Contents
1.
2.
3.
4.
5.
6.
7.
Introduction.................................................................................................................................1
1.1.
Research goals ....................................................................................................................1
1.2.
Outline of the thesis ..........................................................................................................2
Hard coatings...............................................................................................................................3
2.1.
Titanium Nitride ................................................................................................................5
2.2.
Titanium Aluminium Nitride ...........................................................................................5
2.3.
Aluminium Nitride.............................................................................................................7
Multilayer coatings ......................................................................................................................9
3.1.
Structure ..............................................................................................................................9
3.2.
Mechanical properties .....................................................................................................10
3.3.
Multilayer coatings on cutting tool................................................................................12
Phase transformation and decomposition.............................................................................15
4.1.
Nucleation and growth....................................................................................................15
4.2.
Spinodal decomposition..................................................................................................16
Coating deposition....................................................................................................................23
5.1.
Cathodic arc evaporation................................................................................................23
5.2.
Multilayer growth.............................................................................................................26
Characterization ........................................................................................................................29
6.1.
Nanoindentation ..............................................................................................................29
6.2.
Thermal analysis and calorimetry ..................................................................................31
6.3.
Atom probe tomography ................................................................................................33
6.4.
Scanning electron microscopy .......................................................................................34
6.5.
Transmission electron microscopy................................................................................35
6.6.
Energy dispersive spectroscopy.....................................................................................36
6.7.
Wide angle x-ray scattering.............................................................................................36
6.8.
Small angle x-ray scattering ............................................................................................38
Phase-field simulations.............................................................................................................43
Contents
7.1.
The Cahn-Hilliard phase-field model............................................................................43
7.2.
Microstructure evolution of monolithic and multilayer Ti1-xAlxN ............................45
8.
Metal cutting ..............................................................................................................................47
8.1.
Conditions during cutting...............................................................................................47
8.2.
Wear mechanisms ............................................................................................................48
8.3.
Cutting performance of TiAlN coatings ......................................................................50
8.4.
Cutting performance of multilayer coatings ................................................................51
9.
Stabilization of c-Ti0.25Al0.75N ..................................................................................................53
9.1.
Deposition conditions.....................................................................................................53
9.2.
Microstructure ..................................................................................................................54
9.3.
Mechanical properties .....................................................................................................55
10.
Summary of papers and contribution to the field............................................................59
10.1.
Paper 1...............................................................................................................................59
10.2.
Paper 2...............................................................................................................................60
10.3.
Paper 3...............................................................................................................................60
10.4.
Paper 4...............................................................................................................................61
10.5.
Paper 5...............................................................................................................................61
11.
Future work...........................................................................................................................63
11.1.
In-situ decomposition studies..........................................................................................63
11.2.
Wear behavior ..................................................................................................................63
11.3.
Mechanical properties .....................................................................................................64
11.4.
Surface directed spinodal decomposition.....................................................................64
11.5.
Improved thermal stability by alloying .........................................................................65
12.
Bibliography ..........................................................................................................................67
Paper 1 .................................................................................................................................................77
Paper 2 .................................................................................................................................................83
Paper 3 .................................................................................................................................................93
Paper 4 ...............................................................................................................................................101
Paper 5 ...............................................................................................................................................123
1. Introduction
Contrary to popular belief, thin films are all around us in our daily life. It can be in the
appearance of a decorative and wear resistance coating on a phone or a wristwatch, or as an
electrical contact inside an electrical gadget. New applications using thin film technology are
constantly evolving and the world market is growing. The important market for this thesis is
the protective coatings used in the cutting tool industry. These coatings are expected to have
high hardness and stiffness, in combination with good chemical inertness. One of the
coatings fulfilling these requirements is Ti1-xAlxN, the material explored in this thesis.
1.1. Research goals
The main objective of this thesis is to understand the behavior of Ti1-xAlxN/TiN multilayer
coatings, or more explicitly, the influence from the lamellar structure on the mechanical
properties, thermal stability and cutting performance. The unstable c-Ti1-xAlxN transforms to
nano-sized domains rich in AlN and TiN by spinodal decomposition, which results in
improved mechanical properties. I herein study the details of this isostructural
decomposition which are not fully understood. I also examine the possibilities to control the
decomposition behavior with a multilayer architecture, and through this improve the
mechanical properties. Hence, with this work some light might be shed on how internal
interfaces influence the high temperature behavior of Ti1-xAlxN. The specimens were
deposited with the physical vapor deposition (PVD) technique reactive cathodic arc
evaporation, using a full-scale industrial system. Characterization was performed by analytical
transmission electron microscopy, x-ray diffractometry and atom probe tomography to
obtain information of the microstructure and the composition. Furthermore,
nanoindentation and cutting tests were performed to investigate the mechanical properties
and differential scanning calorimetry was used to examine the thermal stability.
