp. 359–369, 1954.

[80] I. Bakonyi, “Atomic volumes and local structure of metallic glasses,” Acta Materialia, vol.

53, no. 8, pp. 2509–2520, 2005.

[81] G. Abadias, P. Djemia, and L. Belliard, “Alloying effects on the structure and elastic

properties of hard coatings based on ternary transition metal (M = Ti, Zr or Ta) nitrides,”

Surface and Coatings Technology, vol. 257, pp. 129–137, 2014.

[82] P. Djemia, F. Ganot, C. Dugautier, and M. Quilichini, “Brillouin scattering from the

icosahedral quasicrystal Al70.4Mn8.4Pd21.2,” Solid State Communications, vol. 106, no. 7,

pp. 459–461, 1998.

[83] R. Ristić, M. Stubičar, and E. Babić, “Correlation between mechanical, thermal and

electronic properties in Zr–Ni, Cu amorphous alloys,” Philosophical Magazine, vol. 87, no.

35, pp. 5629–5637, 2007.

[84] D. B. Miracle, W. S. Sanders, and O. N. Senkov, “The influence of efficient atomic packing

on the constitution of metallic glasses,” Philosophical Magazine, vol. 83, no. 20, pp. 2409–

2428, 2003.

[85] P. Yu and H. Y. Bai, “Poisson’s ratio and plasticity in CuZrAl bulk metallic glasses,”

Materials Science and Engineering: A, vol. 485, no. 1–2, pp. 1–4, 2008.

[86] D. B. Miracle, “The efficient cluster packing model – An atomic structural model for

metallic glasses,” Acta Materialia, no. 16, pp. 4317–4336, 2006.

[87] C. Zwikker, Physical properties of solid materials. New York: Interscience Publishers, 1954.

[88] H. M. Ledbetter, “Temperature behaviour of Young’s moduli of forty engineering alloys,”

Cryogenics, vol. 22, no. 12, pp. 653–656, 1982.

[89] S. F. Pugh, “XCII. Relations between the elastic moduli and the plastic properties of

polycrystalline pure metals,” Philosophical Magazine Series 7, vol. 45, no. 367, pp. 823–843,

1954.

[90] J. R. Rice and R. Thomson, “Ductile versus brittle behaviour of crystals,” Philosophical

Magazine, vol. 29, no. 1, pp. 73–97, 1974.

[91] J. A. Charles and G. C. Smith, Eds., Advances in Physical Metallurgy: The Professor Sir Alan

Cottrell Symposium, First Edition edition. London: Maney Publishing, 1990.

[92] D. G. Pettifor, “Theoretical predictions of structure and related properties of

intermetallics,” Mater. Sci. Technol., vol. 8, no. 4, pp. 345–349, 1992.

[93] J. J. Lewandowski *, W. H. Wang, and A. L. Greer, “Intrinsic plasticity or brittleness of

metallic glasses,” Philosophical Magazine Letters, vol. 85, no. 2, pp. 77–87, 2005.

[94] X. J. Gu, A. G. McDermott, S. J. Poon, and G. J. Shiflet, “Critical Poisson’s ratio for

plasticity in Fe–Mo–C–B–Ln bulk amorphous steel,” Applied Physics Letters, vol. 88, no. 21,

p. 211905, 2006.



188

[95] G. N. Greaves, A. L. Greer, R. S. Lakes, and T. Rouxel, “Poisson’s ratio and modern

materials,” Nat Mater, vol. 10, no. 11, pp. 823–837, 2011.

[96] F. Spaepen, “A microscopic mechanism for steady state inhomogeneous flow in metallic

glasses,” Acta Metallurgica, vol. 25, no. 4, pp. 407–415, 1977.

[97] P. Yu and H. Y. Bai, “Poisson’s ratio and plasticity in CuZrAl bulk metallic glasses,”

Materials Science and Engineering: A, vol. 485, no. 1–2, pp. 1–4, 2008.

[98] X. F. Liu, R. J. Wang, and W. H. Wang, “Poisson’s ratio of metallic glasses under pressure

and low temperature,” Scripta Materialia, vol. 62, no. 5, pp. 254–257, 2010.

[99] W. H. Wang, “Elastic moduli and behaviors of metallic glasses,” Journal of Non-Crystalline

Solids, vol. 351, no. 16–17, pp. 1481–1485, 2005.

[100] A. Kawashima, K. Ohmura, Y. Yokoyama, and A. Inoue, “The corrosion behaviour of Zr-

based bulk metallic glasses in 0.5 M NaCl solution,” Corrosion Science, vol. 53, no. 9, pp.

2778–2784, 2011.

[101] A. Gebert, K. Buchholz, A. Leonhard, K. Mummert, J. Eckert, and L. Schultz,

“Investigations on the electrochemical behaviour of Zr-based bulk metallic glasses,”

Materials Science and Engineering: A, vol. 267, no. 2, pp. 294–300, 1999.

