Simulation of Diffusion Processes in Turbine Blades
and
Large Area Deposition of MAX Phase Thin Films with PVD
Von der Fakultät für Georessourcen und Materialtechnik der
Rheinisch-Westfälischen Technischen Hochschule Aachen
zur Erlangung es akademischen Grades eines
Doktors der Ingenieurwissenschaften
genehmigte Dissertation
vorgelegt von Diplom-Ingenieurin
Claudia Walter
aus Moers
Berichter:
Univ.-Prof. Jochen M. Schneider, Ph. D.
Univ.-Prof. Dr.-Ing. Lorenz Singheiser
Tag der mündlichen Prüfung: 28. Juni 2005
Danksagung
Die vorliegende Arbeit entstand während meiner Tätigkeit am Lehrstuhl für Werkstoffchemie der Rheinisch-Westfälischen Technischen Hochschule Aachen. Für die
Förderung durch die Deutsche Forschungsgemeinschaft im Rahmen des Sonderforschungsbereiches 370: “Integrative Werkstoffmodellierung” möchte ich mich an
dieser Stelle bedanken.
Ich bedanke mich bei Herrn Prof. Schneider für die Möglichkeit an etwas zu arbeiten,
was mir wirklich Spass gemacht hat, die sehr gute Betreuung und die angenehme
Arbeitsatmosphäre.
Herrn Prof. Singheiser danke ich herzlich für die Übernahme des Korreferats.
Vielen Dank an alle Kollegen des Lehrstuhls und innerhalb des Sonderforschungsbereiches, die durch ihre Kreativität, Erfahrung und Kooperation zum Gelingen meiner
Arbeit beigetragen haben.
Mein besonderer Dank gilt Peter Franke und Bengt Hallstedt für die unerschockene
Betreuung auch der seltsamsten Simulationsprojekte, und Reimund Swoma für
Engagement über das Selbstverständliche hinaus bei technischen Problemen und
kontroversen Diskussionen.
Ein Hoch auf die Werkstätten, die auch das Unmögliche möglich gemacht haben!
Ich danke meiner Familie und ganz besonders meinen Eltern für Liebe, Geduld und
Unterstützung in allen Lebenslagen.
El beso más grande a Carlos por su apoyo y amor.
II
Contents
I
Simulation of Diffusion Processes in Turbine Blades
1
1 Introduction
1.1 Motivation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
3
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2 State of the Art
2.1 Turbine Blades . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
2.2 Microstructure Simulation of Turbine Blades . . . . . . . . . . . . . . .
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3 Simulation Methods
3.1 Equilibrium Calculations . . . . .
3.2 Scheil Calculations . . . . . . . .
3.3 Diffusion Calculations . . . . . .
3.3.1 Moving Boundary Systems
3.3.2 Dispersed Systems . . . .
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4 Results and Discussion
4.1 Solidification and Cooling of CMSX-4 . . . . . . . . .
4.1.1 Fraction Solid . . . . . . . . . . . . . . . . . .
4.1.2 Microsegregation . . . . . . . . . . . . . . . .
4.1.3 Solid State Transformation . . . . . . . . . . .
4.2 β Phase Depletion in the Bond Coat . . . . . . . . .
4.3 Interdiffusion between Bond Coat and Base Material
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5 Summary and Conclusions
35
II Towards Large Area Deposition of MAX Phase Thin
Films with PVD
37
6 Introduction and Motivation
6.1 MAX phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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7 Thin Film Growth
7.1 Magnetron Sputtering . . . . . . . . . . . . . . . . . . . . . . . . . . .
7.2 Experimental Setup . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
43
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III
8 Characterization Methods
8.1 XRD Analysis of the Crystal Structure . . . . . . . . . . . . . . . . . .
8.2 XPS and WDX Analysis of the Chemical Composition . . . . . . . . .
8.3 Nanoindentation Measurements of the Mechanical Properties . . . . . .
47
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49
49
9 Results and Discussion
9.1 Ti2 AlC . . . . . . . . . . . . . . . . . . . . . .
9.1.1 Compound Target . . . . . . . . . . .
9.1.2 Compound Target + Ti Target . . . .
9.1.3 Single Target of Ti and Ti2 AlC Wedges
9.1.4 Compound Target with Cathodic Arc .
9.2 Cr2 AlC . . . . . . . . . . . . . . . . . . . . . .
9.2.1 Compound Target . . . . . . . . . . .
9.3 Ti2 AlN . . . . . . . . . . . . . . . . . . . . . .
9.3.1 Ti Target + Al Target + N2 Gas . . .
9.3.2 Powder Target . . . . . . . . . . . . .
9.3.3 Ti Target + AlN Target + N2 Gas . .
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10 Summary and Conclusions
75
11 Future Work
77
IV
List of Figures
2.1
Material system of a turbine blade. . . . . . . . . . . . . . . . . . . . .
3.1
3.2
3.3
3.4
3.5
Schematic drawing of the principle of a Scheil calculation. .
Simulation setup for a moving phase boundary simulation.
Schematic of a moving phase boundary simulation. . . . .
Simulation setup for a dispersed system. . . . . . . . . . .
Schematic of a dispersed system simulation. . . . . . . . .
4.1
4.2
4.3
4.4
4.5
4.6
4.7
4.8
4.9
4.10
4.11
4.12
4.13
4.14
4.15
4.16
4.17
4.18
4.19
4.20
Micrograph of a transverse section of a quenched sample. . . . . . . . .
Comparison of fraction solid from different calculation methods. . . . .
Solute distribution plotted vs. distance for Co, W, and Re. . . . . . . .
Solute distribution plotted vs. distance for Ti, Al, and Cr. . . . . . . .
Solute distribution plotted vs. distance for Ta. . . . . . . . . . . . . . .
Calculated mole fraction of γ’ vs. distance. . . . . . . . . . . . . . . . .
Micrograph of a horizontal section of quenched CMSX-4. . . . . . . . .
Micrograph of a horizontal section of quenched CMSX-4. . . . . . . . .
Microstructure of the bond coat after high temperature exposure. . . .
Measured oxidation kinetics for the MCrAlY bond coat. . . . . . . . .
Simulated β depletion for different replacements of the oxidized Al. . .
Simulated mole fraction of β at 1223 K. . . . . . . . . . . . . . . . . . .
Simulated mole fraction of β at 1273 K. . . . . . . . . . . . . . . . . . .
Simulated mole fraction of β at 1323 K. . . . . . . . . . . . . . . . . . .
Micrograph of the bond coat at different times of heat treatment at 1323 K.
Simulated mole fraction of β at 1323 K with a factor 1.5. . . . . . . . .
Initial concentration profiles. . . . . . . . . . . . . . . . . . . . . . . . .
Concentration profiles after annealing treatment at 1273 K. . . . . . . .
Concentration profiles after annealing treatment at 1373 K. . . . . . . .
Concentration profiles after annealing treatment at 1473 K. . . . . . . .
15
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6.1
6.2
6.3
6.4
Periodic table with elements forming MAX phases highlighted.
MAX phases reported in the literature [38]. . . . . . . . . . .
Unit cell of a 211 MAX phase. . . . . . . . . . . . . . . . . . .
Charge-density contours for a) TiC, b) TiAl, and c) Ti2 AlC. .
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7.1
7.2
Schematic of magnetron sputtering. . . . . . . . . . . . . . . . . . . . .
Structure zone model proposed by Movchan and Demchishin [48]. . . .
44
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V
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7.3
7.4
Structure zone model proposed by Thornton [49]. . . . . . . . . . . . .
Arrangement of cathodes and substrate in the vacuum chamber. . . . .
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8.1
8.2
8.3
Illustration of Bragg’s law. . . . . . . . . . . . . . . . . . . . . . . . . .
X-ray diffractometer. . . . . . . . . . . . . . . . . . . . . . . . . . . . .
Cross section of an indent before, during, and after indentation. . . . .
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9.1
9.2
9.3
9.4
9.5
9.6
9.7
9.8
9.9
9.10
9.11
9.12
9.13
9.14
9.15
9.16
9.17
9.18
9.19
9.20
9.21
9.22
9.23
9.24
9.25
XRD of the Ti2 AlC compound target material. . . . . . . . . . . . . . .
WDX data of the Ti2 AlC compound target material. . . . . . . . . . .
XRD of thin films deposited with the Ti2 AlC target at 400 K and 1153 K.
Ti/Al ratio of thin films deposited at different Ar-pressures. . . . . . .
Ti/Al ratio of thin films deposited at different Ti target power densities.
XRD of a thin film deposited with optimized parameters. . . . . . . . .
Ti/Al ratio of thin films deposited at different temperatures. . . . . . .
Segmented target consisting of a variable number of Ti and Ti2 AlC wedges.
Position of the metallic Ti wedges in the segmented target. . . . . . . .
XRD for thin films deposited with a variable number of Ti wedges. . .
XRD for a film at a substrate temperature of 1173 K with 4 Ti wedges.
XRD for films at different temperatures with the segmented target. . .
Ti2 AlC cathode for the cathodic arc experiment. . . . . . . . . . . . . .
Eroded Ti2 AlC cathode after the cathodic arc experiment. . . . . . . .
Cr2 AlC compound target. . . . . . . . . . . . . . . . . . . . . . . . . .
XRD of the Cr2 AlC compound target material. . . . . . . . . . . . . .
XRD of a thin film deposited with the Cr2 AlC compound target at 1123 K.
XRD of thin films deposited at different substrate temperatures (Cr2 AlC).
Effect of N2 poisoning on the target voltage. . . . . . . . . . . . . . . .
Target with AlN powder and ring construction. . . . . . . . . . . . . .
XRD of a thin film deposited with the AlN powder target. . . . . . . .
XRD of a thin film deposited with the sintered AlN target. . . . . . . .
XRD showing first traces of the Ti2 AlN MAX phase. . . . . . . . . . .
XRD of a thin film deposited with optimized parameters (Ti2 AlN). . .
XRD with high spatial resolution of a Ti2 AlN film. . . . . . . . . . . .
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VI
Part I
Simulation of Diffusion Processes in
Turbine Blades
1
Chapter 1
Introduction
Today metallic materials have been known and used for thousands of years. However,
most of the knowledge about metals has been created only in the last century. Only
recently it has been understood that the macroscopic properties of a material are determined by its microstructure. Most properties can be varied within a wide range by
changing the microstructure of a material with a constant overall chemical composition. Additionally it was found that the microstructure undergoes changes during the
processing of a material. Together these relations are nowadays employed to produce
materials with tailor made properties through optimized processing.
Another important development, especially in the last 40 years of the last century,
was the vast increase of computational power. Computers have become an important
tool in materials science. While in earlier times knowledge could be acquired only
via experiments, it is nowadays attempted to replace some of these experiments by
computer simulations. This has the advantage of saving time and money. Moreover,
simulations can give insights into areas which are not accessible experimentally. If they
are based on physical models they can also contribute to a deeper understanding of
the basic physics behind a problem.
It follows directly from these developments that the simulation of the microstructure
during processing of a material is a major topic in materials science today. This work
contributes to this effort for selected processes in the life-cycle of a turbine blade. A
turbine blade is exposed to high mechanical loads at high temperatures which requires
a highly developed material system. An ordinary turbine blade today consists of a
base material of excellent high temperature strength coated with several protective
layers. The manufacturing involves many different processing techniques. As base
material either a single crystalline or a directionally solidified Ni-based superalloy is
used. A typical coating series consists of a diffusion barrier deposited by chemical
vapor deposition (CVD), which is followed by a bond coat deposited by physical vapor
deposition (PVD) and topped by a thermal barrier layer applied by atmospheric plasma
spraying (APS).
This work has been performed in the context of an integrative materials modeling of
the whole manufacturing process, which is the aim of the collaborative research center
SFB 370 of the Deutsche Forschungsgemeinschaft (SFB 370 “Integrative Werkstoff3
4
CHAPTER 1. INTRODUCTION
modellierung”). This work is focused on the solidification of the base material and the
behavior of the bond coat under operating conditions. First, aspects of the casting
process of a typical turbine blade base material, the Ni-base superalloy CMSX-4, are
simulated. Phase formation and microsegregation are investigated during solidification
and cooling. Secondly, phase decomposition in the bond coat due to diffusion processes
between bond coat and thermally grown oxide layer (TGO) and interdiffusion processes
between bond coat and base material are investigated under operating conditions of
the blade.
1.1
Motivation
This work is part of a collaborative research center which aims at a continuous simulation of the microstructure for a product throughout the entire manufacturing process from its raw materials to its final shape. The motivation is to understand the
microstructure evolution during processing in order to control the properties of the
product. To show the correlation between processing steps and final microstructure
every step needs to be modeled and integrated into the chain of a continuous simulation. This shall be demonstrated for the example of a turbine blade. This work wants
to contribute towards understanding by means of thermochemical simulations. A tool
for describing diffusion controlled phase transformations in multicomponent systems is
applied to understand the selected aspects of microstructure evolution.
Chapter 2
State of the Art
2.1
Turbine Blades
A state-of-the-art turbine blade for gas turbines is run with hot gas at a temperature
of up to 1470 K under high mechanical load. This imposes extreme demands on the
material system. As base material single crystalline Ni or Co based superalloys are in
use, because they possess the necessary high temperature hardness/creep resistance.
In addition to an effective cooling system the blade is coated with a thermal barrier
layer on top which typically consists of partially yttria stabilized zirconia (PYSZ). In
between thermal barrier and base material a bond coat is needed. The bond coat
serves as an oxidation barrier and enhances the adhesion of the coatings. It consists
of MCrAlY, where M stands for Ni or Co. Under working conditions the bond coat
will form an impervious alumina layer, which stops the diffusion of oxygen into the
underlying material. This alumina layer is also known as thermally grown oxide layer
(TGO). Finally a thin alumina layer is deposited between base material and bond
coat to prevent interdiffusion processes between these materials. An overview over the
system of base material and coatings is shown in Fig. 2.1.
The efficiency of a turbine can be improved by increasing the inlet temperature of the
gas. Thus research is directed towards more efficient cooling systems as well as towards
the development of new superalloys with improved high temperature properties. For
the latter the processing of the alloys has been developed from forging over casting
to directional solidification and the production of single crystalline material. Simultaneously the composition of the alloys has been tuned by the addition of refractory
elements like Re, W, Ta, and Mo. The here investigated superalloy CMSX-4 is a single
crystalline superalloy with low contents of Re, W, and Ta, which is already in use
industrially. A new generation with increased content of refractory elements is being
developed at the moment.
Often the lifetime of a blade is limited by the coating integrity. Diffusion processes
under working conditions change the microstructure of the coatings which may be
partly responsible for coating failure. In this work diffusion processes in the bond coat
have been investigated which cause depletion of the dispersed β phase in the bond coat
material. The results will be incorporated in a model for lifetime prediction which is
5
6
CHAPTER 2. STATE OF THE ART
Figure 2.1: Material system of a turbine blade consisting of the base material and
several protective layers.
being developed within the framework of the collaborative research center SFB 370.
2.2
Microstructure Simulation of Turbine Blades
A number of approaches has been developed to simulate the formation of microstructure during solidification or heat treatment within the past years. A general overview
over simulation methods for systems with up to 3 components is given by Kraft and
Exner [9, 10]. For multiphase and multicomponent alloys a very good review has been
published recently by Hecht et al. [11]. Since many factors influence the microstructure, simulation methods cannot yet include all, but need to focus on selected aspects.