1
Introduction
1.2. Outline of the thesis
The second chapter gives an introduction to the materials used for hard coatings, especially
the ones of interest for this work. The structure of multilayers and the resulting hardening
mechanisms are described in chapter 3. Chapters 4 and 5 deal with phase transformations,
decomposition behavior and how the coatings in this work were deposited. This is followed
by a description of the characterization techniques used in this thesis. Chapter 7 gives a short
introduction to phase-field simulations and how the method has supported the experimental
results in this work. Chapter 8 explains the wear mechanisms studied in this work and
chapter 9 presents data, not found in the appended papers, showing epitaxial stabilization
effects and mechanical properties of an arc evaporated c-Ti0.25Al0.75N/TiN multilayer.
Chapter 10 gives a summary of the appended papers and their contribution to the field. The
final chapter contains some suggestions of future work, based on the results in this thesis.
2
2. Hard coatings
To give a perspective of how hard the coatings in this work are, an overview with their
hardnesses in comparison to steel, c-BN and diamond is given in Figure 1. It is seen that
steel, a common engineering material, and cemented carbide (WC), a typical cutting tool
material, is softer compared to the hard coatings. At the right end of the graph the hardest
material known is found, diamond. The second hardest material in the graph is c-BN, a
modern cutting tool material.
Figure 1. Approximate hardness of stainless steel [1],
AlN[2], WC, TiN, TiAlN, TiAlN/TiN [paper 2], c-BN [3]
and diamond [4].
Hard ceramic coatings are found in a broad range of markets such as in aerospace, automotive, medical technology, optics and electronics. They were introduced in the cutting tool
industry in the 1970´s and today about 90% of the inserts for metal cutting are coated. The
reason for this is simply due to the increased performance and lifetime of a coated tool in
3
Hard coatings
comparison to an uncoated tool. The hard coatings have evolved from the chemical vapor
deposition (CVD) TiC coatings to today’s more complex quaternary and multilayer coatings.
An overview of when important hard coating has been introduced to the market is seen in
Figure 2.
Figure 2. Year of market introductions of coatings for cutting
tools, based on Ref. [5] For details on the two deposition
techniques CVD and PVD see chapter 5.
In the manufacturing industry increased productivity is always desired. In terms of cutting
parameters, this means higher cutting speeds and feeding rates during operation. The effect
of such demands is requirements of improved thermal stability, mechanical properties and
oxidation resistance of the protective coatings. The development has often been driven by
the desire to control the microstructure and composition in such way that the properties are
improved or tailored for a specific cutting operation. The desired microstructure can evolve
during deposition or at elevated temperature, often referred to as self-organization. The selforganization has been achieved by selecting a system with a miscibility gap, where the atoms
are forced into a supersaturated unstable solid solution. This results in a microstructural
transformation upon exposure to elevated temperatures or during the cutting operation.
Much of the focus has been on the unstable Ti1-xAlxN system, which has the ability for age
hardening through decomposition at elevated temperatures [6-8]. The materials relevant for
the hard coatings investigated in this work are described below.
4
Hard coatings
2.1. Titanium Nitride
TiN is a hard ceramic material with a NaCl crystal structure, as illustrated in Figure 3. The
lattice parameter of TiN has been measured to a=4.24 [9] and the bonding structure is
reported to be a mixture of covalent, metallic and ionic bonds [10]. The covalent bonding is
the explanation for the high hardness of ~20 GPa, measured on single crystals [11]. When
TiN is deposited with arc evaporation, the technique used in this work, the hardness is
measured to ~26-30 GPa due to lattice defects induced by the deposition conditions [12, 13].