[102] D. Zander, B. Heisterkamp, and I. Gallino, “Corrosion resistance of Cu–Zr–Al–Y and Zr–

Cu–Ni–Al–Nb bulk metallic glasses,” Journal of Alloys and Compounds, vol. 434–435, pp.

234–236, mai 2007.

[103] A. Kawashima, H. Habazaki, and K. Hashimoto, “Highly corrosion-resistant Ni-based

bulk amorphous alloys,” Materials Science and Engineering: A, 304–306, pp. 753–757, 2001.

[104] C.-F. Hsu, W. Kai, H.-M. Lin, C.-K. Lin, and P.-Y. Lee, “Fabrication and corrosion

behavior of Ti-based bulk metallic glass composites containing carbon nanotubes,” Journal

of Alloys and Compounds, vol. 504, Supplement 1, pp. S176–S179, 2010.

[105] C. Li, W. Chen, Q. Jiang, L. Wang, and D. Luo, “Corrosion resistance of Ti-based metallic

glass coating in concentrated nitric acid,” Materials Chemistry and Physics, vol. 143, no. 3,

pp. 900–903, 2014.

[106] G. Xie, F. Qin, S. Zhu, and D. V. Louzguine-Lugzin, “Corrosion behaviour of porous Ni-

free Ti-based bulk metallic glass produced by spark plasma sintering in Hanks’ solution,”

Intermetallics, vol. 44, pp. 55–59,2014.

[107] H. Bala and S. Szymura, “Acid corrosion of amorphous and crystalline Cu-Zr alloys,”

Applied Surface Science, vol. 35, no. 1, pp. 41–51, 1988.

[108] Y. N. Zhang, J. L. Zi, M. S. Zheng, and J. W. Zhu, “Corrosion behavior of copper with

minor alloying addition in chloride solution,” Journal of Alloys and Compounds, vol. 462,

no. 1–2, pp. 240–243, 2008.



Thin Solid Films, vol. 529, pp. 338–341, 2013.


Content

5.1 Are ceramic and metallic films compatible? ................................................................................. 189

5.2 What films do we associate? ............................................................................................................ 190

5.3 Towards an optimized physico-chemical protection through a duplex CrAlYN/Zr - 44 at% Ni

TFMG ........................................................................................................................................................ 194

Chapter summary ................................................................................................................................... 198

Chapter 5

Innovative approach in surface engineering to combine

high temperature resistance with corrosion protection

Overview

This chapter proposes an innovative way to protect the machining tools at both ambient

and high temperatures.

The corrosion resistance of the two types of PVD thin films will be discussed in light of

their chemical and structural differences.


Innovative approach in surface engineering

to combine high temperature resistance with corrosion protection

189

In the previous chapters we have presented the results related to the two different types of

PVD thin films with respect to their deposition method and associated properties.

These materials have a great potential to serve as surface engineering solutions for wide

applications. The deposition parameters as well as the aging resistance in temperatures

influence properties, as a consequence of structural thermal changes. Due to a high chemical

inertness combined with a stability of the structure, ceramic coatings are able to protect

surfaces at very high temperatures (950 °C), while the TFMGs have defects-less

morphologies and structural/chemical homogeneity, of great interest for a physico-chemical

resistance at moderate temperature.

The present chapter is willing to compare these two PVD thin films in order to determine

whether their own interesting characteristics might be combined modifying the architecture

of the composite film.

5.1 Are ceramic and metallic films compatible?

The physical, chemical and mechanical properties of the composite coating/substrate have to

be constantly optimized for an improvement of performance and lifetime of

components/parts. Adopting an approach of “surface engineering”, the ceramic and metallic

glass thin films proposed within the framework of the current thesis can represent some

potential solutions for innovative combined systems. The PVD processes have their own

applications, advantages, limitations, specificities shortly discussed in what follows.

For instance, in the case of the hard coatings, often deposited by an arc evaporation

technique, the films are well adherent with remarkable final properties. These CrN-based

coatings are used mainly for protection of machining tools against oxidation at high

temperature.

In chapter 3 we presented the most important characteristics in terms of deposition

conditions, their influence on the films’ growth and the oxidation behavior they afford to

the M2 steel to sustain high temperature environments. However, it has also been pointed

out a significant disadvantage of the process related to “poor” covering rates of the coated

surface, especially in the case of CrAlYN, which must be overcome.

The arc-deposition indeed induces formation of growth defects (droplets, open porosities)

leading to zones where the coated metal is not fully protected. Such a weak point is not only

due to the deposition process itself, but also is dependent on the targets’ nature through

their melting temperature especially. Fortunately, the results presented and discussed

previously showed and demonstrated that despite their poor surface states, arc-evaporated

ceramic coatings efficiently protected the tool at high temperatures.

In the case of the Zr-Cu TFMGs, the PVD magnetron sputtering is well adapted to obtain

defects-free thin films with still good adhesion and interesting oriented-properties.

From the chapter 4, we saw that these films have an outstanding covering rate of the

substrate and a very smooth surface. It has been established that ceramic films were very

oxidation resistant while metallic amorphous films resisted well to humid corrosion.