Generally speaking, two main directions are pursued. One is the correct simulation
of phase fractions and constitution, while the other is more concerned about the morphological stability and dynamic behavior of phase patterns during coupled growth.
For the latter phase field simulations have established themselves as the method of
choice in the last few years. On the other hand side computational thermodynamics
is the most advanced method for the simulation of phase fractions and constitution if
morphological aspects are of minor importance.
The calculations in this work belong to the field of computational thermodynamics.
Computational thermodynamics are based on the concept of minimizing Gibbs energy. Today multicomponent databases exist for the most important material systems
providing a parametric description of the Gibbs energies of the phases as functions
of composition and temperature. These databases have been assessed from experimental data using the CALPHAD method, which has been discussed extensively in
the literature [4]. The development of databases and a suitable software to calculate
phase equilibria started in the 1950’s and computational thermodynamics are nowadays applied in many fields of materials science. Generally they can be divided into
equilibrium calculations and diffusion simulations, which also consider kinetic effects.
2.2. MICROSTRUCTURE SIMULATION OF TURBINE BLADES
7
Most frequently used are simple equilibrium and Scheil calculations. Both rely on
thermodynamic data only. In the early 1990’s the software DICTRA was developed,
which combines thermodynamic calculations with the kinetics of diffusion controlled
phase transformations. It is limited to 1D systems, but offers an accurate treatment of
multicomponent diffusion and phase equilibria.
The application of computational thermodynamics in turbine materials started in 1996,
when Saunders developed a data base for Ni-based superalloys [12]. Since then a number of successful thermodynamic calculations for Ni-based superalloys has been published [13] - [16]. Most of these calculations are equilibrium calculations. For the alloy
706 the solidification has been successfully modeled with a Scheil calculation [15]. The
latest development in this field is aimed at coupling the thermodynamic calculations
with physical models in order to model the physical properties of a material [17, 18].
Non-equilibrium processes are treated with the Scheil model. Matan and coworkers
developed a coupled thermodynamic/kinetic model for diffusional processes in superalloys, which has so far been tested for ternary diffusion couples only [19]. A multicomponent diffusion mobility database for Ni-based superalloys has been developed by
Campbell and tested on diffusion couples of Ni/Rene-88 and IN718/Rene-88 [5, 20].
This work includes equilibrium and Scheil calculations for CMSX-4, but focuses on
diffusion simulations, which include the kinetics of diffusion controlled phase transformations. It could be shown that results improved significantly for CMSX-4 when
diffusion simulations were used instead of the simplified Scheil model. The components
Al, Co, Cr, Ni, Re, Ta, Ti, and W were included in this simulation. Furthermore, this
work comprises diffusion simulations of the β-phase depletion in the bond coat. The
bond coat material, a so-called MCrAlY coating, is simulated as a two phase material
consisting of the components Co, Cr, Al, and Ni. The small amount of Y has to be neglected due to a lack of data. The bond coat interacts with the base material CMSX-4
on one side and with a thermally grown oxide layer (TGO) on the other side. Only
the latter has been investigated in the simulations.
Lately the simulation of microstructure with phase field simulations has become increasingly popular - especially in areas where a 1D simulation is not sufficient. As a
future trend the coupling of phase field methods and computational thermodynamics seems promising for a locally resolved simulation of multicomponent alloys [21].
Warnken and coworkers calculated first results for the solidification of a five component
Ni-based superalloy and good qualitative agreement was achieved for the segregation
profiles [22].
8
CHAPTER 2. STATE OF THE ART
Chapter 3
Simulation Methods
Three different types of calculations have been performed: equilibrium calculations,
Scheil calculations, and diffusion simulations. Equilibrium and Scheil calculations were
performed using the software Thermo-Calc [3]. For the diffusion simulations the software package DICTRA was used [3]. With these calculations phase formation, fraction
solid, and microsegregations during solidification of the base material have been treated
as well as phase transformations at high temperatures in the bond coat of the turbine
blade. Equilibrium and Scheil calculations are rather rough estimates for most real
problems and can only give an upper and a lower limit to the problem solution. The
diffusion simulations describe the problems more accurately, but on the other hand
they need more input information and are more time consuming.
3.1
Equilibrium Calculations
Phase equilibria are calculated by minimizing the Gibbs energy of a system. The molar
Gibbs energy of a solution phase can be described as
Gφm = Gφ,ref + Gφ,id + Gφ,ex .
(3.1)
P
It consists of a reference term from the pure elements, i, Gφ,ref = i xφi Gφ,0
i , a term
P φ
φ
φ,id
describing ideal mixing G
= RT i xi ln(xi ), and an term for the non-ideal, or soφ,ex
called excess, mixing G . Different models exist to describe the Gibbs excess energy
for multi-component systems. In this work the approach by Muggianu was used. For
a system of three components a, b, and c it expands to:
Gφ,ex = xa xb [L0ab +L1ab (xa −xb )]+xa xc [L0ac +L1ac (xa −xc )]+xb xc [L0bc +L1bc (xb −xc )]+xa xb xc L0abc .
(3.2)
The interaction parameters L0ij and L1ij and the Gibbs energies of the pure components
Gφ,0 are available as polynomial functions of temperature in data bases. The data have
been assessed according to experimentally available data, e.g. measured enthalpies,
activities, and transformation temperatures. This method of data assessment is known
as CALPHAD method and has been discussed extensively in the literature [4]. The
9
10
CHAPTER 3. SIMULATION METHODS
data used here for the turbine blade material system stem from the National Institute
of Standards and Technology (NIST) [6].
Based on this description the software Thermo-Calc calculates the equilibrium state
of a system by minimizing its Gibbs energy. The calculation yields phase fractions,
compositions, and transformation temperatures as well as other thermodynamic quantities.
3.2
Scheil Calculations
The Scheil calculation describes the case of maximum segregation for a solidification
process. It is named after one of his inventors: E. Scheil and G.H. Gulliver [7] [8]. Scheil
calculations assume a homogeneous liquid phase, no diffusion in the solid phase and
local equilibrium at the phase boundary. Under these conditions maximum segregation
takes place in the solid phase. In most cases real solidification will be located in between
the equilibrium and the Scheil calculations. The advantages of Scheil calculations are
the small computational effort and the fact that no kinetic data are needed. Therefore
Scheil calculations are frequently used to estimate the solidification sequence and the
fraction solid as a function of temperature.
Fig. 3.1 demonstrates schematically the solidification path according to a Scheil calculation in a binary phase diagram. An alloy of composition c0 is solidified in temperature
steps of ∆T , which are small compared to the solidification interval. After every step
the composition of the solid cα and the melt cm are determined by an equilibrium calculation. Since diffusion in the solid phase is neglected, the solid remains unchanged and
the solid material grows in concentration layers as shown on the left side of Fig. 3.1. It
does not participate further in the solidification process. The calculated composition of
the melt serves as a new starting point and the procedure is repeated until the eutectic
temperature of the system or a threshold amount of liquid has been reached. A typical
ca2
ca3
c0
Temperature
ca1
ca1
ca2
ca3
melt
cm1
cm2
cm3
a
a+b
A
Mole fraction B
B
Figure 3.1: Schematic drawing of the principle of a Scheil calculation for a binary
system.
3.3. DIFFUSION CALCULATIONS
11
threshold value for the residual amount of liquid is 1 vol pct.
3.3
Diffusion Calculations
In real alloys the rate of phase transformations is often controlled by diffusion processes. Thus time is an important variable and kinetic data need to be included in
an accurate model for solidification and high temperature behavior. A suitable software package to treat such diffusion controlled phase transformations is DICTRA. It is
closely linked with the Thermo-Calc software, which provides all necessary thermodynamic calculations. Additionally it incorporates kinetic data in the form of mobilities
for the elements in multicomponent systems. The data used here for the base material
and the bond coat of the turbine blade stem from NIST [5]. Thus phase fractions, segregation profiles, etc. can be simulated resolved in time and one dimension of space.
This restriction to one dimension demands a careful modeling of the problem setup if
more than one solid phase is involved. For solidification a maximum number of two
solid phases can be simulated in contact with a liquid phase. In this case the liquid
is located in the middle and shares one moving phase boundary with a solid to the
left and another with the second solid to the right. Redistribution of the elements in
the liquid is very fast and processes like eutectic growth can be modeled. Generally
phases can be treated as regions separated by moving phase boundaries or as one region with several phases present as dispersed particles in a matrix phase. The two
models cannot be combined. If a simulation setup contains a moving phase boundary
it cannot contain any dispersed phases. This is a further restriction by the software.
However, many problems can be represented meaningfully by one of the models or by
a sequence of such. A general description of the used simulation models will be given
in the following two chapters.
3.3.1
Moving Boundary Systems
A simple setup for a moving phase boundary calculation is shown in Fig. 3.2. It shows
two phases, α and β, separated by a phase boundary moving with a velocity v α . The
concentration profiles cαk and cβk of a component k in the two phases are indicated. To
Figure 3.2: Simulation setup for a moving phase boundary simulation.
12
CHAPTER 3. SIMULATION METHODS
Figure 3.3: Schematic of a moving phase boundary simulation.
find the velocity v α , with which the phase boundary moves during a time step ∆t, the
procedure shown in Fig. 3.3 is followed. The simulation relies on thermodynamic and
kinetic data stored in data bases. The thermodynamic calculations of multicomponent
Gibbs energies and local equilibrium function as described in section 3.1. The same
data are used.
The kinetic data needed to calculate the multicomponent diffusion coefficients is stored
in the form of mobilities of the elements i:
Mi =
Mi0
−Qi
exp
.
RT
RT
(3.3)
The diffusional flux of a component k in a direction z for an isothermal, isobaric onephase binary alloy is given by Fick’s first law
Jk = −Dk
∂ck
.
∂z
(3.4)
It denotes the concentration gradient as driving force for diffusion. Today it is generally understood that the gradient in the chemical potential is the true driving force.
Nevertheless concentration gradients are still frequently used because they are a valid
approximation for many materials and more convenient to handle. A concentration
gradient can be transformed into a chemical potential gradient multiplying by the
i
, where γi denotes the activity coefficient of the element
thermodynamic factor 1 + dlnγ
dlnci
i.
However, for multicomponent systems it was found that the diffusion of one component does not only depend on its own chemical potential gradient, but can also be
driven by the chemical potential gradient of any other component in the system. Thus
a multicomponent extension of Fick’s first law was formulated by Onsager [1], who
established a linear relationship between a flux and every thermodynamic driving force
3.3. DIFFUSION CALCULATIONS
13
in the system
Jk = −
n
X
L0ki
i=1
∂µi
.
∂z
(3.5)
Here µi is the chemical potential of an element i and depends only on the composition. L0kj is a phenomenological coefficient, which depends on the mobilities Mi of the
individual species, which is defined as:
L0kj
=
n
X
(δik − ck Vi )ci yva Mi
(3.6)
i=1
where δik is the Kronecker delta: i.e. δik = 1 when i = k and δik = 0 otherwise. ck and
ci are the amounts of k and i per unit volume, Vi is the partial molar volume of element
i, and yva is the fraction of vacant lattice sites on the sublattice, where i is dissolved.
Many times it is found more convenient to work with concentration gradients instead
of gradients in chemical potential. In this case eq. 3.5 can be rearranged to a form
resembling Fick’s first law:
n
X
∂cj
Jk = −
Dkj
,
(3.7)
∂z
j=1
with
Dkj = −
n
X
i=1
L0kj
∂µi
.
∂cj
(3.8)
Together with the continuity equation
∂ck
∂
=
(−Jk )
∂t
∂z
(3.9)
this gives a system of coupled partial differential equations, which can be solved numerically to determine the flux of each component [2].
The concentrations cαk and cβk at the phase boundary are calculated assuming local
equilibrium and now the flux balance at the phase boundary
v α (cαk − cβk ) = Jkα − Jkβ
(3.10)
can be solved for v α . This way the concentration profiles in each phase and the migration of the phase boundary are calculated for each time step ∆t.
3.3.2
Dispersed Systems
A simple simulation setup for a system with dispersed phases is shown in Fig. 3.4. A
phase β exist as dispersed particles in a continuous matrix phase α which shows a gradient in composition as indicated for one element i. Diffusion is assumed to take place
only in the matrix phase. The dispersed phase acts as a point sink or source for solute
atoms. The amount and composition of the dispersed phase is calculated assuming
local equilibrium at a number of node points along the space axis. A correction factor,
14
CHAPTER 3. SIMULATION METHODS
Figure 3.4: Simulation setup for a dispersed system.
the labyrinth factor, can be chosen proportional to the volume fraction of the dispersed
particles and accounts for their influence on the diffusion in the matrix.
Fig. 3.5 shows the calculation scheme. For a time step ∆t the diffusion equations for the
matrix phase are solved as described in the previous section. From this the new average
composition at each node point can be calculated. Then the new composition of the
matrix phase is calculated at each node point assuming local equilibrium. Now the
diffusion equations are solved again for the new composition and a new time step ∆t,
etc. This way the concentration profiles in the matrix phase and the phase fractions at
each node point can be simulated for different temperature-time exposure of a material.
Figure 3.5: Schematic of a dispersed system simulation.
Chapter 4
Results and Discussion
Since the modeling itself was a challenge and not straight forward, a detailed description
of the specific parameters for each case is heading the actual simulation results and
discussion.
4.1
Solidification and Cooling of CMSX-4
Solidification and cooling were simulated in a temperature interval from liquidus temperature down to 800 K. This temperature range covers the solidification and also the
phase transformations in the solid state of CMSX-4. The diffusion simulation was
performed in three steps; the primary solidification from the liquidus until γ’ starts to
form, the secondary (final) solidification of the γ/γ’ eutectic, and the precipitation of
γ’ during further cooling.
Solidification was modeled as a moving boundary problem. The simulation volume
is an infinitely high cylinder with the radius of half a primary dendrite arm spacing.
The simulation setup is shown in Fig. 4.1. During primary solidification the dendrite
Figure 4.1: Micrograph of a transverse section in the mushy zone of a quenched sample
at a fraction solid of about 0.91 [24]. Measuring paths of the EDX line scans and the
corresponding simulation set-up are shown. The <100> direction corresponds to the
horizontal/vertical direction of the micrograph.
15
16
CHAPTER 4. RESULTS AND DISCUSSION
matrix phase γ grows from the dendrite core (bottom of the cell) into the liquid phase.
Secondary solidification starts when the ordered phase γ’ nucleates as part of the γ/γ’
eutectic (at the top of the cell) and the two solid phases grow towards each other until all liquid phase disappears. The solidification simulation had to be split up into
primary and secondary solidification because of the phase description of γ and γ’ in
the database. For γ a simple fcc-A1 description was used consisting of one sublattice
for the substitutional elements. To describe the order-disorder transformation a fccL12 model is needed with two sublattices, where the substitutional elements can be
arranged either randomly or in an ordered position. The disordered fcc-L12 is thermodynamically identical to the fcc-A1 description. fcc-A1 was chosen for the γ-phase,
because the available diffusion data are formulated for one sublattice for the substitutional elements. For the ordered γ’ the fcc-L12 model had to be taken, which includes
the possibility for this phase to appear as γ if the disordered modification is more stable. Therefore primary solidification with the γ-phase growing from the bottom of the
cell in Fig. 4.1 into the liquid phase was simulated and equilibrium calculations were
performed for the liquid concentration at the other end of the cell to detect the moment
of nucleation for the ordered γ’-phase. Then a new simulation was started with both
solid phases present. Concentration profiles for γ and liquid were taken across from
primary solidification and a small nucleus of γ’ was added at the top of the cell with a
concentration according to the equilibrium calculation. Otherwise it was not possible
to prevent disordered γ to grow from both sides into the liquid.