As seen in Figure 2, TiN was one of the first coating materials used in the cutting tools
industry and it is still used as diffusion barriers and for decorative coatings. The material can
be deposited as hard or protective coatings utilizing both physical vapor deposition (PVD)
and CVD. It has a shiny golden appearance and, like most other ceramic materials, relatively
good mechanical and thermal properties. TiN has been shown to oxidize at a rather high rate
above 450 °C, which is one of its main disadvantages when used as a tool coating. Annealing
of an arc evaporated TiN film in an inert atmosphere, results in a decrease of the hardness
towards its intrinsic hardness due to defect annihilation and stress relaxation [13, 14].
Figure 3. The NaCl-structure. Bright spheres
correspond to N and dark to Ti or Al.
2.2. Titanium Aluminium Nitride
If the Ti in the TiN matrix, Figure 3, is partially replaced randomly by x percent of Al it will
result in a cubic Ti1-xAlxN. The material is used in a wide range of applications such as
protective and wear resistance coatings [8, 15-17], diffusion barriers [18, 19] and optics in
solar devices [20].
In this work mainly three compositions of Ti1-xAlxN have been investigated, x= 0.50, 0.67
and 0.75. I has been shown that only a few percent of AlN can be dissolved in the cubic TiN
[21, 22] at equilibrium conditions. However, it is possible to deposit c- Ti1-xAlxN with as high
Al content as ~67% by reactive cathodic arc evaporation [7] (see paragraph 5.1). The ability
of the technique to incorporate such high Al content has been attributed to the combination
5
Hard coatings
of low deposition temperatures and the highly ionized plasma. A higher Al content than
~67% will result in growth of a mixture of hexagonal and cubic phases, see chapter 9 or
Refs. [7, 23, 24] for more details on this. When the Al content is increased, the lattice
parameter of the ternary will decrease and approaches the one of pure AlN [25, 26], as seen
in Figure 4. It should be noted that both the calculated and experimentally measured lattice
parameters are deviating from the linear Vegard's law. The color of the coating will change
from the TiN-golden to a dark blue/grey with higher Al content. The color change has been
attributed to the change in valence electron band structure [24].
The as-deposited c-Ti1-xAlxN is unstable and decomposes into the binary phases in two
steps upon heat treatments, first via spinodal decomposition to domains rich of c-TiN and
metastable c-AlN. Further annealing results in a transformation of the c-AlN to its
equilibrium phase h-AlN [6, 7, 13, 26-28]. The decomposition pathway can be summarized
as
c-TiAlN → c-TiN + c-AlN → c-TiN + h-AlN
The first step is believed to be a spinodal decomposition because of the miscibility gap and
that ab initio calculations show a negative second derivative of Gibbs’ free energy [22, 25, 29],
which is typical for systems phase separating with this mechanism. Also experimental results
show features typical for the spinodal decomposition, such as coherent domains [7, 30, 31] a
widespread decomposition [26, 27] and a constant domain size over a period of time during
decomposition [paper 4]. The hardness of arc evaporated c-TiAlN has been measured to
~32 GPa depending on the deposition conditions and composition [6, 8, 27, 32]. Age
hardening after thermal annealing of the coating is seen and associated to the decomposition
of the coating [6, 8, 13, 27].
During the last decade several investigations have been performed on TiAlN alloyed with a
third metal, such as e.g. Cr [33-35], Ta [36], Hf [37], or Zr [38, 39]. The motivation for this is
the potential improvements of the thermal stability, oxidation resistance and mechanical
properties through alloying.
6
Hard coatings
Figure 4. Lattice parameter of c-Ti1-xAlxN, determined experimentally (circles) from
powder x-ray diffraction and theoretically from ab-initio calculations (squares)[22, 40].
The value for TiN (open circle) is from Ref. [41]. Reprinted from Ref. [26] with
permission.
2.3. Aluminium Nitride
In stable state, AlN is in a hexagonal wurtzite structure, here denoted as h-AlN, illustrated
schematically in Figure 5. The material has been shown to have relatively high thermal
conductivity, and is used both as an electrical insulator and semiconductor. Important for
this work, the metastable NaCl c-AlN with a lattice parameter of a≈4.05 Å, is found both as
a decomposition product of c-TiAlN and at high pressure and high temperature (HPHT)
conditions [42]. c-AlN can be grown by PVD in a cubic state utilizing multilayer epitaxial
stabilization effect [43]. Its relatively low lattice mismatch to TiN is of great importance for
the age hardening of TiAlN [6, 7, 26, 27, 32]. AlN has also been observed experimentally in
the metastable zinc-blende phase [44].