Chapter 5

190

The risk to combine these two conductive coatings relies on their own electrochemical

characteristics. Indeed, it is well known that partially covering nitride coatings deposited

onto steel can suffer from a severe galvanic coupling leading to a strong local attack (Figure

5.1 a, b).

Figure 5.1. CrN coated steel after a 72-hour corrosion test in a 30 g/L NaCl solution. Surface’s

aspect (a), cross section views (b ,c) [1]

A solution sometimes involved to overcome such a drawback consists in depositing, prior to

the PVD hard coating, a corrosion-resistant sublayer such as an electroless nickel coating for

instance (Figure 5.1.c).

In our case, we would like to apply the same architecture modifies strategy, associating a

hard coating as sublayer to sustain the high temperature due to the machining operation,

covered by a corrosion-resistant TFMG intended to protect the tool against the industrial

aggressive storage atmosphere.

5.2 What films do we associate?

Objective of this chapter is to determine in which conditions both metallic and ceramic

coatings may be associated, giving rise to a synergistic relationship. From the investigated

samples we have chosen the CrAlYN as the representant of hard coatings, and Zr-48.3 at.%

Cu from the TFMGs group. We considered that these films are the best ones in their

category, as it has been highlighted in the numerous sections of chapter 3 and 4. We saw

indeed that CrAlYN has the finer structure and the best oxidation resistance at high

temperatures, while the Zr-48.3 at.% Cu is the only composition with a fully amorphous

structure whatever the adopted investigation scale (XRD as well as TEM).

Besides, this film belongs to the regime B, where the disorder level is the highest leading to

high mechanical properties and to the longest supercooled region.

Before comparing these two materials, it is worth mentioning some main differences with

respect to deposition method and their properties. From a structural point of view, the hard

coatings have a crystalline structure, as compared with the amorphous glassy structure of

the Zr-Cu TFMGs.




191

Even more, the latter films are chemically and structurally homogenous, characteristics that

are required for a better corrosion protection. Their lack of grain boundaries and crystalline

structural defects (such as holes) or distortions are also valuable.

In order to determine the opportunity to associate CrAlYN with the metallic glass film, a

comparison of their potentiodynamic electrochemical curves is shown in figure 5.2.

Figure 5.2. Potentiodynamic curves for CrAlYN and Zr-48.3 at. % Cu coatings deposited onto

glass

Cathodic parts of curves for both PVD films are comparable, which implies similar

reactivity towards the oxygen’s reduction. In the anodic part, a typical long plateau is also

observed for both thin films, which means that the exposed surface becomes passivated and

highly protected. Presence of such passive film is explained by the zirconium element

contained in the metallic glass film, and by the chromium, and aluminum especially, for the

hard coating [2 - 5].

From the figure we may observe the differences in terms of corrosion potential from – 220

mV/SCE for Zr-48.3 at.% Cu, to a more noble potential for CrAlYN coating at - 110 mV/SCE.

The shift of the potential of the film to the positive direction implies a still better corrosion

resistance for the nitride. This assumption is confirmed further by the corrosion current

densities, one decade lower for the ceramic film with respect to the metallic one ( 1.1 µA/cm2

for Zr-48.3 at.% Cu and 0.14 µA/cm2 for CrAlYN).

However, considering the Tafel’s plots the passive film formed on Zr-48.3 at.% Cu TFMG

(ZrO2-rich) seems to be more stable than the one for CrAlYN hard coating (Al2O3-rich), since

the passive plateau is larger and almost constant at roughly 10 µA/cm2, whereas the

quaternary coating evidences a depassivation from 0.8 V/SCE.

The depassivation of the ceramic film’s surface could be due to formation of the Cr2O3 oxide

which has a less compact structure [6]. The growth of this second oxide may cause the

rupture of the diffusion barrier, leading to a porous morphology and paths for the diffusion

of corrosive species. The SEM micrographs presented in figure 5.3. allow corroborating such

behavior, since some cracks are observed on the exposed area.


Chapter 5

192

Figure 5.3. SEM micrographs of CrAlYN after corrosion test in Na2SO4. Inserted figures: the A

represents the exposed area, in B is highlighted an example of the initiated cracks due to

corroded surface and C is showing the surface of the film in the unexposed zone to electrolyte

From the figure we may observe, first, that the depassivation recorded on the

potentiodynamic curve (Figure 5.2.) could be due the surface cracking (insert B) and

instability of the passive film. We assume that once the Cr3O2 is formed, its porous structure

may lead to the breakthrough of the Al2O3 protective layer and therefore cracks can initiate

on the surface.

In a second level, it can be observed also a change of the surface’s morphology when

comparing corroded zones with respect to unexposed ones (insert A vs C). The modification

of the surface roughness is due to corrosion products. Their formation favors the

development of stresses, still inducing a cracks’ propagation, contributing to the earlier

depassivation.