During further cooling after solidification γ’ starts to form finely dispersed precipitates
inside the γ dendrite. For this solid state transformation a model for dispersed systems
was used. It is a limitation of the simulation software that dispersed phases and moving
phase boundaries cannot be handled at the same time. Therefore only the dendritic γphase with its segregation profiles was adopted from the solidification simulation. The
influence of the interdendritic γ’ during further cooling had to be neglected. This end of
the simulation cell was simply cut off. Instead γ’ was introduced as a dispersed phase in
the γ matrix. To avoid having the phase to precipitating in the disordered state (due to
the phase-description, as mentioned above) starting values for the ordered composition
were calculated with an equilibrium calculation. The influence of the precipitates on
the diffusion in the matrix phase was modeled with the so-called labyrinth factor. This
factor can be defined as a function of the volume fraction of the dispersed phase and
is multiplied by the rate of diffusion. As all phases present have a similar diffusivity,
this factor was set to one [23].
The eight most abundant elements of CMSX-4 were included in the simulation with a
composition according to Table 4.1. The thermodynamic and kinetic data used in the
Table 4.1: Composition of CMSX-4 in wt pct as used in the simulation
Al Co
5.6 9.6
Cr
6.5
Re Ta
2.9 6.5
Ti
1.0
W
6.4
Ni
bal.
4.1. SOLIDIFICATION AND COOLING OF CMSX-4
17
simulations stem from NIST (National Institute of Standards and Technology) [5, 6].
It was primarily evaluated to treat the solidification of Ni-base superalloys into γ and
γ/γ’ structures. Other phases, e.g. bcc, possibly appearing at lower temperatures, were
given less attention. Diffusion in the ordered phase was neglected. The simulation
parameters were chosen to be comparable to solidification experiments by Ma and
Grafe [24]. They investigated the directional solidification of CMSX-4 samples at a
constant thermal gradient of 11.1 K/mm and a withdrawal speed of 0.125 mm/min,
which was the critical velocity for the morphological transition from cells to dendrites.
This results in an effective cooling rate of 0.023 K/s. The primary dendrite/cell spacing
was 320 µm.
4.1.1
Fraction Solid
Fig. 4.2 shows a comparison between the fraction solid calculated by equilibrium, Scheil
and diffusion simulations. The onset of γ’ formation during Scheil and diffusion simulations is indicated. The diffusion simulation results in a fraction solid between Scheil and
equilibrium calculations. Differences between Scheil and diffusion simulation become
more pronounced with progressing solidification.
An equilibrium calculation describes an infinitely slow solidification and always overestimates the fraction solid. Scheil calculations on the other hand usually underestimate
it by neglecting back diffusion in the solid phases. Nevertheless solidification is described quite well with the Scheil model in some systems, at least if no interstitial
elements are present. The applicability of Scheil calculations to superalloys has been
Figure 4.2: Comparison of fraction solid from different calculation methods.
18
CHAPTER 4. RESULTS AND DISCUSSION
studied in several cases and was generally found to be valid [4, 25, 26]. In this investigation this was not found so for the calculation of the fraction solid of CMSX-4.
As expected the diffusion calculation predicts a fraction solid that lies between the
equilibrium and Scheil predictions. However the results from Scheil and DICTRA calculations deviate significantly. When the alloy is completely solidified according to the
diffusion simulation, the Scheil calculation still indicates 17 vol pct liquid. The solidification interval from the equilibrium calculation is 44 K and about 80 K from the
diffusion simulation. In comparison the Scheil calculation indicates an unreasonably
large solidification interval of more than 200 K. In this case complete solidification is
assumed when less than 1 vol pct liquid is present in the calculation. During the Scheil
calculation the bcc phase had to be suppressed, since it otherwise would have appeared
in the solidification sequence.
The temperatures of beginning and complete solidification calculated by the equilibrium calculation correspond to the solidus and liquidus temperatures of the material.
Literature values for solidus and liquidus temperatures of CMSX-4 are sparse and
scatter, especially for the solidus temperature. Differential thermal analysis (DTA)
by Wilson et al. revealed a solidus temperature of 1617 K [27]. Palumbo et al. determined a solidus temperature of 1598±3 K from differential scanning calorimetry
(DSC) at different rates [28]. The equilibrium calculation shows a reasonably small
deviation from these experimental data. The calculated solidus temperature of 1599 K
lies well within the scatter of the literature data. The calculated liquidus temperature
of 1643 K lies below but within a reasonable distance from the experimental values
of 1657 K [28] and 1659 K [27]. There are general doubts about the measurability of
a solidus temperature of multicomponent materials by means of DTA and DSC [29].
Due to lack of other data, it will still be regarded a rough estimate here.
The onset of secondary solidification is a third representative temperature for the solidification process of CMSX-4. Palumbo and coworkers did differential scanning calorimetry (DSC) experiments for CMSX-4 at cooling rates between 1 and 10 K/min [28]. They
detected a peak that may indicate the onset of γ’ formation 64 K below the liquidus
temperature with both temperatures being independent of the cooling rate. According
to the diffusion simulation the onset of γ’ formation lies 61 K below the liquidus temperature, and according to the Scheil model it lies 66 K below the liquidus temperature.
Both calculations are in excellent agreement with the measurements.
4.1.2
Microsegregation
CMSX-4 undergoes an extensive heat treatment after casting to dissolve the eutectic
and refine the γ’ precipitation inside the dendrites. The solutioning temperature has to
be approached stepwise in order to avoid local melting due to microsegregations in the
as-cast structure. Microsegregation in CMSX-4 has been investigated experimentally
by Ma and Grafe [24]. They determined the concentration profiles across a dendrite
by energy dispersive X-ray analysis (EDX) towards the end of solidification as pointed
out in Fig. 4.1. Figs. 4.3 - 4.5 show the results of their measurements for the most
abundant elements of CMSX-4 in comparison to results from Scheil calculation and
4.1. SOLIDIFICATION AND COOLING OF CMSX-4
19
diffusion simulation.
The dendrite center is situated at the left hand side and the interdendritic γ’ would
form at the right end of the cell, but is not included, since the measurements were
done on samples quenched before the end of primary solidification. All plotted profiles
represent the state at the end of primary solidification. At this point the amount of
residual liquid is larger for the Scheil calculation than for the diffusion simulation and
therefore Scheil segregation profiles are shorter.
Despite significant differences in the calculation of the fraction solid, both Scheil and
diffusion calculations yield similar microsegregation profiles. Compared to the measured data the profiles for Co, Cr, Ti, and W are within the precision of the measurement. For Re the Scheil calculation deviates somewhat more from the measurement,
but remains in excellent agreement with the diffusion calculation. For Ta similar behavior is observed with a slightly stronger segregation tendency for the Scheil model
compared to the diffusion simulation. For these heavy and slow diffusing elements the
good agreement between Scheil and diffusion calculations does not surprise. For the
light element Al the differences in composition at the center of the dendrite between
the two calculation methods can be explained by the effect of back diffusion in the
diffusion simulation.
In the diffusion simulation the general trend for accumulation or depletion in the liquid
of each single element is predicted correctly. Figs. 4.3 - 4.5 show that the calculated
profiles for Co, Cr, Ti, and W are within the precision of the measurement. The profiles for Re, Ta, and Al deviate somewhat more. Towards the interdendritic region with
Figure 4.3: Solute distribution plotted vs. distance from the dendrite center, along
<100> direction. Results from measurement (symbols), Scheil calculation (dashed
line) and diffusion simulation (solid line).
20
CHAPTER 4. RESULTS AND DISCUSSION
Figure 4.4: Solute distribution plotted vs. distance from the dendrite center, along
<100> direction. Results from measurement (symbols), Scheil calculation (dashed
line) and diffusion simulation (solid line).
Figure 4.5: Solute distribution plotted vs. distance from the dendrite center, along
<100> direction. Results from measurement (symbols), Scheil calculation (dashed
line) and diffusion simulation (solid line).
4.1. SOLIDIFICATION AND COOLING OF CMSX-4
21
the eutectic γ/γ’ microstructure the measurements show steeper slopes over a shorter
distance than the calculated concentration profiles exhibit. This may be due to an
influence of the interdendritic γ’ phase on the EDX measurement, so that measured
compositions close to the interdendritic region represent a mixture of γ and γ’. The
scatter in the measurement is also larger in this region.
Ma and Grafe developed an analytical 2-dimensional model taking account of the
anisotropy of crystal growth and compared results to their measurements cited
here [30] [31]. They treat the morphology of the dendrite in much more detail, but
base their multicomponent thermodynamics on linearized phase diagrams. Using the
Scheil model to determine the fraction solid their calculated segregation profiles were
qualitatively right, but deviated quantitatively to the measurements [31]. Using the
model of Bower, Brody and Flemings to determine the fraction solid, their results are
of comparable quality to the diffusion simulations here [30]. Keeping in mind that the
diffusion simulations are based on a simple one-dimensional model, not accounting for
the morphology this is a very good result. Nevertheless results could still improve with
improved data, especially for less common elements like Ta.
4.1.3
Solid State Transformation
After complete solidification CMSX-4 experiences an additional phase transformation
in the solid state. γ’ precipitates finely dispersed inside the dendrites and the resulting
γ/γ’ structure is responsible for the elevated temperature strength after refinement
through a heat treatment.
Fig. 4.6 shows the amount of γ’ at different temperatures throughout the dendrite.
The microsegregation profiles showed increasing concentrations towards the eutectic
for the elements partitioning preferentially into γ’. In agreement with this the ordered
phase starts to precipitate close to the eutectic before it forms throughout the whole
dendrite with increasing undercooling. The initial difference in mole fraction of 0.3
between dendrite tip and center equalizes to 0.15 with progressing cooling. Below a
temperature of 800 K there are no more significant changes. Finally a fraction of 62
vol pct of γ’ formed in the core of the dendrite and more than 70 vol pct towards the
interdendritic region.
It can be observed in Fig. 4.6 that the amount of γ’ is not increasing monotonically, but
reaches its maximum shortly before the eutectic region. The location of the maximum
approximately coincides with the position of the solidification front at the end of the
primary solidification. This phenomenon can also be observed experimentally. Fig. 4.7
shows a transverse section of a directionally solidified sample in which the individual
dendrites can be clearly distinguished. The shade of the dendrites changes from white
in the center to a dark grey close to the interdendritic regions. Darker color indicates
a higher amount of γ’ precipitates. Fig. 4.8 shows this microstructure with higher
magnification. It can be observed that the γ/γ’ eutectic of the interdendritic regions
in many places is surrounded by a narrow white band before the dark grey marks
the presence of a high amount of finely disperse γ’. Thus, the general trend of the
calculated amount of γ’ is in agreement with the observed microstructure.
22
CHAPTER 4. RESULTS AND DISCUSSION
Figure 4.6: Calculated mole fraction of γ’ vs. distance from the dendrite center at various temperatures. The vertical dashed line indicates the position of the solidification
front at the end of the primary solidification.
Figure 4.7: Light optical micrograph (LOM) of a horizontal section of a directionally
solidified and quenched sample of CMSX-4 after etching, just above the bottom of
the dendrites. The primary γ dendrites and the rosettelike interdendritic eutectic can
easily be distinguished.
4.1. SOLIDIFICATION AND COOLING OF CMSX-4
23
Figure 4.8: LOM of a representative interdendritic region from the same sample as in
Fig. 4.7 at a higher magnification. Coarse γ’ precipitates have formed from the solid,
but close to the eutectic, no precipitates can be found.
24
4.2
CHAPTER 4. RESULTS AND DISCUSSION
β Phase Depletion in the Bond Coat
The bond coat is a layer of MCrAlY between base material and thermal barrier coating
that enhances the adhesion of the coatings to the base material and serves as an
oxidation barrier. MCrAlY coatings typically exhibit a two-phase microstructure of a
dispersed β phase in a γ matrix. The β phase possesses an ordered bcc-B2 structure
with NiAl stoichiometry. The MCrAlY coating hinders oxygen to enter the underlying
material by forming an impervious alumina layer under working conditions (TGO).
For this process aluminum diffuses towards the outside of the bond coat to form the
oxide and causes a depletion of aluminum in the MCrAlY. Additionally aluminum
diffuses from the bond coat into the base material if no diffusion barrier is present.
As the amount of Al decreases, the β phase tends to dissolve. For this reason, it is
often described as an aluminum reservoir, and coating life often measured in terms
of depletion of β. Fig. 4.9 shows a MCrAlY bond coat after several hours of high
temperature exposure. The β depletion zone can be clearly distinguished from the
γ+β microstructure of the intact bond coat material.
Figure 4.9: Microstructure of the bond coat after several hours of high temperature
exposure.
To model the β depletion in the bond coat the setup for dispersed systems was used.
The cell size was 200 µm, which corresponds to the thickness of the bond coat. As
mentioned before the used software cannot handle dispersed phases and moving cell
boundaries in the same simulation setup. Thus the cell size and volume were fixed for
this simulation. The kinetics of the oxide formation have been measured experimentally for the commercially available MCrAlY alloy PWA 286 at different temperatures.
Fig. 4.10 shows the thickness of the oxide layer over time at different temperatures.
Assuming that the oxide formed was pure alumina with a density of 3980 kg/m3 the
resulting flow of aluminum from the bond coat was calculated. This flow was then
defined to leave the simulated bond coat volume across one cell boundary.
Since aluminum is not an interstitial component it has a volume contribution. The
4.2. β PHASE DEPLETION IN THE BOND COAT
25
Figure 4.10: Measured oxidation kinetics for the MCrAlY bond coat at different temperatures
Al flow leaving the simulated volume therefore needs to be replaced, since the overall
volume is fixed. By default the majority component, here Ni, is used to replenish the
simulated volume. This will then change the composition and consequently influence
the thermodynamics and kinetics and alter the simulation results. The replacement
of the oxidized aluminum is unavoidable due to the simulation method. To minimize
the introduced error, an alternative to the replacement by Ni has been developed. The
oxidized Al has been compensated by an inward flow of Ni, Co, and Cr and the ratio of
the amount of these components has been chosen equal to their ratio in the bond coat
material. This way the MCrAlY composition is not altered by an excess of Ni. For
the implementation new components have been defined with the exact thermodynamic
properties of Ni, Co, and Cr. For numerical simplicity all three components were given
the mobility of Al, since they enter the cell at the boundary and have to diffuse in
opposite direction to the Al. Fig. 4.11 shows a comparison of the two replacement
alternatives. Plotted is the mole fraction of β throughout the bond coat initially and
after 5000 h at 1373 K. Starting and end point for the β distribution after 5000 h are
nearly identical for the two alternatives. The simulation with Ni replacement yields a
steeper transition from the zone with a constant mole fraction of β to the depletion
zone. It differs by maximum 0.1 mole fraction from the simulation with a Ni, Co, Cr
replacement. The later is numerically unstable due to the complex boundary conditions
and opposed diffusion flows and needs calculation times in the order of days, while the
simulation with Ni replacement is stable and has a running time in the order of minutes.