Figure 5. Hexagonal structure of AlN where
dark spheres correspond to Al and bright to N.
7
3. Multilayer coatings
In many cases the specification profile for modern coatings are complex and can only be met
by advanced alterations of the material. One technique which has been utilized successful to
do this is the design and growth of multilayer structures. The complex multilayers are
justified by their improved properties compared with monolithic systems. It has been shown
that e.g. electrical [45, 46], optical [47], tribological [48], and oxidation resistance [49]
properties can be influenced by the structure. The sections below elucidate multilayer
coatings, a subclass of thin films, with attractive features relevant for this work.
Figure 6. A schematic of a multilayer structure.
3.1. Structure
A multilayer structure is grown when different materials (A and B) are deposited alternatively
and repeatedly as illustrated in Figure 6. The thickness of two consecutive layers in this work
referred to as the period (Λ) of the multilayer. If a layer consists of a single plane of atoms it
is referred to as a monolayer. The multilayer structures are, interestingly, not only found in
materials made by man but also in nature. One example of this is the Cicindela scutellaris
beetle, seen in Figure 7. It is believed that the beetle uses the multilayer structure both as
9
Multilayer coatings
mechanical protection and to attain an unique and attractive color which increases the
possibilities for mating [50].
A superlattice is a special case of the multilayer where the film is grown as a single crystal,
i.e. no grain boundaries and coherent interfaces throughout. If certain conditions are fulfilled
in the superlattice coatings, such as similar chemical bonding and similar atomic radii of
constituents, entirely new materials with properties and characteristics not directly related to
the layer materials can be attained [51, 52]. The superlattices are considered as a separate
class of thin films because of the possibility that that they will exhibit unique properties.
Epitaxial layers also allow for growth of unstable phases, not found in the phase diagram, as
for example c-AlN [43]. This mechanism is called the epitaxial stabilization effect and has
been applied to several material systems [53-55] and it is used also in this work (see chapter
9).
Figure 7. (a) TEM cross sectional view of the (b) Cicindela
scutellaris beetle, reprinted with permission [50].
3.2. Mechanical properties
In 1970 J. S. Koehler wrote “We would like to propose a composite material which is rather different
from previous suggestions. Suppose that a specimen is prepared by epitaxial crystal growth which consists of
alternate layers of crystals A and B.” in his theoretical work “Attempt to make a strong solid”
[56]. He also stated some suggestion for the choice of materials for A and B, such as that the
lattice parameter should be nearly equal, the elastic constant should differ and the thickness
of the layers should be thin, i.e. in the order of 100 atom layers. In 1978 Lehoczky grew a
multilayer coating of Al-Cu laminates, based on Koehlers design, which confirmed his
10
Multilayer coatings
theories [57, 58]. Enhancements of the hardness in layered structures have since then been
observed in a wide range of multilayer classes such as metal-ceramic [59-61], metal-metal
[62-64], and ceramic-ceramic [11, 30, 40]. The physics behind the hardness alteration is based
on hindering of dislocation glide across the interfaces due to the difference in E-modulus
and the lattice mismatch of the layers.
One of the first theories adapted to explain the multilayer hardening was the Hall-Petch
relationship for grain size. This theory was published independently around 1950 by Hall [65]
and Petch [66] who essentially established the same thing, namely a relationship between
yield strength(σy) and grain size (d) as
σ y =σ0 +
K
d
Eq. 1
where K and σ 0 are material dependent constant. The hardness increase is connected to
grain size reduction, which results in increased grain boundary areas that hinder the
dislocation motions in the form of locked up Frank-Read sources and dislocation pile-ups.
The Hall-Petch relationship has been shown to be valid down to very fine grains of only
few nm [67]. To adopt this theory to multilayers a layer in the stack is considered as a grain,
i.e. the Λ/2 of the multilayer as d. However, the Hall-Petch relationship does not give an
exact estimation of what hardness to expect, but rather the relative increase from a decreased
layer period. In addition, nano-scale multilayers have shown deviations from the relationship
in Eq. 1, see e.g. Refs. [68-70]. There are also studies showing that an inverse Hall-Petch
relationship exists beyond a critical multilayer period, especially for nitiride multilayers. This
hardness decrease was shown early by Helmerson et al. [71] in a multilayer consisting of TiN
and VN layers. The decrease in hardness is not fully understood but has been associated to
incomplete layers and intermixing resulting in a broken imperfect multilayer stack. Limited
intermixing by using immiscible layers has been shown to lower this hardness decrease [72].