In opposite, regarding the Zr-48.3 at.% Cu TFMG, the passive film is more stable and

resistant. Moreover, corrosion test does not significantly affect the exposed surface, as


Figure 5.4. SEM micrographs of Zr-48.3 at.% Cu TFMG after corrosion test in Na2SO4. Inserted

figures: the A represents the exposed area, and B is showing the surface of the film in the

unexposed zone to electrolyte




193

The morphology of the Zr-Cu TFMG is indeed not changing much after exposure to the

Na2SO4 electrolyte. From the insert A we may observe that the exposed zone is very slightly

affected, without any corrosion products and continuity defect. This confirms again the high

stability of the passive Zr-rich film formed on the surface.

In our vision we proposed that Zr-Cu TFMG could be deposited on the already coated tool

with CrAlYN coating and forming therefore a bilayer coating. The metallic glass film has to

behave as a sacrificial layer once the machining tool is used at high temperatures.

When stored, the glassy film will be responsible for the corrosion protection of the tool,

while once exposed at high temperatures it will rapidly crystallize and then probably will be

fast destroyed.

We consider that such easy-to-do “hybrid” approach would directly increase the lifetime of

the protected surfaces at room temperature as well as when exposed at high temperatures.

Such type of stratification approach has already been involved for some applications. For

instance, the sacrificial thin films are commonly used for the antibacterial protection [7], in

case of corrosion resistance improvement [8, 9] and even for polymers [10].

Very recently, Fenker and colleagues reported a complex review related to the corrosion

protection of the hard coatings [11] where they presented among other results, a study of

Purushotham et al. [12]. In this paper, authors described effects of a Zr-implantation of PVD

TiN films on their corrosion behavior. The polarization curves of the Zr-implanted TiN-

coated mild steel showed a decreased of the corrosion current density icorr with a shift of the

corrosion potential Ecorr to more noble values, thus implying a lowered reactivity. The

protection was attributed to a crystal structural modification and formation of an oxinitride,

transformation that was able to delay access of the electrolyte to the substrate.

It is actually clear the high interest hold by researchers to new and innovative solutions to

improve the corrosion resistance of the hard coatings developed for high temperature

applications. Our proposed multilayer coating could be an innovative solution for

corrosion-less of the CrAlYN as schematically represented in the figure 5.5.

Figure 5.5. Schematic representation of the CrAlYN ceramic film deposited on tool steel without

(a) and with a further metallic glass top-coat (b)


Chapter 5

194

The figure shows that, ideally, the metallic glass film once deposited on the top will fulfill

and hide the holes and growth defects of the subjacent coating. The interlayer zone will

have probably a mixed composition, but at the deposition temperature no significant

diffusion into the CrAlYN should probably be generated during the second deposition of

the metallic layer. We believe that this solution could be interesting to be tested especially

for an enhanced corrosion protection purpose. Indeed, in terms of mechanical properties,

neither the CrAlYN properties, nor its oxidation resistance at high temperatures will be

affected.

It can then be concluded that CrAlYN may be used as the oxidation protective underlayer,

while the metallic glass coating could be exploited for its ability to passivation.

However, such association leads to two weak points. First, if the Zr-48.3 at.% TFMG has the

ability to passivate, its passive film reduces only moderately the dissolution rate estimated

around 1 µA/cm2. Moreover, both samples show a difference of potential of about 110 mV

(Figure 5.2), which is high enough to consider a galvanic corrosion. Such a galvanic

coupling will induce a preferential sacrificial corrosion of the metallic top-coat, leaving then

the nitride sublayer unprotected.

For a transitory protection, such association could be acceptable, but more efficient

alternatives can be found regarding the nature of the metallic layer.

5.3 Towards an optimized physico-chemical protection through a duplex

CrAlYN/Zr - 44 at% Ni TFMG

It is important to optimize the protection using another TFMG susceptible to act as a

potential sacrificial layer.

In the previous figure 5.2. we noticed that the Zr-48.3 at.% Cu TFMG evidenced a high

stability of the protective ZrO2 passive layer, even if the copper does not passive.

The query would be: what can happen if the copper is replaced in binary matrix by another

corrosion-resistant metal like Ni? Will the corrosion response of the Zr-based binary TFMG

be affected?

These are some questions on which we will focus on the next section of this chapter.

Considering the experience of the LERMPS laboratory in deposition methodology and the

corrosion expertise from CorrIS laboratory, we decided to look further into this issue and

we come across with several other points that deserved to be mentioned.

A large number of electrochemical studies have been reported about the superior corrosion

resistance of nickel with respect to copper or iron due to its higher passivation

characteristics [13 -15].

In our continuous attempt to improve the corrosion and structural stability of the Zr-based

binary TFMGs we decided to deposit a Zr-Ni TFMG with a similar composition as Zr-Cu

TFMG (Zr- 44 at.% Ni) and to compare its electrochemical behavior with the CrAlYN

coating.

Some recent studies related to Zr-Ni TFMGs (few compositions) proposed that the binary

Zr-based deposited by PVD magnetron sputtering have good structural and mechanical

properties, but the authors did not consider also the corrosion resistance [16, 17].