Therefore the developed alternative has been discarded accepting the introduced error
by a Ni replacement of the oxidized Al.
The final simulations were conducted for a bond coat of MCrAlY with a composition
as shown in Table 4.2. This composition corresponds approximately to PWA 286. Due
26
CHAPTER 4. RESULTS AND DISCUSSION
Figure 4.11: Simulated β depletion after 0 h and 5000 h at 1373 K, when the oxidized
Al is replaced by Ni only or by a mixture of Ni, Co, and Cr.
to a lack of thermodynamic and kinetic data small amounts of yttrium (0.6 wt pct),
silicon (0.3 wt pct), and hafnium (0.14 wt pct) were neglected in the simulation. The
labyrinth factor, that describes the influence of the amount of dispersed phase on the
diffusion, was set to 1 since diffusion may be assumed equally fast in β and γ.
Table 4.2: Composition of MCrAlY in wt pct as used in the simulation
Al
12
Co
21
Cr
17
Ni
bal.
The phase transformations in the bond coat were simulated for a time interval of
5000 h at 1223 K, 1273 K, and 1323 K. The calculated amount of β phase in the
bond coat at different times for these temperatures is shown in Figs. 4.12- 4.14. The
aluminum flow that is oxidized to form the TGO leaves the simulation cell across the
right hand side boundary. Two effects can be observed. Higher temperature favors the
oxidation reaction and the Al diffusion. The measured oxidation kinetics in Fig. 4.10
show that the TGO at 1323 K is approximately twice as thick compared to 1223 K.
A proportional decrease in the overall fraction of β can be observed in Figs. 4.124.14. Secondly faster diffusion results in flatter concentration profiles throughout the
material. The oxidized aluminum is not supplied by the closest region only, but the
diffusion processes involve more of the underlying material. At 1223 K in Fig. 4.12
the mole fraction of β throughout the MCrAlY remains almost unchanged at a value
of 0.6 and then drops to 0 within 30 µm from the TGO. Even after several hundred
hours the mole fraction of β at the left hand side of the cell remains unaffected and the
region where β starts being depleted has extended to 100 µm. At 1273 K the transition
4.2. β PHASE DEPLETION IN THE BOND COAT
Figure 4.12: Simulated mole fraction of β at 1223 K.
Figure 4.13: Simulated mole fraction of β at 1273 K.
Figure 4.14: Simulated mole fraction of β at 1323 K.
27
28
CHAPTER 4. RESULTS AND DISCUSSION
from constant mole fraction of β to the depletion zone is smoother and after 2784 h
the amount of β has dropped from 0.6 to 0.5 at the left hand side of the cell. At 1323
K the bond coat is involved in the diffusion process with full depth after less than 100
h and the amount of β starts to decrease at the far end before it is fully depleted in
the closest region to the TGO. Thus the simulation delivers qualitatively physically
meaningful results.
For further evaluation the simulation results were compared to experimental data.
Fig. 4.15 shows micrographs of a compound of base material, bond coat, TGO and
TBC (bottom to top) at different times during heat treatment at 1323 K. The bond
coat in the middle is PWA 286 with a thickness of approximately 200 µm. Its two
phase microstructure can easily be identified. The lighter shaded material is the γ
phase and the grey shaded is β. The thin dark grey band on top of the bond coat is
the TGO followed by a thick layer of TBC. After 96 h a thin β-depleted zone with
an average thickness of approximately 12 µm is visible close to the TGO. There is a
sharp transition from β-rich to β-depleted material. After 384 h the depletion zone
has grown to approximately 20 µm and the β-rich zone starts to show some coarsening
and a slight decrease in the amount of β. After 960 h the depletion zone extends
over approximately 40 µm and the decreasing β content and coarsening in the β-rich
Figure 4.15: Optical microscopy images of the bond coat PWA 286 after 96 h, 384 h,
960 h, and 2783 h of heat treatment at 1323 K.
4.2. β PHASE DEPLETION IN THE BOND COAT
29
material is now clearly visible. After 2784 h the depletion zone has reached an extension
of approximately 55 µm and only a smaller number of large isles of β is left in the rest
of the material. Since in this case there is no diffusion barrier between bond coat and
base material, a similar diffusion and depletion process takes place between these two
phases, but to a smaller extent. Fig. 4.15 shows this clearly and it will be taken into
account by the simulation in future work. If these experimental results are compared
to the simulation at 1323 K in Fig. 4.14, only an agreement for the general trend can
be found. In the simulation the depletion zone is less pronounced, the transition to the
β-rich zone is not as sharp and the overall decrease in β content is less than observed
experimentally.
Fig. 4.16 shows the simulation results if the flow of aluminum to be oxidized is multiplied by a factor of 1.5. The increased flow of aluminum leads not only to a stronger
overall decrease of β content, but also causes a sharper transition from β-rich to βdepleted zone at short times and a more pronounced depletion zone. After 2784 h the
simulation yields a depletion zone with a thickness of 45 µm which is in the range
of quantitative agreement with the approximate average thickness of 55 µm measured
experimentally. This agreement would be more than satisfying for the simplicity of the
here presented model compared to the complexity of the material. The factor 1.5 for
the oxidized aluminum could be explained by the meandering surface of TGO and bond
coat. The micrographs in Fig. 4.15 clearly show the corrugated nature of the interface
between TGO and bond coat. This has not been taken into account when evaluating
the oxidation kinetics from the thickness of the TGO, and thus an inaccuracy of the
factor 1.5 could easily be explained.
Figure 4.16: Simulated mole fraction of β at 1323 K if the measured oxidation kinetics
is multiplied by 1.5.
30
4.3
CHAPTER 4. RESULTS AND DISCUSSION
Interdiffusion between Bond Coat and Base
Material
The interdiffusion processes between bond coat and base material were already mentioned in the foregoing section about β depletion. The main reason for the formation of
β-depleted zones is the diffusion of aluminum out of the bond coat due to the formation
of a protective alumina layer. This process takes place at the interface between bond
coat and the thermal barrier coating (PYSZ). However, Fig. 4.15 showed that a similar
process was also taking place at the interface between bond coat and base material. To
a small extend the formation of a β-depleted zone could be observed at this interface,
too, which must be the result of interdiffusion processes. To prevent this interaction a
small diffusion barrier of alumina can be applied as shown in Fig. 2.1. Nevertheless one
question within the framework of the collaborative research center was the simulation
of a worst case scenario, where bond coat and base material are in direct contact due
to failure of the diffusion barrier.
Both bond coat and base material contain dispersed phases in a fcc-structured matrix.
For the simulation they were modeled as one region with a fcc-structured matrix and
a change in concentration at the position of the interface between the two materials.
This has the advantage of omitting a moving phase boundary in the simulation setup.
As already explained early moving phase boundaries and dispersed phases can not be
handled in the same simulation setup due to a software limitation. At the current
state of the simulations no dispersed phases are actually included because of numerical
problems, but this possibility exists for future work. Fig. 4.17 shows the initial concentrations for the simulation. The interface between bond coat and base material is
located at a distance of 20 µm. The bond coat lies to the left and the base material to
Figure 4.17: Initial concentration profiles as used in the simulation (lines) and as
measured with WDX (symbols)
4.3. INTERDIFFUSION BETWEEN BOND COAT AND BASE MATERIAL
31
the right of the interface. The composition of the MCrAlY and the Ni-based superalloy
were chosen comparable to experiments done in the collaborative research center and
are listed in Table 4.3 A Si content of approximately 2.5 at pct in the MCrAlY had to
be neglected in the simulations.
Table 4.3: Compositions in at pct as used in the simulation
Al
Co
bond coat
15.15 25.89
base material 13.06
-
Cr
Re
Ta
22.82 3.42
10.49
2.67
W
2.92
Ni
bal.
bal.
For this diffusion couple annealing experiments were done at 1273 K, 1373 K, and
1472 K for 10 h respectively. Afterwards the composition of the samples was analyzed
with wavelength dispersive X-ray microprobe analysis (WDX) along a linescan across
the interface. Experiments and analysis were done by J. Mueller within the framework
of the collaborative research center.
Figs. 4.18-4.20 show a comparison of simulated and measured concentration profiles.
The initial position of the interface is indicated by the dashed vertical line. The measured values show a much larger scatter after the heat treatment than before, especially
in the MCrAlY to the left of the interface (compare to Fig. 4.17). This may be attributed to coarsening of the dispersed phases, such that the measured composition
is not the average composition, but depends on whether the measuring signal is acquired mainly from matrix or dispersed material. For the elements Ni, Al, and Cr
the agreement between measurement and simulation is satisfactory. The equalization
Figure 4.18: Simulated (lines) and measured (symbols) concentration profiles after
annealing treatment for 10 h at 1273 K.
32
CHAPTER 4. RESULTS AND DISCUSSION
Figure 4.19: Simulated (lines) and measured (symbols) concentration profiles after
annealing treatment for 10 h at 1373 K.
Figure 4.20: Simulated (lines) and measured (symbols) concentration profiles after
annealing treatment for 10 h at 1473 K.
4.3. INTERDIFFUSION BETWEEN BOND COAT AND BASE MATERIAL
33
of the concentration profiles throughout the materials with increasing temperature is
reproduced well. The diffusion depth of Co from the bond coat into the base material is slightly overestimated by the simulation. Nevertheless the simulated overall
decrease in Co in the bond coat is in good agreement with the measurements. The
rate of diffusion of the refractory elements Re, Ta, and W is generally underestimated
by the simulation. This may be due to a lack of diffusion data for these less common
elements. Overall improvement of the simulation results may be achieved by including
the dispersed phases in the modeling. This has been unsuccessful so far because of
numerical difficulties.
34
CHAPTER 4. RESULTS AND DISCUSSION
Chapter 5
Summary and Conclusions
Diffusion processes during solidification and further cooling of the base material and
in the bond coat of a turbine blade under working conditions were simulated. Using a
fairly simple one-dimensional model for diffusion controlled phase transformations the
modeling of this complex material system was not straight forward. Strategies were
developed to break down and simplify the simulated processes and meaningful results
could be calculated with the available methods.
It was found that in the case of CMSX-4 the frequently used Scheil model severely
overestimates the solidification interval. Nevertheless the microsegregation profiles
calculated with the Scheil model were in good agreement with experimental results
and of similar quality compared to the diffusion simulation.
Microsegregation and fraction solid from the diffusion simulation were in good agreement with experimental results. Compared to other simulation methods the here presented method was of better or comparable accuracy for the prediction of microsegregation.
For the precipitation of γ’ in the solid state the experimentally observed formation of
a γ’-depleted zone next to the eutectic could be reproduced successfully.
The β-phase depletion in the bond coat at different temperatures was simulated showing only a general qualitative agreement with experimentally observed behavior. However, it was found that very good agreement between simulation and experiment could
be achieved multiplying the measured oxidation kinetics by a factor 1.5 for the simulation. An inaccuracy in the measurement of the oxidation kinetics due to the meandering
interface between bond coat and oxide layer was identified as possible explanation for
this factor.
First steps have been made to simulate the interdiffusion between bond coat and base
material. Results are promising, but further improvement of the modeling is needed.
The method used relies heavily on the quality of the thermodynamic and kinetic data
used. Thus further improvement of the simulation quality can be expected for improved
data.
Future work for the bond coat will include the β-phase depletion due to interaction
between bond coat and base material. In the frame work of the SFB 370 equivalent
simulations and more validating experiments will be performed for PWA 1386. The
35
36
CHAPTER 5. SUMMARY AND CONCLUSIONS
Composition of this bond coat differs only slightly from the here investigated PWA
286 and similar results are expected. Since experimental data for PWA 1386 are not
available yet, PWA 286 has been used to develop the method.
Part II
Towards Large Area Deposition of
MAX Phase Thin Films with PVD
37
Chapter 6
Introduction and Motivation
Metals and ceramics are two common classifications of inorganic materials. Typical
properties of a metal are strength, ductility, machinability, electric conductivity, and
low corrosion resistance. Ceramics can be described as stiff, brittle, stable at high
temperature, and with excellent corrosion resistance. Indeed most inorganic materials
can clearly be assigned to either the metals or the ceramics group and the properties
of these two groups seem to be opposed to each other. Transition metal carbides and
nitrides like TiC were among the few known ceramics, that also possessed some typical
metallic properties, e.g. electrical and thermal conductivity, until in the 1960s the first
MAX phases were discovered by Nowotny and Jeitschko [32, 33].
MAX phases are a group of materials, that combine metallic and ceramic properties to
a previously unknown extend. They are characterized by high elastic moduli, they are
machineable, they exhibit good damage tolerance, excellent thermal shock resistance,
good corrosion resistance, and they are good thermal and electrical conductors. This
unique set of properties makes the MAX phases extremely interesting and a high effort
is made to explore their potential applications in industry and domestic life.
This work explores the possibilities to synthesize thin films of MAX phases by physical vapor deposition. Long-term objective is the large area deposition of pure MAX
phase thin films, which will be the key for an industrial application of these materials.
Therefore three challenges need to be addressed:
• the deposition of single phase coatings,
• the homogeneous deposition over a large area,
• the deposition at low temperature.
The most efficient way to ensure homogeneous coatings over a large area is the
deposition from a single target, which has never been tried for the MAX phases. This
work is the first to focus on a process development directed towards these challenges.
The investigated MAX phases are Ti2 AlC, Ti2 AlN, and Cr2 AlC. This selection enables
an investigation of the effect of an increasing number of valence electrons from titanium
to chromium and from carbon to nitrogen on the properties of the MAX phases, which
39
40
CHAPTER 6. INTRODUCTION AND MOTIVATION
has already been studied theoretically by ab initio calculations [34, 35, 36].
6.1
MAX phases
MAX phases are ternary carbides or nitrides and their name is derived from their
constituents, which are early transition metals (M), A-group elements (A), and carbon
or nitrogen (X), see Fig. 6.1. Nowotny and Jeitschko discovered more than 100 ternary
carbides and nitrides in the 1960s, among them more than 30 phases that would later
be classified as MAX phases. Back then the experimental means for the synthesis of
sufficient amounts of phase pure MAX phase material were not available in order to
examine its properties. 30 years later in the 1990s Barsoum and El-Raghy succeeded in
producing phase pure Ti3 SiC2 bulk material by reactive hot pressing [37]. Since then
some of these materials have been produced in bulk form, but until today only a few
MAX phases, such as Ti3 SiC2 , are well characterized materials.
So far MAX phases with the stoichiometry M2 A1 X1 (211), M3 A1 X2 (312), and M4 A1 X3
(413) have been reported, see Fig. 6.2. They all possess a nanolaminated structure
where layers of metal-carbide are interleaved with A-group element layers as illustrated
Figure 6.1: Periodic table with elements forming MAX phases highlighted.
Figure 6.2: MAX phases reported in the literature [38].
6.1. MAX PHASES
41
by Fig. 6.3 showing the unit cell of a 211 MAX phase. Fig. 6.4 shows the charge-density
contours for TiC, TiAl, and Ti2 AlC, respectively, as calculated with ab initio methods.