Coherency stress hardening is a recognized hardening effect seen in the multilayer structures [73,
74] and is believed to be present also in the coatings investigated in paper 2. This effect will
be active if an epitaxial multilayer is grown and the constituents have dissimilar lattice
parameters i.e. the lattices has to be distorted to be coherent across the interface. The
hardness increase, originating from the coherency, is explained by the resulting stress field
which restricts the dislocation movements. This theory is related to the work by Cahn
investigating the effect of internal stresses produced by coherent phases on dislocation
movements [75].
11
Multilayer coatings
It has been suggested that an effect similar to Orowan strengthening can be an active
hardening mechanism in nanometer scale lamellar structures. The theory was originally
developed for understanding the hardening resulting from precipitation, where a dislocation
is stopped and have to “loop” around the obstacle leaving a so called Orowan loop [76]. The
Orowan-like strengthening in multilayers instead suggests that the hardness increase is an
effect of plastic deformation, occurring by dislocation motion and bowing inside the layer.
The presence of an Orowan-type mechanism present in multilayer structures has been
confirmed in Refs. [70, 77-79].
As seen above, there are numerous effects explaining the alteration of the mechanical
properties of lamellar structures. However, it is unlikely, that the improvements seen in the
copious number of publications on multilayers is attributed to one particular type of
hardening mechanism, but instead are due to a combination of the above explained theories.
It should also be noted, that even if the stated specifications are fulfilled it is not sure that a
multilayer show improved mechanical properties [80, 81].
There are also hardening effects active in monolithic coatings such as e.g. strain hardening
(also referred to as work hardening) which one could expect also in the multilayer coating in
this work. Here, the hardening increase is essentially an effect from a dramatic increase of
the number of dislocations-dislocation interactions and the resulting reduced dislocation
mobility. The creation of defects during deformation is similar to the large amounts of
defects which are introduced during coating growth with arc evaporation. A hardness
increase with an increased defect density was for example seen for the arc evaporated
TiCxN1−x coating in Ref. [82]. The dislocation mechanism behind the hardening is reported
to be dislocation pile-ups and production of sessile dislocations [83].
3.3. Multilayer coatings on cutting tool
The cutting tool companies showed an interest for the multilayer coatings already in
beginning of 1980s. The multilayers were then believed to have potential to adapt to, or
compensate for, mechanical stresses and thermal loads at the high tool temperatures during
metal machining. It had also been shown that the laminated structure could reduce the
diffusion processes, which in some cases are detrimental to the tool lifetime [84]. The first
use of a multilayer in the cutting tool industry was reported by Hara et al. [85]. The multilayer
were deposited on cemented carbide and consisted of a titanium based (TiC or TiN) seedlayer and had intermediate layers with a non given composition. It was reported that the
coating had a better performance for high-speed metal cutting operations compared to single
or double layered coatings. 1989 Sandvik Coromant introduced their first multilayer,
designed with layers of TiC, TiN and Al2O3 [86]. About one year later Kennametal UK
reported that they also had designed and deposited a multilayer coating consisting of TiCN,
12
Multilayer coatings
Al2O3 and TiN-layers [87]. Today most of the major cutting tool companies, like Seco Tools
AB, Kennametal, and Sandvik have some sort of multilayer in their product line, even if the
used materials are rarely stated. A multilayer coating does not only allow for tuning of the
mechanical properties, but in addition gives an impression of being an advanced “hi-tech”
coating, which is a good selling argument. The CVD multilayer coatings found on the
market today are commonly arrangements of between 3 and 13 layers, while for PVD
coatings multilayer stacks consisting of more than 4000 individual layers are reported.
13
4. Phase transformation and decomposition
Phase transformations have been used for materials engineering by mankind for more than
3000 years. One example is the hardening of steel swords and armor, where the phase
transformations were achieved by heat and/or mechanical treatments [88]. It is, however,
only during the past centuries that we have obtained the understanding of the underlying
mechanisms, and that it is indeed a phase transformation occurring in the material during the
treatments. This insight is a result of the extensive research, primarily performed on metals.