195

Our intention is not to initiate a new and complex investigation on this system, but only to

see whether such a Zr-Ni system could be an efficient solution for replacing the Zr-Cu

TFMG for the protection against corrosion.

One of the prerequisite is to verify the amorphous nature of the binary Zr-Ni film. Figure

5.6. presents the XRD pattern for the as-deposited Zr-44 at.% Ni TFMG.

Figure 5.6. XRD pattern of as-deposited Zr- 44 at.% Ni TFMG

The XRD pattern exhibits no Bragg diffraction peaks but only a wide hump similar to that

already displayed in the case of Zr-Cu coatings. The Zr- 44 at.% Ni film can then be

considered as a TFMG.

As stated earlier, this thin film has been prepared with the intention to protect the nitride

coated tool against corrosion. Potentiodynamic curves were then drawn in figure 5.7 for

both metallic and ceramic coatings.

Figure 5.7. Potentiodynamic curves for the CrAlYN coating and the Zr-44 at. % Ni

From the figure we may notice that the electrochemical potential for both electrodes is very

similar, avoiding occurrence of a harmful galvanic coupling between both stratified

structures. Moreover, the passive ability afforded by the zirconium element is also

preserved with the nickel.

Zr- 44 at.% Ni

20 25 30 35 40 45 50 55 60 65 70 75

2deg

Inte

ns

ity

/a

rb.

un

its


Chapter 5

196

All corrosion characteristics are summarized in table 5.1. It appears that the corrosion

behavior due to the Ni addition is far better than that obtained with the Cu alloying (higher

potential values and corrosion rates two decades lower). Ni-containing film affords a high

protection and could be a better alternative than Zr-Cu TFMG for the bilayer system

proposed in figure 5.5.

Table 5.1. Comparison between the corrosion characteristics of Zr-based TFMGs and CrAlYN

Film

Corrosion

potential

(Ecorr) mV/SCE

Corrosion

current

(icorr) µA/cm2

Zr-44 at.% Ni -78 0.02

CrAlYN -110 0.14

Zr-48.3 at.% Cu -220 1.10

If we consider further the electrochemical impedance spectrometry (EIS) the Zr-Ni TFMG as

compared with CrAlYN has a much higher capacitive loop, as illustrated in figure 5.8.

Figure 5.9. Nyquist representation for amorphous Zr- 44 at.% Ni TFMG and crystalline CrAlYN

coating

From Nyquist diagrams the capacitive loops are not well separated, but even so, the

superiority of the Zr-Ni TFMG at lower frequencies and beyond is obvious. The smaller

capacitive loop for the CrAlYN may imply a much higher global kinetics of reactivity of the

surface due probably to its crystalline state. The differences between these two PVD films

are notable and infer that the Zr-Ni TFMG could be an ideal solution for protection against

corrosion.

Therefore, we are able to sustain the choice of the Zr- 44 at.% Ni TFMG to protect the surface

of the CrAlYN coating as a part of a multilayer film. When we were directly compared the

Cu-containing films with the ceramic film, we saw that there is still place for improvements,

hence our second proposed solution with the Ni-containing TFMG could be better for

corrosion protection.




197

The combination of these two thin films (ceramic vs metallic glass) in bilayer architecture,

combining the oxidation resistance of the sublayer with the corrosion protection of the top-

coat could be an efficient answer to the industry needs, in continuous search for enhanced

oriented-properties.

Such multilayer could be deposited in a single run using a hybrid deposition system (arc-

PVD/magnetron sputtering), or by means of recent HIPIMS process presenting the

advantages of arc-PVD in terms of ionized vapor’s level but without its drawbacks linked to

droplets’ emission. Another solution could be to proceed in two steps in two separate

deposition chambers involving the two most suitable process: ion plating for ceramic films,

sputtering for metallic ones.

If the latter solution may appear questionable from a commercial viewpoint, it nevertheless

allows the industrial partner a better flexibility proposing, or not, to the nitride coated tool

further corrosion protectiveness, susceptible to increase thereby the added value of the part.

In the commercial field, it can be underlined the favorable aesthetic impact that can be got to

the coated part after the further deposition of a bright colored layer to recover a mat shadow

AlN-rich coating.


Chapter 5

198

Chapter summary

Chapters 3 and 4 were devoted to present characteristics of the two types of films

investigated within the framework of this PhD. Hard ceramic films were identified to well

sustain high temperature oxidation and an improved resistance for highly severe

environments was specifically found for the quaternary CrAlYN coating.

A second class of metallic layers was also studied, obtained by magnetron sputtering, and

characterized by an amorphous homogeneous structure. Thanks to the propensity of

zirconium to form a surface protective passive film, reactivity of these TFMGs is very low, in

particular for the film Zr-44 at.% Ni.

Adopting a surface engineering approach, we have proposed to change the architecture of

the coating in order to take benefit of the high temperature resistance of the nitride based

sublayer, in addition to the efficient passivation of the metallic glass top-coat.

Electrochemical properties of such “composite” layer are simulated and have showed its

efficiency with tiny corrosion rates and suppression of the galvanic effect involved between

both layers.