It can be seen that the character of the binary metal-carbide and intermetallic, Fig. 6.4
a and b, is conserved in the ternary MAX phase, Fig. 6.4 c. Thus MAX phases exhibit
a stronger ionic/covalent bonding in the metal-carbide layer and a weaker metallic
bonding between the layers of metal-carbide and A-group element atoms. The relatively
weak bonding between the layers allows for gliding and is believed to be the reason for
the ductility and machinability of the material [39]. The stronger ionic bonds between
transition metal and carbon or nitrogen preserve the ceramic properties, e.g. phase
stability. Another consequence of the nanolaminated structure is a strong anisotropy
of all properties.
Figure 6.3:
Unit cell of a
211 MAX phase. Blue spheres
indicate the A-group element,
yellow spheres the transition
metal and black spheres the carbon/nitrogen atoms.
Figure 6.4: Charge-density contours for a)
TiC, b) TiAl, and c) Ti2 AlC.
The synthesis of MAX phase thin films was first reported by Nickl in 1972 [40]. He deposited Ti3 SiC2 by chemical vapor deposition (CVD). He observed that the phase had
no solubility range. Trying to deposit single phase Ti3 SiC2 , small changes in deposition parameters resulted in multiphase coatings. Therefore he suggested an alternative
route for the production of phase pure Ti3 SiC2 powder depositing a phase mixture of
Ti3 SiC2 and TiSi2 and afterwards dissolving the TiSi2 with hydrofluoric acid. He also
observed strong plastic deformation behavior for this phase which was not found for
the sintered material. In 2002 Palmquist reported the first successful deposition of
MAX phase thin films by physical vapor deposition (PVD) [41]. He used magnetron
42
CHAPTER 6. INTRODUCTION AND MOTIVATION
sputtering from two individual elemental targets of Ti and Si and co-evaporated C60
for the synthesis of epitaxial single-phase Ti3 SiC2 thin films. Since then a number of
MAX phase thin films has been produced successfully by PVD and only the relevant
publications for the material systems in this work will be mentioned in the following
section.
Cr2 AlC and Ti2 AlC were both discovered by Jeitschko in 1963 [32]. Thin films of
Cr2 AlC were successfully deposited by magnetron sputtering from three individual
targets of Cr, Al, and graphite at 1123 K [42, 43]. The equilibrium volume of the
thin film phase was compared to bulk material and ab initio calculations and good
agreement was found.
The only reported PVD experiment for the deposition of Ti2 AlC thin films was conducted by Wilhelmsson [44]. Epitaxial films containing Ti2 AlC and Ti3 AlC2 were
grown by magnetron sputtering from three elemental targets. A strong temperature
dependence for the formation of the MAX phases was observed with optimum conditions at 1173 K. Ti2 AlC could not be obtained as phase pure material.
A successful deposition of Ti2 AlN thin films has only been reported by Joelsson [45].
Films were prepared by reactive magnetron sputtering from an alloy 2Ti:Al target in
a mixed Ar/N2 discharge at a temperature of 1103 K. A hardness of 16.1 ± 1.4 GPa
and a Young modulus of 270 GPa were determined by nanoindentation, which is in
good agreement with the theoretical Young modulus of 293 GPa calculated by ab initio
methods [46].
The literature review shows that for all three phases research still focuses on the synthesis of single phase material from multiple targets. To reach the long-term goal of an
industrially applicable large area deposition of MAX phase thin films the deposition of
single phase material from a single target at low temperatures needs to be achieved,
which is the focus of this work.
Chapter 7
Thin Film Growth
7.1
Magnetron Sputtering
Magnetron sputtering is one of the most frequently used technologies of physical vapor
deposition (PVD). A general introduction to some of the scientific principles involved in
these processes was given by Chapman [47]. A magnetron consists of the target material
and a set of magnets with alternating polarity behind the target. The deposition takes
place in a vacuum chamber which is pumped down to the base pressure and refilled
with the discharge gas. Argon is commonly used as discharge gas due to its extremely
low reactivity. It can be ionized by e.g. collisions with electrons (electron impact
ionization) to form Ar+ and Ar2+ . If a negative voltage is applied to the target charged
particles are accelerated along the electric field. Upon impact of the discharge gas ion
on the target more electrons may be emitted from the target which again increase the
probability of ionization of a discharge gas atom. This finally leads to a steady state
where the number of ions and electrons produced becomes the same and the plasma
becomes self-sustaining. When the discharge gas ions collide with the target surface,
atoms with an average kinetic energy of 4-6 eV are knocked out of the surface, which
is called sputtering. When these atoms collide with the substrate they may form the
coating. The set of magnets behind the target creates a magnetic field as shown in
Fig. 7.1. The purpose of the magnetic field is to decrease the loss of electrons to the
walls but keep them close to the target to increase the amount of ionization. The
electrons emitted by the target are trapped by the magnetic field so they return to the
target directly or after performing a cycloidal motion around the magnetic field lines
as indicated in Fig. 7.1. This reduces the loss of electrons to the walls and increases the
path a single electron travels. This again increases the probability of ionizing collisions
with discharge gas atoms and results in a higher deposition rate of the process.
The effects of deposition parameters like substrate temperature, inert gas pressure, and
target-substrate positioning on thin film growth can be described by so called structure
zone models. Movchan and Demchishin developed a model that differentiates between
three zones with their own characteristic structure and physical properties depending
on the ratio of substrate temperature T to melting point of the sputtered material
Tm (see Fig. 7.1) [48]. In zone 1 (T/Tm ≤ 0.3) surface diffusion is limited and this
43
44
CHAPTER 7. THIN FILM GROWTH
Figure 7.1: Schematic of magnetron sputtering.
results in a porous film with high dislocation density consisting of tapered crystals with
domed tops which are separated by voided boundaries. Zone 2 (0.3 ≤ T/Tm ≤ 0.5)
allows for surface diffusion and columnar grains with increased grain size constitute
a dense film. In zone 3 (0.5 ≤ T/Tm ) bulk diffusion is enabled. The film consists
of equiaxed grains with a bright surface and an even larger grain size corresponding
to the structure of a fully annealed metal. Movchan and Demchishin developed their
model based on thin films deposited by electron beam evaporation. Thorton extended
the zone classification to sputtering processes. His diagram shown in Fig. 7.1 includes
an axis to account for the influence of the sputtering gas pressure. Lower sputtering
gas pressure results in less scattering of the particles in the gas phase. Thus the
particles arrive with more energy at the substrate which can then be used for diffusion
processes. Therefore decreasing sputtering gas pressure moves the transition between
zones to lower substrate temperatures.
It can be expected that in order to arrange in MAX phase structure (see unit cell
Figure 7.2: Structure zone model
proposed by Movchan and Demchishin [48].
Figure 7.3: Structure zone model
proposed by Thornton [49].
7.2. EXPERIMENTAL SETUP
45
Fig. 6.1) the adatoms will need sufficient energy for surface diffusion. Thus an elevated
substrate temperature and low inert gas pressure are assumed as suitable deposition
parameters for a successful deposition of MAX phase thin films.
7.2
Experimental Setup
For all experiments the sputtering device L 560 (Leybold AG) was used. Fig. 7.4
shows the arrangement of the cathodes and the substrate in the vacuum chamber. The
substrate is situated above the two cathodes at a distance of 75 mm to the target
surfaces. The cathodes were tilted such that a perpendicular on the center of the
target surfaces pointed at the center of the substrate. If both cathodes are used this
setup causes a concentration gradient across the substrate and a range of chemical
compositions can be investigated on the same substrate. This combinatorial approach
accelerates the deposition parameter search for a specific chemical composition and
structure [50].
Round targets with a diameter of 90 mm were used. Metallic and compound targets
were sputtered with a direct current (DC) power supply. For the non-conductive AlN
target a alternating power supply at radio frequency (RF) was used. The sputtering
processes were run at constant power. The power supplies to the two cathodes could
be controlled individually. When only one cathode was used the other cathode was
covered to avoid the deposition of material on the inactive cathode, but the positions
remained unchanged.
Initial tests with silicon wafers as substrates showed a considerable amount of silicon
diffusion into the thin films at substrate temperatures around 1073 K. A formation of
silicon containing phases in the films was observed. Therefore Si wafers were only used
at substrate temperatures below 800 K. At higher temperatures 15x15 mm alumina
plates were chosen as substrates because of their high temperature stability. The alu-
Figure 7.4: Arrangement of cathodes and substrate in the vacuum chamber.
46
CHAPTER 7. THIN FILM GROWTH
mina plates were made of polycrystalline and unpolished material. Selected samples,
e.g. for nanoindentation, were deposited on polished surfaces of single crystalline alumina plates with a root mean square roughness of ≤ 0.7 nm. It proofed experimentally
that similar results on single and polycrystalline substrates could only be achieved, if a
polycrystalline plate was placed adjacent to the single crystalline substrate inbetween
single crystal and heating system. This may be due to different thermal properties of
single and polycrystalline material in combination with the strong temperature dependence of the thin film structure.
The heating system is situated behind the substrate holder. It is capable of heating the
substrate to a temperature of 1273 K. To determine the substrate temperature during
deposition a calibration experiment was performed. For a series of heating settings
the temperature of a graphite coated silicon wafer substrate was measured with a pyrometer. Six points were measured within a temperature interval from 973 to 1273 K
showing a linear relationship that allows for interpolation. The effect of plasma heating
was assumed to be negligible at these high temperatures.
The vacuum system was pumped to a base pressure of 5x10−6 mbar. During deposition
the pressure was 6x10−3 mbar. Argon was used as inert gas and nitrogen gas could be
added for reactive sputtering.
Chapter 8
Characterization Methods
The methods employed for the analysis of the chemical composition and the structure
of the thin films were x-ray photoelectron spectroscopy (XPS) and x-ray diffraction
(XRD). Additionally wavelength dispersive x-ray analysis (WDX) has been performed
on selected samples to obtain spatially resolved chemical composition information.
Mechanical properties, e.g. Young modulus, have been measured by nanoindentation.
8.1
XRD Analysis of the Crystal Structure
The crystal structure of the thin films has been determined by XRD with a Siemens
D 500 diffractometer. A detailed description of diffraction and related topics can be
found in [51]. This section will provide a short description of the grazing incidence
technique which was used in this work. Generally for diffraction peak positions and
intensities of the measured signal depend among others on symmetry, dimensions,
and atomic positions of the unit cell of the sample. They are characteristic for each
crystalline material and the material and its structure can be identified by comparison
to standards. The correlation between the peak positions and the characteristic dvalues (atomic plane spacings) of the analyzed samples is given by Bragg’s law:
d=
nλ
,
2sinθ
(8.1)
where n is an integer. The principle of Bragg’s law is illustrated in Fig. 8.1.
For the grazing incidence technique the incident angle α of the x-rays on the sample
surface is constant and small (here α = 3◦ ). This ensures a lower penetration depth
of the x-rays and amplifies the signal received from the thin film compared to the
underlying substrate. Cu-Kα radiation with a wavelength of λ=1.54186 Å was used.
A beam of this radiation is emitted by the x-ray source and hits the sample at the
adjusted incident angle. The diffracted beam is then collected by a detector. The
detector moves around the sample on the 2θ orbit at constant velocity, see Fig. 8.2.
47
48
CHAPTER 8. CHARACTERIZATION METHODS
Figure 8.1: A Bragg reflection from a particular family of lattice planes, separated by a
distance d. Incident and reflected rays are shown for the two neighboring planes. The
path difference is 2d sinθ.
Figure 8.2: X-ray diffractometer.
8.2. XPS AND WDX ANALYSIS OF THE CHEMICAL COMPOSITION
8.2
49
XPS and WDX Analysis of the Chemical Composition
XPS is a technique to determine the chemical composition of a material and the binding states of its components. The basic concepts of this technique are discussed in
detail in [52]. Photons of the arriving x-rays interact with the sample which leads
to the ejection of photons, auger-electrons or photoelectrons from the sample surface.
The kinetic energy Ekin of these photoelectrons is then measured by an electron-energy
analyzer. The kinetic energy only depends on the excitation energy of the x-rays hν
and the binding energy of the electron inside the atom Ebin :
Ekin = hν + Ebin .
(8.2)
So the binding states of the electrons, which depend on the chemical composition of
the sample, can be directly determined [53].
The spectrometer Nanoscan 50 with the electron-energy analyzer MAC 3 from Cameca
was used. The excitation was produced by Mg-Kα radiation with a mean energy of
1253.6 eV.
This technique is extremely surface sensitive, because the photoelectrons are ejected
only from the first few atomic layers of the irradiated sample. Since the measurements
had to be conducted ex-situ, surface impurities, especially from hydrocarbons, could
not be avoided. The overall measurement error is in the order of 3-5 at pct. For a more
accurate measurement of the carbon content wavelength dispersive x-ray analysis was
conducted on selected samples in addition to the XPS measurements.
The WDX analysis was done at the Gemeinschaftslabor für Elektronenmikroskopie of
RWTH Aachen University. The device was a Cameca SX 100 equipped with four wavelength dispersive spectrometers with primary energies of 10 and 15 keV, respectively.
The current was 100 nA and the relative measuring accuracy under these conditions is
1-5 at pct. For the metals chromium, titanium, and aluminum the pure elements were
used as standards, while the standards for carbon, oxygen, and nitrogen were TiC,
Al2 O3 , and Fe4 N, respectively.
8.3
Nanoindentation Measurements of the Mechanical Properties
Nanoindentation is a method to determine plastic and elastic properties on the nanometer scale. This technique is useful for thin film analysis because of the low penetration
depth of the diamond tip. The applied force and resulting penetration depth of a diamond tip into a sample is measured. This yields a load-displacement curve which is
characteristic for the material. Mechanical properties of the material, e.g. hardness
and reduced elastic modulus, can then be derived from its load-displacement curve.
According to the method of Oliver and Pharr the reduced elastic modulus Er can be
50
CHAPTER 8. CHARACTERIZATION METHODS
calculated as follows:
√
π
Er = p
S,
2 A(hc )
(8.3)
where S denotes the stiffness and A(hc ) the contact area under peak load (see
Fig. 8.3) [54]. The reduced elastic modulus Er and the Young modulus E of the
sample are correlated by the following equation:
1
1 − v2
1 − v2
=(
)sample + (
)indenter ,
Er
E
E
(8.4)
where v denotes the Poisson ratio.
Figure 8.3: Cross section of an indent before, during, and after indentation.
The hardness H is defined as ratio of load to area. According to Oliver and Pharr
the respective definition of hardness for nanoindentation may be written as ratio of
maximum load Fmax to contact area under peak load:
H=
Fmax
.
A(hc )
(8.5)
The used device was a Triboindenter (Hysitron Inc.). The indents were made with a
so called Berkovich tip which has pyramidal shape with an included tip angle of 143.2◦
between its three faces.
Chapter 9
Results and Discussion
MAX phases are a group of materials that combine metallic and ceramic properties
to a previously unknown extend and a high effort is made to explore their potential
applications in industry and domestic life. So far for all three phases investigated in
this work research still focuses on the synthesis of single phase material from multiple
targets. The results presented in the next sections contribute to bridging the gap from
the state of the art research-lab scale multiple target deposition to an industrially
applicable large area deposition of MAX phase thin films. Therefore three challenges
need to be addressed:
• the deposition of single phase coatings,
• the homogeneous deposition over a large area,
• the deposition at low temperature.
The three MAX phases investigated in this work are Ti2 AlC, Cr2 AlC, and Ti2 AlN and
in the following section results will be presented for each phase.