The work has resulted in thousands of publications, both theoretical and experimental,
dealing with phase transformations. The theories have over the last 20 years been adapted to
the new ceramic materials, such as the hard coatings investigated within this work. This
chapter deals with the two most common mechanisms seen in the ternary nitrides in the
field of ceramic hard coatings, nucleation and growth and spinodal decomposition. The phase
transformations may change the material properties in both positive and negative ways. For
example, heat treatment of solid solution c-TiAlN results in a positive evolution of the
mechanical properties during spinodal decomposition and a negative evolution of the same
properties during the nucleation and growth transformation to h-AlN.
4.1. Nucleation and growth
All phase transformations are driven by minimization of the total energy of the system. The
generation of a new phase with a lower free energy than the matrix is one mechanism to do
this. The nucleation of the new phase occurs via the formation of a small embryo inside the
original matrix. The nucleation can be either homogenous or heterogeneous. In the case of
heterogeneous nucleation, the nucleus is formed at a defect such as a grain boundary, an
elemental inhomogenity or a particle. Logically, some energy is required for the generation of
the nucleus which means a passage of an energy barrier. The original matrix is thus said to be
in a metastable equilibrium. The free energy per mole of the newly generated phase is of
coarse lower than in the matrix. However, if taking into account the surface energy, the
15
Phase transformations and decomposition
nucleus can have a larger total energy per molecule than the matrix. To decrease the total
energy the nucleus has to start growing and enter the growth stage. In the growth stage,
atoms are flowing towards the nucleus with downhill diffusion. A critical radius, where the
growth is energetically favorable, has to be reached. Therefore, there is possible that the new
nucleus instead of growing is dissolved in the matrix. In the most common case,
heterogeneous nucleation, the surface energy is a less dominating factor. Then the nucleus is
formed in shapes with less surface energy, such as e.g. hemispheres. From a microstructural
point of view, nucleation and growth is usually seen to result in relatively few large
precipitates with sharp interfaces.
It has been found, that the nucleation and growth mechanism can be suppressed in thin
layers or films by preventing the nuclei to reach the critical size [89]. This will result in a
higher temperature needed for the phase transformation to occur, and has been reported e.g.
for HfSiON films [90]. A similar effect is seen for the transformation of the c-AlN to h-AlN
in the multilayer coatings in paper 3.
4.2. Spinodal decomposition
Spinodal decomposition was observed experimentally first in 1940 in a Cu-Ni-Fe alloy [91],
which showed signs of a periodic elemental fluctuations in a initially homogenous alloy. The
mechanism behind the microstructural change was however not understood until M. Hillert
gave a theoretical explanation of the decomposition behavior in 1955. Around 1960, J.W.
Cahn improved the theory in two highly cited articles [92, 93]. The presence of spinodal
decomposition in the Ti1-xAlxN system has been confirmed both experimentally and
theoretically [6, 7, 22, 25, 31]. This paragraph gives a brief theoretical explanation of
decomposition type based on the Ref. [94].
Material systems which are immiscible exist, i.e. it is unfavorable for their constituents (A
and B) to mix. Figure 8 shows a phase diagram and free energy curve of such binary alloy
with a miscibility gap. If the single α-phase alloy with composition XAB at temperature T1 is
quenched into the two phase region (α1+α2) to temperature T2 the system will be unstable
and in a local maximum, with the free energy G1. Small elemental fluctuations or defects will
result in a decrease of the total free energy to G2A and G2B , and A- and B-rich regions will
consequently be formed. Hence, there is no energy barrier associated to the spinodal
decomposition, and it is often seen to occur spontaneously over large volumes. The
decomposition will, in contrast to nucleation and growth, occur through up-hill diffusion, in
which the atoms diffuse toward regions which are already enriched with the diffusing atom.
The diffusion process is characterized by a negative second derivative of the free energy.
Outside the spinodal the phase transformation can proceed only through the more
16
Phase transformations and decomposition
“conventional” nucleation and growth, as explained in previous section, by down hill
diffusion.
Temperature
T1
α
spinodal
T2
α1+α2
xAB
Gibbs Free Energy
xA
xB
G1
∂ G
≤0
∂x 2
2
G2B
G2A
Composition
A
B
Figure 8. Schematic phase diagram of a binary
alloy and corresponding free energy curve.
A typical composition profile of the spinodal decomposition can be seen in Figure 9. The
profile is, compared to the nucleation and growth, more subtle but over a larger volume.