Furthermore, such a stratified structure may enable filling pores and hide the harmful

droplets of the subjacent arc-evaporated coating. The metallic smooth top-coat may also

give to the part an aesthetic interesting aspect, of great potential importance from a

commercial viewpoint.




199

Lists of references

[1] P. Steyer, S. Anderbouhr, and P. Jacquot, “Effect of Ni Interlayer on Corrosion Behaviour

of Crn Pvd Coating Deposited onto Plastic Mould Steels,” Surface Engineering, vol. 17, no.

4, pp. 327–331, 2001.


reactive direct current magnetron sputtered CrAlN and CrN coatings,” Surface and

Coatings Technology, vol. 201, no. 6, pp. 2193–2201, 2006.



Thin Solid Films, 529, 338-341, 2013

[4] C. W. Chu, J. S. C. Jang, S. M. Chiu, and J. P. Chu, “Study of the characteristics and

corrosion behavior for the Zr-based metallic glass thin film fabricated by pulse magnetron

sputtering process,” Thin Solid Films, 517, 17, pp. 4930–4933, 2009.

[5] B. Subramanian, S. Yugeswaran, A. Kobayashi, and M. Jayachandran, “Fabrication of

amorphous Zr48Cu36Al8Ag8 thin films by ion beam sputtering and their corrosion

behavior in SBF for bio implants,” Journal of Alloys and Compounds, 572, pp. 163–169, 2013.

[6] H.-Y. Chen and F.-H. Lu, “Oxidation behavior of chromium nitride films,” Thin Solid

Films, vol. 515, no. 4, pp. 2179–2184, décembre 2006.

[7] M. Köller, C. Sengstock, Y. Motemani, C. Khare, P. J. S. Buenconsejo, J. Geukes, T. A.

Schildhauer, and A. Ludwig, “Antibacterial activity of microstructured Ag/Au sacrificial

anode thin films,” Materials Science and Engineering: C, vol. 46, pp. 276–280, Jan. 2015.

[8] C. H. Chang, Y. Rheem, Y.-H. Choa, D.-Y. Park, and N. V. Myung, “Galvanic

displacement of BixTey thin films from sacrificial iron group thin films,” Electrochimica

Acta, vol. 55, no. 3, pp. 1072–1080, 2010.

[9] D. Elazem, H. Jung, T. Wu, J.-H. Lim, K.-H. Lee, and N. V. Myung, “Morphology change

of galvanically displaced one-dimensional tellurium nanostructures via controlling the

microstructure of sacrificial Ni thin films,” Electrochimica Acta, vol. 106, pp. 447–452, 2013.

[10] D. Reuter, A. Bertz, M. Nowack, and T. Gessner, “Thin film encapsulation technology for

harms using sacrificial CF-polymer,” Sensors and Actuators A: Physical, 145, 316–322, 2008.

[11] M. Fenker, M. Balzer, and H. Kappl, “Corrosion protection with hard coatings on steel:

Past approaches and current research efforts,” Surface and Coatings Technology, vol. 257,

pp. 182–205, 2014.

[12] K. P. Purushotham, L. P. Ward, N. Brack, P. J. Pigram, P. Evans, H. Noorman, and R. R.

Manory, “The effect of MEVVA ion implantation of Zr on the corrosion behaviour of PVD

TiN coatings,” Corrosion Science, vol. 50, no. 1, pp. 8–14, 2008.

[13] E. Sikora and D. D. Macdonald, “Nature of the passive film on nickel,” Electrochimica Acta,

vol. 48, no. 1, pp. 69–77, 2002.

[14] J. L. Trompette, L. Massot, and H. Vergnes, “Influence of the oxyanion nature of the

electrolyte on the corrosion/passivation behaviour of nickel,” Corrosion Science, 74, pp.

187–193, 2013.

[15] B. Ter-Ovanessian, C. Alemany-Dumont, and B. Normand, “Single frequency

electrochemical impedance investigation of zero charge potential for different surface

states of Cu–Ni alloys,” J Appl Electrochem, vol. 44, no. 3, pp. 399–410, 2014.






Alloys Compd. 615, S348-S351, 2014


200

General conclusions

The prime aim of this thesis was the thermal investigation of two different thin films

systems, namely ceramic and metallic glass, produced by PVD techniques. The focus was on

their structural and functional stability, adopting a specific multi-scale approach with the

purpose to improve their in-service physico-chemical durability.

On this basis, CrN-based coatings and Zr-Cu thin film metallic glasses (TFMGs) were

deposited on several different substrates and a detailed characterization of their most

relevant properties was carried out.

The first chapter was dedicated to the state-of-the-art on the two categories of coatings.

It appeared that it is a topical subject with many recent papers devoted to how to synthesize

these surfaces and for which applications? However, both films need to be more basically

investigated in order to optimize their characteristics and performance, objective of the PhD.

The results from first part were devoted to the hard CrN-based coatings. The aim was to

determine the role of both Al and Y alloying elements on the oxidation behaviors.