9.1
Ti2AlC
Ti2 AlC is one of the few MAX phases which have been successfully synthesized as
single-phase bulk material (impurities ≤ 4 vol pct) [55]. Therefore a number of
bulk properties are known for this phase, e.g. thermal-expansion coefficient (8.2x10−6
◦ −1
C [55]), electrical conductivity (2.7x10−6 Ω−1 m−1 [55]), Vickers hardness (≈ 4.5
GPa [55]), stress-strain curves [55], and oxidation behavior [56].
For the deposition of Ti2 AlC thin films only one experiment is reported in the literature.
Wilhemsson et al. grew epitaxial films on MgO substrates by magnetron sputtering
from individual targets [44]. They reported a strong temperature dependence for the
formation of the MAX phase. It only formed at substrate temperatures above 973 K.
Their films contained a considerable amount of ancillary phases.
The experiments conducted in this work differ from the before mentioned in the use of
a compound target and the non-epitaxial growth mode. Different experimental setups
51
52
CHAPTER 9. RESULTS AND DISCUSSION
were used in order to approach the goal of the deposition of single phase material from
a single target and the results are reported in the following sections.
9.1.1
Compound Target
The bulk material for the compound target was supplied by 3-one-2 LLC [57]. Due
to its unique combination of ceramic and metallic properties this ceramic could be
machined into a round target (90 mm diameter, 6mm thick) and screwed to a copper
back-plate. The x-ray diffraction data in Fig. 9.1 shows that the target does not
consist of phase pure Ti2 AlC. The peak positions indicate a phase mixture with two
MAX phases, Ti2 AlC and Ti3 AlC2 , predominantly present.
Figure 9.1: X-ray diffraction data of the Ti2 AlC compound target material.
The chemical composition of the compound target was determined by a WDX linescan
shown in Fig. 9.2. The large deviations in compositions along the scanned direction,
observed best for aluminum, confirmed that the target material was not homogeneous.
The average composition calculated from the data amounted to 44 at pct titanium,
26 at pct carbon, 25 at pct aluminum, and 5 at pct oxygen. The target material
has a Ti/Al ratio of 1.76. Considering a relative measurement error of 2-5 pct, the
average composition of the target lacks approximately 5 at pct titanium compared to
the required 2:1:1 stoichiometry.
For the deposition experiments the compound target was run at a power density of
3.14 W/cm2 and with an argon pressure of 6x10−3 mbar. The deposition time was 30
min. The first experiment was conducted without additional substrate heating which
results in a substrate temperature of ≤ 500 K. A single-crystalline Si wafer was used as
substrate. X-ray diffraction data for this thin film is shown in Fig. 9.3 at the bottom.
The broad peaks indicate that the crystal size is very small. This is consistent with the
9.1. TI2 ALC
53
Figure 9.2: WDX data of the Ti2 AlC compound target material.
basic concepts behind the structure zone models discussed in chapter 7. Ti2 AlC melts
incongruently at Tm = 1898 ± 10 K [60] and a substrate temperature of T = 500 K is
not exceeded at the used deposition parameters without extra heating. Therefore the
ratio T/Tm is ≤ 0.3. According to the structure zone models this region is classified as
Zone 1, where surface diffusion is very limited. Due to the low deposition temperature
the particles have little mobility when they reach the substrate surface, which results
in amorphous films or films with very small crystal size. Most peaks can be attributed
to the presence of TiC. This agrees with observations made by Wilhelmsson et al [44].
They observed only TiC peaks for films deposited at a substrate temperature of 573
K using similar deposition parameters that yielded MAX phase films at substrate
temperatures above 1023 K. A Ti/Al ratio of 0.85 was determined by XPS for the
film in this work. Wilhelmsson et al. suggested either a metastable solid solution of
the Al in the TiC structure or the formation of Al-rich precipitations. From the XRD
data it can be seen that the measured peaks are slightly shifted towards greater angles
compared to the symbols indicating the peak positions of pure TiC according to JCPDS
data (see Fig. 9.3). The peak shift corresponds to a 2 pct change in lattice parameter
of the cubic phase. The measured peak positions indicate a smaller lattice parameter
compared to stoichiometric TiC, which may be due to the incorporation of Al into the
TiC structure and the formation of a metastable solution. A comparable shift has been
observed e.g. for the metastable solution of Al in the cubic TiN phase [58].
A second film was deposited under equal conditions, but at a substrate temperature
of 1153 K. Due to the high temperature an alumina substrate had to be used to avoid
the diffusion of silicon into the film. X-ray diffraction data are shown in Fig. 9.3 at
the top. Besides the alumina substrate peaks several diffraction peaks are present that
can be attributed to the film. Most likely these peaks are diffractions from Ti3 AlC2 ,
54
CHAPTER 9. RESULTS AND DISCUSSION
Figure 9.3: X-ray diffraction data of thin films deposited with the Ti2 AlC compound
target at 400 K and 1153 K.
Al3 Ti, and TiC crystals. No clear evidence was found for the presence of the Ti2 AlC
MAX phase.
A series of experiments was conducted varying the deposition pressure in order to
investigate whether the change in chemical composition from the target to the film was
associated to the particle transport through the discharge gas. The effect of pressure
on the sputtering process can be understood in terms of collisions of the sputtered
particles before they arrive at the substrate. The number of collisions is small or zero
in a low pressure atmosphere and the sputtered particles preserve their energy until they
arrive at the substrate. At higher pressure an increasing number of collisions causes
a thermalization of the sputtered flux. Thermalization means that an incremental
fraction of the sputtered flux has a low thermal energy [59]. Significant thermalization
can be expected if approximately 5 collisions per atom take place [61]. The effect of
the energy of the arriving atoms on film morphology has been described by Thornton
(see Fig. 7.1) [49]. For multicomponent sputtering it must be taken into account
that particles from different components have different weights and collisional cross
sections. Therefore they also have different spatial scattering properties and the number
of collisions also effects the chemical composition of the sputtered flux arriving at the
target.
The pressure in the chamber was varied from 0.2 Pa to 6 Pa at a constant temperature
of 1123 K. The compound target was operated at a power density of 3.14 W/cm2
for 10 min for each deposition. Fig. 9.4 shows the change of film composition with
9.1. TI2 ALC
55
pressure measured with XPS. The C content is not displayed because its measurement
is affected by C surface impurities. However, measurements on several selected samples
where the surface layers had been removed by ion etching confirmed that the ratio of
Al/C is 0.86±0.04. Fig. 9.4 shows that no significant changes of the Ti/Al ratio can
be observed within the investigated pressure range. The top axis gives an estimated
average number of collisions between particles for the respective pressure, which was
calculated from the distance between plasma source and substrate divided by the mean
free path of the particles for the respective pressure and temperature.
Figure 9.4: Ti/Al ratio of thin films deposited at different Ar-pressures with the Ti2 AlC
compound target (black squares), Ti/Al ratio of the Ti2 AlC MAX phase (dashed line),
and Ti/Al ratio of the target (dash-dotted line).
Fig. 9.4 shows that the film composition differs notably from the target composition.
The Ti/Al ratio in the films is less than half of the Ti/Al ratio of the target indicated
by the dashed line. It can also be seen that this effect is independent of the pressure.
Therefore the change of the Ti/Al ratio does not seem to be associated to the particle
transport through the discharge gas.
According to the conducted experiments a deposition of Ti2 AlC MAX phase thin films
from a Ti2 AlC compound target seems not a viable option under the tested conditions. Therefore further experiments were conducted with the compound target and
an additional target of elemental Ti.
9.1.2
Compound Target + Ti Target
The additional Ti was supplied by an elemental target of 99.9 pct purity. The two
cathodes were arranged as shown in Fig. 7.4.
56
CHAPTER 9. RESULTS AND DISCUSSION
In a series of experiments the necessary operating power for the Ti target was determined in order to yield the 211 stoichiometry of the MAX phase on the substrate.
A number of thin films were deposited at constant temperature (1123 K), pressure
(3x10−6 mbar), and power density of the Ti2 AlC compound target (3.14 W/cm2 ),
while the power density of the Ti target was varied from 1.57 to 3.14 W/cm2 . Fig. 9.5
shows the Ti/Al ratio of the thin films as measured by XPS. As the Ti target power
density is increased, the Ti/Al ratio increases exponentially. The desired Ti/Al ratio
for the MAX phase stoichiometry is indicated by a dashed line.
Figure 9.5: Ti/Al ratio of thin films deposited at different power densities at the Ti
target using both a Ti2 AlC compound target and a Ti target (black squares), and
Ti/Al ratio of the Ti2 AlC MAX phase (dashed line).
Fig. 9.6 shows the x-ray diffractogram for the respective film deposited on a singlecrystalline alumina substrate. Approximately 90 pct of the integrated intensity can
be attributed to the presence of the Ti2 AlC MAX phase. The origin of the three
remaining diffraction peaks cannot be determined unambiguously. Smaller amounts of
Ti2 AlC as part of a phase mixture can be found in all films with a Ti target power
density between 1.89 and 3.14 W/cm2 . The lattice parameters a and c have been
determined from the (100) and (103) diffraction peaks, respectively. No systematic
change of lattice parameters with chemical composition could be observed and the
average equilibrium volume was determined to be 108.2±0.6 Å3 . The deviation is less
than 3 pct compared to our ab initio calculations and the JCPDS data for Ti2 AlC,
suggesting excellent agreement between theory and experiment [34, 62].
With this set of optimized parameters a second series of experiments was conducted in
order to investigate the influence of the substrate temperature on the chemical composition of the thin films. The substrate temperature was varied from approximately
373 K (no heating applied) to 1123 K. Fig. 9.7 shows the Ti/Al ratio measured by
9.1. TI2 ALC
57
Figure 9.6: X-ray diffraction data of a thin film deposited with optimized parameters
using both a Ti2 AlC compound target and a Ti target at 1123 K.
XPS as a function of the substrate temperature. A linear decrease of the Ti/Al ratio with increasing substrate temperature can be observed for a temperature range of
373 - 1123 K. The Ti/Al ratio between 1123 and 1273 K is constant within the error
of the measurement.
Figure 9.7: Ti/Al ratio of thin films deposited at different temperatures using both a
Ti2 AlC compound target and a Ti target.
58
CHAPTER 9. RESULTS AND DISCUSSION
The dependency of film composition on the substrate temperature would support the
assumption that the change in chemical composition does not seem to be associated
to the particle transport through the discharge gas but rather to processes at the substrate. Possible reasons may be Ti desorption from the substrate. Understoichiometric
incorporation of Ti was also observed for the sputtering of other Ti compounds, e.g.
TiB2 [63, 64]. Another source for the changing Ti/Al ratio could be the alumina substrates. Due to the high solubility of oxygen in Ti, interdiffusion between alumina and
Ti was reported in the literature for temperatures between 1073 and 1373 K [65, 66].
If this may also be applicable to Ti2 AlC and the alumina substrates used in this work
needs further investigation.
9.1.3
Single Target of Ti and Ti2 AlC Wedges
For the coating of large areas the deposition from a single target is the simplest way
to achieve a homogeneous coating. Therefore a target was constructed from wedges of
the metallic Ti and Ti2 AlC compound target material used in the previous chapter.
Fig. 9.8 shows the round target segmented like a cake into 16 wedges. Each wedge was
screwed to the copper back plate and could be replaced individually.
Figure 9.8: Segmented target consisting of a variable number of Ti and Ti2 AlC wedges.
A series of experiments was conducted varying the number of Ti2 AlC compound wedges
contained in the target. Based on the ratio of the power densities needed to deposit
MAX phase thin films of high purity from the individual targets in the previous chapter,
2,4,6, and 8 wedges of the compound target were replaced by metallic Ti wedges. To
ensure best possible homogeneity of the films the wedges were arranged as shown
in Fig. 9.9. All films were deposited at an argon pressure of 6x10−3 mbar and a
substrate temperature of 1123 K. The deposition time was 17 min at a power density
of 3.14 W/cm2 . Polycrystalline alumina substrates were used.
9.1. TI2 ALC
59
2
4
6
8
Figure 9.9: Position of the metallic Ti (dark wedges) for experiments with 2,4,6, and
8 wedges of Ti in the segmented Ti2 AlC compound target.
Fig. 9.10 shows a representative section of the x-ray diffractograms of the resulting thin
films. It can be seen that the Ti2 AlC MAX phase peaks reach the highest intensities
for 4 or 6 wedges of Ti. When only 2 wedges of Ti were inserted in the target the MAX
phase peaks are smaller and the film also shows diffraction peaks from an intermetallic
phase, most likely Al2 Ti. Using 4 wedges of Ti the peaks from ancillary phases are
smaller and could not be uniquely identified. When 6 wedges were used ancillary phase
peaks are still very small and can most likely be attributed to AlTi3 . For 8 wedges
Figure 9.10: X-ray diffraction data for thin films deposited with the segmented Ti2 AlC
compound target with a) 2 Ti wedges b) 4 Ti wedges c) 6 Ti wedges d) 8 Ti wedges.
60
CHAPTER 9. RESULTS AND DISCUSSION
they become more pronounced and other unidentified peaks arise. Thus the purest
film with respect to the MAX phase is the film were 4 wedges of Ti were used for the
deposition. Films deposited with less Ti in the target show aluminum rich impurity
phases, whereas films deposited with more Ti contain titanium rich impurity phases.
The presence of small amounts of TiC in all films is very probable and cannot be
excluded due to overlapping peaks with the substrate.
The deposition with 4 Ti wedges was repeated on a single crystalline alumina substrate
to eliminate substrate peaks in the x-ray data and determine the phase purity of the
film. The substrate temperature was 1173 K and the target was sputtered for 20 min
with the remaining parameters unchanged from the former experiments. The resulting
x-ray data are shown in Fig. 9.11 and a phase purity with more than 95 pct Ti2 AlC
MAX phase was determined from the relative peak intensities. Therefore a segmented
target of Ti and Ti2 AlC wedges is a suitable single target for the deposition of phase
pure and homogeneous Ti2 AlC MAX phase thin films.
Figure 9.11: X-ray data for a film deposited at a substrate temperature of 1173 K on
a single crystalline alumina substrate with the segmented target consisting of 4 Ti and
12 Ti2 AlC wedges.
Most substrate materials require rather low deposition temperatures for their properties to remain unchanged during the coating process. Therefore in a last series of
experiments the substrate temperature was decreased from 1173 K down to 873 K in
steps of 100 K. Fig. 9.12 shows how the intensity of the strongest Ti2 AlC peak at
2θ= 39 ◦ decreases drastically with decreasing substrate temperature. At 1073 K the
intensity of the MAX phase peaks has decreased by half and at 2θ= 36 ◦ a peak from
an ancillary phase becomes clearly visible. At 973 K broader peaks of even lower intensities can be observed. The ancillary peak at 2θ= 36 ◦ is now of equal intensity as the
MAX phase peak at 2θ= 39 ◦ . The peak broadening is caused by a smaller grain size
which can be understood based on the structure zone models introduced in section 7.1
9.1. TI2 ALC
61
Figure 9.12: X-ray data for films deposited at different substrate temperatures with
the segmented target consisting of 4 Ti and 12 Ti2 AlC wedges.
as a result of the lower substrate temperature. At 873 K peaks have become so broad
and of low intensity that they cannot be clearly distinguished anymore. This film is
probably nanocrystalline with a nearly amorphous film structure. A similar temperature dependence of the Ti2 AlC MAX phase formation was observed by Wilhelmsson et
al. [44]. They observed no Ti2 AlC at 973 K substrate temperature, a reduced peak of
Ti2 AlC in the x-ray diffractogram at a substrate temperature of 1073 K and a strong
peak at 1173 K. Thus their experiments show a similar trend but for slightly different substrate temperatures which may be due to a different experimental setup or
inaccurate temperature measurement.