From the illustration it is can be seen that the composition wavelength is constant during the
decomposition. An experimental example of this is presented in Figure 10, where the
wavelength is constant during the first 20 min of the isostructural decomposition of
Ti0.50Al0.50N. The wavelength of the modulation at this stage can, accordingly the theory of
Chan [92], be calculated by
λ = 4π
−κ
δ G (Gm , κ )
2
Eq. 2
δxk 2
17
Phase transformations and decomposition
where Gm is the molar free energy of mixing, xk the molar fraction of element k and κ the
gradient energy coefficient. λ for Ti0.50Al0.50N can be calculated by Eq. 2 using
thermodynamic data from Alling et al. [22] to ~2.4 nm for temperatures around 900 °C. This
value is reasonable considering the experimentally measured domain sizes for short
annealing times in paper 4 and 5.
(b)
Composition
Increasing time
(a)
Distance
Distance
Figure 9. Compositional profiles during (a) nucleation and growth and (b)
spinodal decomposition with increasing time downwards.
From a microstructural point of view the spinodal decomposition occurs over large volumes
but the domains are typically smaller than the ones resulting from nucleation and growth. A
good example of the widespread decomposition is seen in Figure 5 in paper 4, where a
periodic elemental modulation is observed in all visible columns. Furthermore, because of
the diffusion type involved in spinodal decomposition, the boundaries between the resulting
phases are usually coherent and have relatively diffuse interfaces, compared the sharp ones
resulting from nucleation and growth. Diffuse interfaces of the spinodally decomposed
TiAlN has been revealed in transmission electron microscopy [31, 95] and atom probe
tomography [27, 96, 97].
18
Phase transformations and decomposition
Figure 10. Compositional wavelength evolution of
Ti0.50Al0.50N during isothermal annealing.
4.2.1.
Coarsening
The spinodal decomposition is followed by a latter stage, coarsening. This results in a lower
number of larger domains when the smaller domains coarsen, i.e. a increase of the
compositional wavelength [93]. This latter stage of the decomposition is driven by the
minimization of the evolving surface and gradient energies. It has been debated if the
spinodal decomposition and coarsening are two separate steps or are overlapping [98]. What
is clear is that it is challenging, from an experimental point of view, to separate the two
mechanisms. From an energetically point of view, it is most likely that there always will be an
overlap of the mechanisms due to presence of inhomogeneities and defects, allowing for an
uneven decomposition process. In paper 5 we observe two stages both in the simulations
and the experimental wavelength evolution of Ti0.50Al0.50N, i.e. spinodal decomposition and
coarsening, as seen in Figure 10. Figure 11 gives the compositional wavelength after
coarsening of Ti0.34Al0.66N, for long annealing times (>1 h) at 800, 850 and 900 °C, extracted
with an autocorrelation function from STEM images. The graph shows that the coarsening
rate is highly dependent on the annealing temperature. The results give a rough estimation of
what compositional wavelength to expect in a Ti0.34Al0.66N coatings subjected to long
annealing times.
19
Phase transformations and decomposition
Figure 11. Compositional wavelengths of annealed
Ti0.34Al0.66N, extracted form STEM images. The
measurements were performed at 5 different locations in
the coating. The value at t=0 corresponds to the calculated
initial wavelength for the composition.
4.2.2.
Surface directed spinodal decomposition
Simulations show that the kinetics of the spinodal decomposition and the resulting evolving
microstructure can be significantly affected by the presence of an interface or a surface [99102]. The feature affecting the decomposition behavior can e.g. be the surface between the
substrate and a thin film [103], a grain boundary [104] or the interfaces in a multilayer stack.
The characteristics of such interface-influenced decomposition are formation of a layered
structure parallel to the interface, i.e. a dominant wave vector directed normal to the surface.
Such decomposition is exemplified by the phase-field simulations in Figure 12. This
phenomenon is known as surface directed spinodal decomposition (SDSD), and has been
experimentally observed several times in polymers [105-107] but also in metals [108]. The
reason for the modified decomposition behavior has been attributed to e.g. coherency
stresses and wetting mechanisms [109, 110]. However, it has been shown that the layered
structure can arise even without influence from surface interaction energies such as wetting
[101].
20
Phase transformations and decomposition
Figure 12. Phase-field simulation of surface directed
spinodal decomposition in of a Ti0.34Al0.66N layer
constrained by TiN. The simulation was performed by Dr.