In addition to conventional post mortem thermogravimetric measurements, the oxidation of

the films has also been followed by means of real time approach using ESEM and in situ

XRD. During those experiments we were able to observe the chemical reactivity of the films’

surfaces and the structural transformations that undergo during exposure in air.

Deposited by arc-evaporation, films are columnar and suffered from a large amount of

growth defects such as droplets and open-pores, especially for the Al-containing coatings.

Such a poor morphology affected the oxidation resistance at moderate temperatures, for

which CrN, the most covering coating, evidenced a high protectiveness.

The chemical and structural analyses have showed that aluminum was inserted into a solid

solution under stresses and with a distortion of the lattice. For these films, the protective

scale formed on the surface is composed by a mixed (Al, Cr)2O3 oxide. This well-adherent

oxide efficiently protects the subjacent coating for temperatures as high as 900 C,

preserving the initial columnar structure. A further addition of small amounts of yttrium

(about 1 at. %) into the CrAlN film improves even more the oxidation resistance for very

severe conditions (950°C).

The main explanation is related to the finer structure and a chemical effect which accelerates

the global reactivity of the species by forming faster the oxide protective scale on the

coating’s surfaces. According to us, the oxidation mechanism would be governed by the

nitrogen release and outward diffusion process of the species from the film to form the

oxide scale. The beneficial effect of Al would then be due to a much higher thermodynamic

stability of AlN versus CrN, making more difficult its oxidation.

Moreover, the further positive influence of yttrium would be the consequence of a still

better stability due to the Y insertion into the CrAlN lattice, and/or to the formation of a

more protective Al enriched mixed oxide.


201

On the second results part, the manuscript is focused on the Zr-Cu TFMGs deposited by

magnetron sputtering using pure Zr and Cu targets in a wide compositional range covering

the entire phase diagram. The metallic glass thin films represent a new class of materials,

since the properties of the more massive and classical bulk metallic glasses can be today

effectively transferred at a micrometer level. The metallic glass thin films are newcomers to

the field of amorphous materials, and the main advantage is that during deposition it is

easier to obtain a random atomic and bonding configuration as compared with classical

quenching methods.

The general idea within the work was basically the development of the Zr - Cu thin films

and to combine their excellent structural and chemical homogeneity with the mechanical

and physic-chemical properties. For this part, the main objective was to better understand

the possibility of amorphization/crystallization and to determine the influence of structural

state on the films properties.

Regarding the most relevant conclusion, it might be worth mentioning that the films are

amorphous in a wider compositional domain (33.3 – 89.1 at.% Cu) compared with classical

BMGs with three different regions and maximum structural stability in the 45 – 75 at. % Cu

range where the films have the best properties.

These different regions were explained taking into account their properties such as

mechanical properties or structural stability determined from the crystallization resistance,

where a microstructural evolution model was proposed for each regime. The combination of

composition, amorphous structure and thermal characteristics was essential to understand

the particular features of the crystallization resistance of the films. The results revealed a

strong dependence on composition and deposition parameters. Thanks to their outstanding

surface morphology in terms of smoothness, homogeneity and grain-boundaries absence,

electrochemical properties of these films was high, leading to a promising corrosion

resistance.

Conclusions of the results chapters 3 and 4 have led to the identification of the best ceramic

and metallic glass candidates. CrAlYN was indeed the most oxidation-protective coating,

while Zr- 48 at.% Cu resist very well to corrosion. The idea developed in chapter 5 consisted

to simulate a composite architecture where the highly defective hard surface would be

protected against corrosion by the smooth perfect metallic glass top-coat. The concept is

validated from an electrochemical point of view.

Besides, another metallic glass system Zr-Ni is proposed, able to avoid the harmful galvanic

coupling involved between the ceramic under-layer and the metallic outer film.

Finally, the results obtained during the course of the present work showed that in the case

of CrN-based coatings, especially the CrAlYN type, the oxidation resistance is significantly

improved which will increase the lifetime of the machining tools utilized for industrial

applications.


202

In what concerns the metallic glass thin films, the continuous research will increase the

possibility to use such materials for diverse applications. The lack of crystalline defects is

one of the main plus of these materials, which will fast propose them as structural materials

and models to develop new applications. The experiments conducted on the Zr-Cu TFMGs

are believed to be transferable also to other binary TFMGs. Their chemical simplicity should

qualify them as very good subjects for theoretical simulation and modeling.

A duplex architecture would be susceptible to combine both beneficial effect of each layer,

the ceramic sublayer affording the high temperature resistance while the metallic glass top-

coat acting as a corrosion sacrificial layer.


203

Outlook and perspectives

The use of thin films in engineering applications relies on the ability of these systems to

meet the design and service requirements to the final properties. The capability of the thin

films to meet these requirements is generally determined by the mechanical and physico-

chemical properties.

In spite of new conclusions reached through this thesis several aspects for future

investigations still remain open, as will discussed in the following.