Since it has been shown earlier in this work that the chemical composition of the Ti2 AlC
films changes with temperature (Fig. 9.7), the composition of the thin films shown in
Fig. 9.12 was analyzed by EDX. For the substrate temperatures from 873 K up to
1073 K a constant Ti/Al ratio of approximately 2.4 was determined, while the film
with a substrate temperature of 1173 K had a Ti/Al ratio of 2.1. This ratio lies closer
to the MAX phase stoichiometry and this observation agrees with the phase purity of
more than 95 pct MAX phase for this film compared to approximately 87 pct for the
film with a substrate temperature of 1073 K. Towards lower temperatures the phase
purity decreases even further, but has not been quantified due to peak broadening and
overlapping.
Furthermore it is interesting to note that the linear dependency of the Ti/Al ratio on
the temperature observed for the setup with two separate cathodes of Ti and Ti2 AlC in
Fig. 9.7 was not found in this experiment with the segmented target. Instead a jump of
the Ti/Al ratio was observed from a value of approximately 2.1 at 1173 K to a value of
approximately 2.4 at 1073 K which then remained constant down to a temperature of
62
CHAPTER 9. RESULTS AND DISCUSSION
873 K. The latter was measured by EDX, while the linear dependency was measured by
XPS. XPS is a surface sensitive technique, which only measures the composition of the
topmost atomic layers, while EDX has a higher penetration depths (500 nm here) into
the samples. Thus different experimental setups and analysis techniques have to be
taken into account, but the true nature of the temperature dependence of the thin film
composition for Ti2 AlC has not been understood yet and needs further investigation.
However, so far an acceptable phase purity could only be achieved at temperatures
above 1123 K, which is too high for a deposition on e.g. steel. Future work could
investigate the nature of the temperature dependence of the thin film composition and
the application of a bias potential to the substrate to lower the deposition temperature.
9.1.4
Compound Target with Cathodic Arc
As an alternative to magnetron sputtering the deposition with the cathodic arc technique was tested in a short experiment. Detailed information about the cathodic arc
technique can be found in [67]. The main difference to magnetron sputtering lies in
the plasma formation. In cathodic arc the plasma is formed through a high current
discharge at discrete micron-sized sites, called cathode spots, at the target surface.
The resulting plasma is characterized by a high degree of ionization (up to 100 pct
compared to less than 10 pct for sputtering). The particles in the arc discharge possess
kinetic energies in the range of 20-160 eV, compared to around 5 eV for sputtered
particles.
The thin film growth was performed using a filtered cathodic arc system with a target
to substrate distance of 454 mm. The top of a conical aluminum cathode (base and
top diameters 51 and 12 mm, respectively, and height 38 mm) has been replaced by the
same bulk Ti2 AlC material as used for the sputtering target, see Fig. 9.1.4. The arc
was generated with a DC power supply with a resulting arc current of 36 A. A straight
filter was used consisting of a coil with 640 turns and a coil current of 6 A. A 10x10x1
mm single-crystalline alumina substrate was used. The substrate was heated to a tem-
Figure 9.13: Ti2 AlC cathode for the
cathodic arc experiment.
Figure 9.14: Eroded Ti2 AlC cathode
after the cathodic arc experiment.
9.2. CR2 ALC
63
perature of 1073 K which was measured with a K-type (NiCr/NiAl) thermocouple.
Before deposition the vacuum system reached a base pressure of 7.5x10−6 mbar. A deposition time of 317 s was reached at an argon pressure of 6x10−3 mbar before the arc
extinguished. The porosity of the Ti2 AlC material was detrimental for the cathodic arc
process, because the cathode spot got trapped in cavities below the target surface and
extinguished. As a consequence the cathode surface was not homogeneously eroded
but shows cavities and pores as can be seen in Fig. 9.1.4.
The chemical composition of the thin film was determined by elastic recoil detection
analysis (ERDA) with 40 MeV Cl7+ ions having an angle of 15 ◦ relative to the sample
surface. Details about the ERDA set-up can be found elsewhere [69]. The measured
film composition was 58 at pct aluminum, 32.6 at pct titanium, 5 at pct carbon, 3.1 at
pct oxygen, and 1.1 at pct nitrogen with a standard deviation of 4 pct and a sensitivity
of 0.5 pct. The film shows a strong deficiency in titanium compared to the target. The
arc-deposited film has a Ti/Al ratio of 0.56. A comparable trend was observed by
XPS for the magnetron sputtering process. Yielding similar results for two different
techniques supports the assumption that the understoichiometric titanium content is
due to processes at the substrate. Furthermore the content of carbon is very low in the
arc-deposited film compared to the target and the magnetron sputtered film. This may
be due to scattering in the gas phase. Both (arc and magnetron sputtering) deposition
processes were run at the same pressure but the target to substrate distance in the arc
setup is 6x as large as in the sputtering setup. Carbon is the lightest of the evaporated
elements and will be affected most by the increased number of collisions. Additionally
the arc process causes much stronger resputtering effects at the substrate because the
arriving particles carry higher energies compared to sputtered particles. This may also
be the reason why the x-ray diffraction of the arc-deposited films showed no signs of
a crystal structure. It is reported in the literature that the impact of high energy
particles on the substrate leads to the formation of amorphous films of an otherwise
crystalline material because high energy collision cascades destroy the crystal structure
of the underlying film material [68].
9.2
Cr2AlC
Compared to Ti2 AlC less is known about the Cr2 AlC MAX phase. It is reported in
the literature that single phase samples (impurities ≤ 5 vol pct) of the bulk material were produced successfully [70]. However, except for the heat capacity (139 J
mol−1 K−1 [70]) no material properties have been measured so far. A bulk modulus
B of Cr2 AlC was reported from theoretical (B = 225 GPa) and semi-empirical studies
(B = 228 GPa) [42, 70].
Thin films of single phase Cr2 AlC MAX phase were deposited successfully by magnetron sputtering from three elementary targets [42, 43]. With the focus on large
area deposition of homogeneous thin films, in this work the deposition from a single
compound target with Cr2 AlC stoichiometry was investigated.
64
9.2.1
CHAPTER 9. RESULTS AND DISCUSSION
Compound Target
The bulk material for the compound target was supplied by 3-one-2 LLC [57]. Electrical discharge machining had to be used to fabricate the target (90 mm diameter, 3
mm thick) from the bulk material because ultra hard inclusions inhibited mechanical
machining. The block of bulk material was cut into plates that were attached to a
copper back plate with aluminum screws. As can be seen in Fig. 9.15 the screws had
to be fixed with aluminum nuts on the target surface, because no screw thread could
be tapped into the material. To avoid sputtering of the nuts they were attached at
the maximum possible distance to the racetrack (8 mm). The chemical composition of
the compound target was determined by WDX to be 49 at pct chromium, 25 at pct
aluminum, 24 at pct carbon, and 2 at pct oxygen. This composition agrees well with
the desired Cr:Al:C ratio of 2:1:1. The x-ray diffraction data in Fig. 9.16 shows that
most peaks coincide with the peak positions of Cr2 AlC indicated by the black squares.
Only two diffraction peaks at 2θ = 42 ◦ and 2θ = 67 ◦ result from impurities and ≥ 90
pct of the integrated intensity can be attributed to the presence of the Cr2 AlC MAX
phase.
For the deposition experiments the compound target was run at a power density of
5.5 W/cm2 and with an argon pressure of 6x10−3 mbar. The deposition time was 30
min. Polished single crystalline alumina substrates were used at a substrate temperature of 1123 K.
X-ray diffraction data of the resulting film are shown in Fig. 9.17. Most peaks can be
attributed to the presence of the Cr2 AlC MAX phase. The residual diffraction peaks
caused by ancillary phases are very small. In comparison to the target material the
ancillary phase peaks of the thin film are smaller and at distinct positions. According
to the integrated intensity the thin film consists of ≥ 90 pct pure MAX phase. The
Figure 9.15: Cr2 AlC compound target constructed of bulk material plates attached to
a copper back plate with aluminum screws and nuts.
9.2. CR2 ALC
65
Figure 9.16: X-ray diffraction data of the Cr2 AlC compound target material.
chemical composition of the thin film was measured by WDX. It consisted of 46 at pct
chromium, 27 at pct aluminum, 24 at pct carbon, and 3 at pct oxygen. A decrease of
3 at pct in chromium compensated by an increase of 2 at pct in aluminum and 1 at
pct in oxygen can be observed compared to the target composition.
The Young modulus of the Cr2 AlC thin films was measured by nanoindentation. After
Figure 9.17: X-ray diffraction data of a thin film deposited with the Cr2 AlC compound
target at 1123 K.
66
CHAPTER 9. RESULTS AND DISCUSSION
deposition the film surface was rough with height differences of up to 370 nm. Before
the indentation films were polished with a solution containing fused alumina and a final
roughness of less than 10 nm was reached. A series of 54 indents was performed with
a force of 500 µN and a resulting maximum indentation depth of 38.8 nm. A Young
modulus of 178 GPa was determined with a standard deviation of 8.8 pct. Compared to
the theoretical value of 357.7 GPa the measured value is very low [35]. Presumably this
may be due to the low density of the films. A strong correlation between film density
and elastic properties has been reported in the literature for e.g. boron suboxide thin
films [71]. Due to the formation of nano-pores a decrease in film density by a factor
of 1.55 caused an elastic modulus reduction by a factor 4.5. Thus the reduction of the
Young modulus by a factor 2 observed for Cr2 AlC may be due to a comparatively small
deviation in film density due to the formation of pores. Two deposition parameters
that affect the film density are substrate temperature and substrate bias potential.
Especially the use of the latter to improve the film density will be explored in future
work.
Since a high deposition temperature like 1123 K may affect the substrate material, it is
desirable to run the deposition process at lower temperatures. Consequently the substrate temperature was decreased stepwise in a series of experiments and the resulting
films were analyzed with XRD. Fig. 9.18 shows that the Cr2 AlC MAX phase is stable
and crystalline down to a deposition temperature of 723 K. Below this temperature
at 623 K the film structure appears to be amorphous. In the whole temperature interval the phase purity remains above 90 pct and no significant amounts of ancillary
phases appear. Compared to the Ti2 AlC MAX phase this is a remarkable result. Phase
pure Ti2 AlC could only be observed at substrate temperatures above 1123 K and the
Figure 9.18: X-ray diffraction data of thin films deposited with the Cr2 AlC compound
target at different substrate temperatures.
9.3. TI2 ALN
67
films lost their crystalline structure at 873 K, which is 200 K above the respective
temperature for the Cr2 AlC films.
Compared to the Ti2 AlC MAX phase Cr2 AlC shows no significant deviation of the
chemical composition for target and thin film. Therefore Cr2 AlC seems suitable for a
large area deposition from a single compound target. The low deposition temperature
of 723 K allows for the coating of a large variety of materials, e.g. steel. What still
needs to be done is an experimental investigation of the properties of this material. A
first suggestion is to investigate whether the low elastic modulus of the films can be
raised by increasing the film density.
9.3
Ti2AlN
Up to now all of the well-established MAX phase systems are carbides, e.g. Ti-Si-C,
Ti-Al-C, Cr-Al-C. Nitrides seem to be more difficult to synthesize, which may be one
reason why they are less explored. Theoretical calculations showed that the bulk modulus of M2 AlN MAX phases is generally higher compared to the corresponding carbide
MAX phase M2 AlC, which may be explained by there being an extra electron in the
nitrides, contributing to stronger chemical bonding [36]. Barsoum et al. reported that
despite their best efforts, samples of Ti2 AlN contained anywhere between 10 and 15
vol pct of ancillary phases [55]. They tried to synthesize bulk material by hot isostatic
pressing, the same technique that was used successfully for the preparation of the single phase carbide Ti2 AlC. They reported the thermal-expansion coefficient (8.8x10−6
◦ −1
C ), the electrical conductivity (4.0x10−6 Ω−1 m−1 ), Vickers hardness (≈ 4.0 GPa),
and stress-strain curves for their Ti2 AlN bulk material [55].
Thin films of Ti2 AlN were prepared by reactive magnetron sputtering from a sintered
Ti-Al target with a 2:1 composition ratio by Joelsson et al. [45]. Oriented MgO substrates were used for an epitaxial growth of the Ti2 AlN thin films. They report singlephase films for a nitrogen pressure of 0.19 mTorr with a hardness of 16.1 GPa and a
Young modulus of 270 GPa measured by nanoindentation, which is in good agreement
with the theoretical Young modulus of 293 GPa calculated by ab initio methods [46].
Epitaxial growth is not suitable for industrial deposition of large area thin films. Therefore this work could not build on any earlier experience with deposition of Ti2 AlN
MAX phase coatings. Since no sintered bulk Ti2 AlN material was available for use as
a compound target, experimental setups with multiple targets were also investigated.
Keeping in mind the long-term objective of large area deposition from a single target,
the first challenge was to find a systematic approach to the deposition of single-phase
Ti2 AlN MAX phase thin films.
9.3.1
Ti Target + Al Target + N2 Gas
A straight forward approach for the deposition of Ti2 AlN is the use of elemental targets
and nitrogen gas. The nitrogen gas will react with the target surface, which leads to
the contamination of the target surface with reactive compounds, called poisoning.
68
CHAPTER 9. RESULTS AND DISCUSSION
Often, poisoning changes the sputtering yield and depends on the partial pressure of
the reactive gas and the power supplied to the target. In the case of insulating reactive
compounds the sputtering mode of the target changes from metallic to dielectric mode
with increasing poisoning. During the reactive sputtering of more than one target the
poisoning of each target depends on the others due to different reactivities with respect
to the reactive gas. In poisoned mode this interaction of the targets is decoupled. The
power of one target can be decreased without changing the sputtering yield of the other.
Therefore it is advantageous to work in poisoned mode while adjusting the target power
densities for stoichiometric Ti2 AlN thin films. The major draw back of deposition in
poisoned mode is the low deposition rate. In a first experiment the poisoning of the two
elemental targets in a nitrogen atmosphere was investigated. The flow of nitrogen into
the chamber was varied from 0 to 50 ml/min monitoring the resulting voltage change
at both cathodes keeping power densities constant. Fig. 9.19 shows the resulting curves
for a power density of 0.79 W/cm2 at the Al target and 1.57 W/cm2 at the Ti target
and a constant argon pressure of 6x10−6 mbar. At a flow of 30 ml/min both targets
are sputtered in poisoned mode.
Next the power densities of the targets were adjusted in order to reach the desired 211
stoichiometry. Films were deposited at a substrate temperature of 773 K, an argon
pressure of 6x10−6 mbar, and a nitrogen flow of 30 ml/min. Through changing the
power densities at the targets and subsequent analysis of the chemical composition by
XPS the desired stoichiometry of the films was stepwise approached. A Ti/Al ratio
of 2 was reached for a power density of 1.26 W/cm2 and 3.77 W/cm2 for the Al and
Ti target, respectively. The nitrogen content of the films lay between 50 and 55 at
pct, independent of the varying power densities. A nitrogen content around 50 at pct
indicates the formation of the binary nitrides or metastable (Ti,AL)N. Increasing the
Figure 9.19: Effect of N2 poisoning on the target voltage.
9.3. TI2 ALN
69
substrate temperature to 1273 K did not change the nitrogen incorporation into the
film significantly.
The experiments showed that the nitrogen partial pressure is crucial for the phase formation during reactive sputtering in the Ti-Al-N system. Sufficient supply of nitrogen
at a Ti/Al ratio of 2 will favor the formation of TiN type phases over formation of the
MAX phase. This is the case when sputtering in poisoned mode from the elemental
targets. Therefore a systematic approach by reactive sputtering in the poisoned mode
seems not viable for the Ti2 AlN MAX phase.
9.3.2
Powder Target
Since reactive sputtering in the poisoned mode did not yield a controllable nitrogen
concentration, nitrogen was introduced to the sputtering process in the solid state as a
nitride. The use of a powder target seemed a promising setup for such an experiment.
By mixing AlN and Ti powder the chemical composition of the target could be adjusted
in order to yield the desired 211 stoichiometry in the thin films. The round target was
replaced by a stainless steel cup with the same diameter (90 mm). The cup was
filled with powder up to the rim with a height of 12 mm. First tests showed that
thin films deposited with this powder target contained up to 11 at pct of iron due to
plasma burning on the rim top. Therefore a ring was constructed that was positioned
approximately 1 mm above the rim of the cup. The ring was on ground potential and
prevented a plasma ignition on the ring, which solved the iron problem. The complete
construction of the powder target is shown in Fig. 9.20.
In first experiments sputtering of the AlN powder was tested. The powder was 99
pct pure. Films were deposited on Si substrates without substrate heating at an argon
Figure 9.20: Target with AlN powder and ring construction to prevent sputtering of
the powder cup.
70
CHAPTER 9. RESULTS AND DISCUSSION
pressure of 6x10−3 mbar. An RF generator was used and the power was slowly increased
carefully observing the powder behavior. At 400 W the surface layer of the white AlN
powder became black and cracks in the surface layer became visible. Consequently
a thin film was deposited running the powder target at 200 W at the same argon
pressure and a substrate temperature of 773 K. The sputtering rate of the AlN powder
was extremely low. After 145 min sputtering at 200 W a film thickness of ≤ 0.7 µm
was reached. X-ray diffraction data for this film are shown in Fig. 9.21. Based on
these data a clear identification of the existing phases was difficult. Some of the peak
positions agree with the existence of γ-Al2 O3 . However, the visible peaks exclude AlN
from being the main component of the film.
Figure 9.21: X-ray diffraction data of a thin film deposited with the AlN powder target.
Chemical analysis by XPS revealed that all films contained up to 60 at pct oxygen.
The source of this oxygen remains unclear. XPS analysis of the chemical composition of the powder was not possible due to missing facilities for the introduction of
a powder into the analysis chamber. X-ray diffraction of the AlN powder showed no
traces of impurities. All visible peaks could be attributed to AlN. The presence of
x-ray amorphous oxides cannot be excluded, but the powder itself seems an unlikely
source for the oxygen impurities in the thin films. More probably the oxygen stems
from the residual gases in the deposition chamber and incorporation into the films is
extremely high due to the low deposition rate of the AlN powder target. Air may also
be trapped in the powder target, because the cup is filled with powder and compressed
under atmospheric conditions.
It was also observed that the Al/N ratio of the thin films lay above 2. This observation
may be connected to the high content of oxygen or due to other reasons, e.g. poor
nitrogen adsorption at the substrate, which could not be clarified.
Deposition from a powder target has advantages that encourage further efforts to de-
9.3. TI2 ALN
71
velop this technique. This work explored only briefly its applicability to the deposition
of MAX phases. The main drawback encountered was the incorporation of an excess
amount of oxygen into the deposited thin films, which makes the use of a powder target
unsuitable for the deposition of phase pure Ti2 AlN.
9.3.3
Ti Target + AlN Target + N2 Gas
A third alternative to reactive sputtering and the use of a powder target is the deposition from a sintered nitride target. A 99.7 pct pure sintered AlN target (diameter
90 mm, thickness 5 mm) with a density of 98 pct was used in combination with the
elemental Ti target.
First only sputtering of the AlN target was examined. It was sputtered for 105 min at
a power density of 3.14 W/cm2 and an argon pressure of 6x10−3 mbar. The substrate
was a Si wafer heated to a temperature of 773 K. A Al/N ratio of ≥ 2 was measured
by XPS on the surface of the deposited thin film. Similar results were found for the
powder AlN target mentioned in the previous chapter. Both targets yield Al/N ratios
above 2 in the thin films but less oxygen is incorporated using the sintered target. This
was confirmed by the x-ray diffraction data for this film shown in Fig. 9.22. It shows
peaks at the same positions as the thin film deposited with the AlN powder target in
Fig. 9.21, but contains additional strong diffraction peaks that may be attributed to
the presence of AlN. The film had a thickness of 1.3 µm. Compared to deposition from
the powder target at a comparable power density the sputtering rate is 2.5 times as
high for the sintered target.
In order to supply the amount of nitrogen needed for the Ti2 AlN stoichiometry, nitrogen
Figure 9.22: X-ray diffraction data of a thin film deposited with the sintered AlN
target.
72
CHAPTER 9. RESULTS AND DISCUSSION
gas had to be added to the deposition process. Going back to reactive sputtering has
the disadvantage that now the sputtering yield of the metallic Ti target depends on the
amount of nitrogen gas used. This makes a systematic variation of the process parameters very difficult. However, a film deposited at a power density of 6.29 W/cm2 and
4.72 W/cm2 at the AlN and Ti target, respectively, an argon pressure of 6x10−3 mbar,
a nitrogen gas flow of 2.5 ml/min and a substrate temperature of 1123 K showed first
traces of the Ti2 AlN MAX phase in the x-ray diffractogram in Fig. 9.23. It can also be
seen that the predominant phase has a cubic TiN structure. Decreasing the nitrogen
gas flow and increasing the power density at the Ti target favored the formation of the
MAX phase in the thin film. Fig. 9.24 shows the diffractogram for a film deposited with
a power density of 5.50 W/cm2 at the Ti target and a nitrogen gas flow of 2 ml/min
(other parameters unchanged from previous deposition). More than 85 pct of the
integrated intensity can be attributed to the presence of the Ti2 AlN MAX phase. The
remaining two minor peaks indicate the presence of a small amount of a cubic TiN
phase.
Figure 9.23: X-ray diffraction data of a thin film deposited with the sintered AlN
target, metallic Ti target, and nitrogen gas, showing first traces of the MAX phase.
Improving the phase purity of the thin films was tried in various experiments. The
substrate temperature was increased up to 1273 K. The power density at the Ti target
was varied between 5.19 W/cm2 and 6.29 W/cm2 . The nitrogen gas flow was varied
between 0 and 7.5 ml/min. Several experiments were done varying both, nitrogen gas
flow and Ti target power density, in order to obtain the right stoichiometry. Although
the phase formation of the thin film was sensitive to all those parameters, the amount
of TiN impurities in the thin films could not be reduced any further. Similar problems
were reported by Barsoum et al., who could not reduce the amount of ancillary phases
to less than 15 vol pct in their bulk samples of Ti2 AlN[55, 72].
9.3. TI2 ALN
73
Figure 9.24: X-ray diffraction data of a thin film deposited with a power density of
6.29 W/cm2 and 5.50 W/cm2 at the AlN and Ti target, respectively, and a substrate
temperature of 1123 K .
Figure 9.25: X-ray diffraction data at five spots (4 mm maximum extension) along the
chemical gradient of the thin film shown in Fig. 9.24.
74
CHAPTER 9. RESULTS AND DISCUSSION
Nevertheless it has to be considered that the film shown in Fig. 9.24 has a gradient
in chemical composition due to the experimental setup with two cathodes. Since the
MAX phase is a stoichiometric phase only a small area of the coating can be expected
to be phase pure. The goniometer used for Fig. 9.24 measures an area of several mm2
which includes a range of chemical compositions. A second x-ray diffraction experiment
was done using a general area detector diffraction system (GADDS). With this device
the maximum extension of the measuring spot could be reduced to 4 mm at an angle of
θ = 9 ◦ . Fig. 9.25 shows the 2 θ scans for 5 points evenly distributed from side to side
on a line along the chemical composition gradient of the wafer (direction of increasing
Ti content indicated by the arrow). All peaks can be attributed to either the Ti2 AlN
MAX phase or cubic TiN. The content of TiN lies below 20 vol pct in all films and the
characteristic TiN peaks at 36.7 ◦ and 42.6 ◦ decrease with decreasing Ti content. The
second scan counting from the bottom up shows almost no traces of TiN. At the same
time the overall peak size also decreases with decreasing Ti content.
Fig. 9.25 indicates that the deposition of larger amounts of phase pure Ti2 AlN is
difficult with this experimental setup. Due to the lack of solubility ranges the chemical
composition gradient causes the formation of ancillary phases. The use of a rotating
substrate holder or a single target, e.g. made of wedges of Ti and AlN, may eliminate
this problem and is suggested for future work.
Chapter 10
Summary and Conclusions
The long-term goal that this work contributed to is an industrially applicable large area
deposition of MAX phase thin films. Therefore three challenges need to be addressed:
• the deposition of single phase coatings,
• the homogeneous deposition over a large area,
• the deposition at low temperature.
This work contributed to only one or all three of these challenges depending on the
material system. The investigated materials were the MAX phases Ti2 AlC, Cr2 AlC,
and Ti2 AlN and the chosen deposition technique was magnetron sputtering.
For all three phases process parameters were determined that allowed for the deposition of thin films consisting of over 85 pct single-phase MAX phase material.
In the case of Cr2 AlC a phase purity of over 90 pct could be achieved for the thin films.
They were deposited from a single target of Cr2 AlC MAX phase bulk material and the
chemical composition of the films did not differ significantly from the target composition. The lowest substrate temperature that still yielded phase-pure and crystalline
MAX phase films was determined to be 723 K. Thus for Cr2 AlC all three challenges
could be addressed successfully and a pathway for an industrially applicable large area
deposition was demonstrated. It remains an open question why the Young modulus
of the thin films, measured by nanoindentation, was only half of the theoretical value.
It is proposed that this may be due to low film density and will be subject to future
research.
Analogous to the successful deposition of Cr2 AlC thin films an experiment was conducted for the deposition of Ti2 AlC thin films from a single target of Ti2 AlC MAX
phase bulk material. In the case of Ti the chemical composition of the deposited thin
films differed significantly from the target composition. A strong lack of Ti in the
films was observed independent of the deposition pressure and the deposition technique (magnetron sputtering or cathodic arc). The Ti content in the films decreased
with increasing substrate temperature. Therefore desorption processes at the substrate
are believed to cause the Ti deficiency of the films. Consequently Ti was added to the
process from an additional target of elemental Ti and with this setup Ti2 AlC MAX
75
76
CHAPTER 10. SUMMARY AND CONCLUSIONS
phase thin films with a phase purity of over 90 pct could be deposited. The next step
was the construction of a single target that was segmented like a cake consisting of
compound Ti2 AlC and metallic Ti wedges. This target design enabled the deposition
of homogeneous Ti2 AlC MAX phase coatings with a phase-purity of over 95 pct at
a substrate temperature of 1173 K. Lower substrate temperatures lead to decreasing
phase-purity and the crystalline structure became undefined at 873 K. Therefore a
pathway for an industrially applicable large area deposition was found for Ti2 AlC, but
only at very high deposition temperatures. Future work may be directed towards decreasing the substrate temperature by e.g. application of a substrate bias potential.
The deposition of Ti2 AlN MAX phase thin films has not been published in the reviewed
literature. The phase was observed in bulk material, but no phase pure material could
be synthesized so far and no single target of Ti2 AlN MAX phase bulk material was
available. In this work several strategies were explored for the deposition of singlephase Ti2 AlN thin films. First deposition experiments were conducted using elemental
targets and nitrogen gas. Working with both metallic targets in poisoned mode resulted
in films containing an excess of nitrogen for the formation of the Ti2 AlN MAX phase.
Deposition from an AlN powder target was extremely slow and significant amounts of
oxygen were incorporated in the films, while the nitrogen content was understoichiometric compared to aluminum. Finally the deposition of Ti2 AlN MAX phase thin films
succeeded using a sintered AlN target in combination with a metallic Ti target and a
small amount of nitrogen gas. Since the elements are co-sputtered from two targets
the resulting films are not homogeneous but show a gradient in chemical composition.
Due to its small solubility range, only small amounts of phase pure Ti2 AlN could be
synthesized with this experimental setup. Thus in a first step suitable deposition parameters have been determined to deposit Ti2 AlN MAX phase thin films. The next
step will be to combine the target materials in a single target for the deposition of
homogeneous films over large areas.
Chapter 11
Future Work
The deposition of MAX phase thin films is a new and emerging research field and
this work only contributed to the very beginning of a detailed investigation of these
fascinating materials. Future work is abundant and some ideas that arose straight
from this thesis are resumed in this chapter. For Cr2 AlC the low Young modulus of
the films remains an open question. As a first assumption this may be caused by low
film density. To enhance the density and the related properties of the Cr2 AlC thin
films the application of a negative bias potential to the substrate is planned.
For Ti2 AlC the reason for the lack of Ti in thin films deposited from the target made
of Ti2 AlC MAX phase bulk material is still not well understood. Further insight could
be gained by an analysis of the particles arriving at the substrate using a mass energy
analyzer.
Future work could also include the use of a negative bias potential at the substrate
to lower the deposition temperature for crystalline Ti2 AlC MAX phase thin films.
Furthermore Young modulus and hardness of the films deposited in this work have not
yet been determined and compared to the calculated values.
Results from ab initio calculations exist for structure and elastic properties of M2 AlC
with M = Ti,V, and Cr. The deposition of V2 AlC thin films would allow for a complete
comparison and a systematic analysis of the influence of the number of valence electrons
on the elastic properties.
A similar ab initio study exists for the nitride MAX phases M2 AlN with M = Ti,V,
and Cr. The deposition of V2 AlN and Cr2 AlN thin films has not been reported in the
literature and is a challenging goal for future work.
For the deposition of homogeneous Ti2 AlN thin films over large areas in future work
a single target has to be constructed. As long as no compound material is available a
segmented target as used for the deposition of Ti2 AlC is suggested. Since it showed for
sputtering AlN that the Al/N ratio in the films was above 2, replacing the metallic Ti
in the target by TiN may be an option to avoid reactive sputtering with nitrogen gas.
Working with AlN and TiN the target will contain more nitrogen than needed for the
MAX phase formation, but the resulting films may still have the desired stoichiometry.
Additionally the nitrogen content could be influenced by the deposition pressure and
the application of a bias potential to the substrate.
77
78
CHAPTER 11. FUTURE WORK
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Curriculum Vitae
Name:
Claudia Walter
Place/Date of birth:
Moers, 11/04/1976
Nationality:
German
Education:
02/2001-04/2005
Doctoral student at RWTH Aachen (Materials Chemistry)
10/1995-12/2000
Diploma student at RWTH Aachen (Mechanical Engineering)
08/1999-06/2000
DAAD scholarship for Carnegie Mellon University, Pittsburgh, USA
(Department of Chemical Engineering)
08/1986-05/1995
High school student at Gymnasium Adolfinum, Moers
83