Regarding the CrN-based hard coatings, we can suggest some future research opportunities:

The mechanical performance of the coatings deposited by arc evaporation is deeply

influenced by their structure, microstructure, phase’s transformations and deposition

defects. One downside of the arc-PVD process is the important amount of growth defects

compared to other classical PVD methods involving the sputtering. Even so, the final

properties are most of the time just slightly influenced by these defects and such coatings

are successfully applied for protection against oxidation of machining tools.

The deposition process could be optimized in a further work to reduce the droplets level

and open porosities, especially in the case of CrAlYN film. A possible solution could be a

filter system positioned in the front of cathodes used as magnetic field to separate the

droplets from the coating flux. Furthermore, new opportunities can be found with the use of

innovative processes such as HiPIMS allowing, as in the case of arc-PVD, a high amount of

energetic cations leading to a good adherence, with less emission of droplets.

Among the CrN-based coatings we demonstrated that CrAlYN had the best

oxidation behavior at high temperatures. The yttrium accelerates the chemical reactivity of

the metal elements and boost up the formation of a well-protective oxide layer on the top

surface that seems hindering the inward diffusion of oxygen. The mechanism that activates

this reactivity is still unclear and further studies are required. The oxidation mechanism of

quaternary CrAlYN coating should be better understood during a study involving different

yttrium contents and different oxidation periods. Based on such new findings an oxidation

model showing the kinetics and influence of yttrium would be likely proposed.

The ceramic coatings studied here showed good mechanical properties and high

oxidation resistance. Exploring further the tribological properties and cracking mechanism

could provide a more detailed image on their operation performance. Application of those

coatings on real cutting tools and test them could be a next further step.


204

Regarding the new class of thin films we concentrated our study on the phase selection

mechanism with relevant phase competition and the associated dynamics of structural

evolution of devitrification in the Zr-Cu system. From these findings some future

suggestions can be proposed as follows.

Thermodynamics and crystallization model proposed can be used to predict

observed phase sequences and transition rates. At the same time, in devitrification studies, it

is important to study the initial stages of phase decomposition leading to crystal formation

as well. In the present study we provided the sequence of phases during crystallization

under continuous heating for all Zr-Cu phase diagram. Hence, using these informations,

such as time and temperature, will permit to model the crystallization stages for other

similar binary TFMGs, such as Zr-Ni alloys especially, that were identified as an interesting

corrosion system.

The devitrification analysis results can be used in controlling the microstructure to

enhance the desired properties of metallic glass thin films. The control of nanocrystalline

domains could be the key for obtaining thin film metallic glasses with higher mechanical

and physical properties. In this sense, post-controlled heating treatments may be involved

with the purpose to manage the nanograins growth, leading to optimized properties.

The finding of relationships between the structural relaxation and the local atomic ordering

in supercooled liquid region should provide the control of glassy structure.

Hence, the behavior of structural relaxation for TFMGs using different annealing time and

temperature below glass transition temperature could be studied in further projects and

compared with other classical BMGs.

The good elastic properties of Zr-Cu TFMGs are a plus of the metallic glass-material.

Their glassy structure can endure greater stresses during mechanical testing; thereby the

plastic deformation energy accumulated before final rupture is higher compared to

crystalline counterparts. It could be interesting to develop a model to assess the formation of

shear bands during plastic deformation and to observe their distribution during tensile or

compressive tests by means of in situ real time observation inside the SEM chamber.

These results could feed a database to be compared with the elastic behavior determined by

photo-acoustic measurements.

Such materials are ideal for protection against corrosion due to their structural and

chemical homogeneity. For the particular Zr-Cu TFMGs studied here, it could be interesting,

first, to better understand and explain the copper influence considering its non-passivating

nature and secondly, to explore if a “semi-amorphous”/crystalline structure influence the

corrosion behavior. Will the creation of partial ordering or full crystalline structure affect or

improve the corrosion resistance? This could be a question to be answered within the

framework of further investigations.


205

Efficiency of the duplex ceramic/metallic architecture susceptible to combine both oxidation

and corrosion resistances have been validated by electrochemical simulated coupling.

For a more realistic configuration, it would be interested to test directly this opportunity

into the latest generation of industrial hybrid arc evaporation/magnetron sputtering reactor.

The challenge will be to overcome the poor rough morphology of the hard sublayer, in

order to reach a good adherence.

Finally, from a pragmatic point of view, the results presented in the current thesis showed

that the thin films are a real surface engineering solution to extend the lifetime of the coated

substrates. The films studied here could be proposed for different types of applications.

For instance, when the temperature domain is below 400 C the Zr-Cu TFMGs could be the

right choice for engineering applications. Their main advantage is the defects-free structure

and chemical homogeneity, which could make them exciting candidates for microelectronics

industry, aesthetic optical applications and the corrosion protection.

On the other hand, when the temperatures are higher than 500 C and the utilization

environment is highly oxidative the ceramic coatings are the better choice due to higher

structural stability and oxidation resistance. Such coatings are preponderantly suggested to

be used for protection of machining tools against oxidation.


Documents

Temperature impact on thermal evolution of advanced PVD ceramic