K. Asp Grönhagen.
For coatings, the publications showing experimental confirmations of this behavior are very
limited. Adibi et al. [103] showed compositionally modulated platelets during growth of
Ti0.50Al0.50N, but reports on a similar behavior during post annealing is lacking in the
literature. In paper 5 the spinodal decomposition after short time annealing in Ti1-xAlXN
enclosed by TiN layers is investigated. The existence of the typical SDSD structure, seen in
Figure 12, is highly dependent on the initial elemental fluctuations. For relatively high initial
fluctuations, as expected in the arc evaporated coatings in this work, the microstructure
typical for the decomposition type will be almost completely dissolved by the isotropic
spinodal decomposition in the “bulk” Ti1-xAlxN. A similar behavior, i.e. a decay of the
layered structure, has been seen in simulations when increasing the thermal noise [105].
Furthermore, it is shown by the simulations that the SDSD is the main reason for the earlier
onset of the decomposition of a multilayer coating in comparison to a monolithic coating,
seen for example by DSC.
21
5. Coating deposition
A wide variety of coating deposition techniques exists but two main classes can be
distinguished, the physical vapor deposition (PVD) and chemical vapor deposition (CVD).
CVD utilizes a gas mixture as the source material, which is heated to chemically react and
form the coating on the substrate. PVD synthesizes thin solid films by evaporating the
coating material in vacuum from a solid material, the target (here cathode). Eventually, the
vaporized material will condense on the substrate and form the solid film. There exist a
number of different PVD techniques such as e.g. sputter deposition, pulsed laser deposition
and cathodic arc evaporation. The deposition technique used to synthesis the coatings for
this thesis is reactive cathodic arc evaporation, which is explained briefly below.
5.1. Cathodic arc evaporation
Cathodic arc evaporation is the most commonly used PVD technique in the cutting tool
industry. The main reason for this is the efficient source of highly ionized material that
produces a dense well adherent coating from a wide range of metals. Furthermore, cathodic
arc evaporation has a relatively high deposition rate compared to many other PVD
techniques which, of course, is important for the industry. The technique is usually described
as a low-voltage, high current plasma discharge between metallic electrodes in a controlled
gas or a vacuum [111]. The material to be deposited, the cathode, is evaporated by an arc
discharge which is ignited at the surface using a mechanical trigger which initiates a voltage
breakdown. This creates non-stationary spots of very high current densities, high
temperatures and low voltage discharges which locally melt the material. These locally
molten areas are called the cathode spots and are usually smaller than 10 µm [112]. The
cathode spot is present only for a short period of time before a new one is ignited, which
results in the characteristic stroboscopic appearance of the arcs moving over the cathode
surface. The electrical current of the arc discharge is transported in the plasma produced by
the discharge itself, and the technique is thus said to be self-sustaining. The molten, and later
23
Coating deposition
evaporated, material from the pool transforms into ions and electrons, i.e. the plasma. The
plasma is highly ionized close to the cathode spot, where an almost 100% ionization has
been measured [113]. The ions will finally migrate to the substrate where they condense and
form the solid coating. The plasma can be controlled and manipulated using magnetic and
electric fields, which allows for some control of the growing film. If there is reactive gas in
the deposition system, the ions will react with it at the substrate, as in the case of N2 during
deposition of e.g. TiN.
The targets used in this work were either elemental (Ti) or compound (Ti1-xAlx) cathodes.
By varying the cathode composition the resulting composition of the coating was altered. It
is not sure that the cathode composition and the resulting film composition will be identical.
Rogström et al. [26] recently measured the composition of Ti1-xAlxN by energy dispersive xray spectroscopy (EDS) to x=0.65 and x=0.47, for the x = 0.67 and x = 0.50 cathodes,
respectively. This has been explained by a different mean charge state of the plasma
depending on the atomic number [113]. Hence, a negatively biased substrate will influence
the plasma differently depending on the cathode composition. The existent data of the
charge state have, however, been measured for elemental cathodes and might vary for the
compound cathodes used in this work [114]. Also the re-sputtering is varied depending on
the element, which in addition might influence the final coating composition [112].
5.1.1.
Film growth
The growth of a polycrystalline coating is basically a phase transformation, the vaporized
atoms from the cathode and the reactive gas condensate on the substrate and forms a solid.
The atoms will move over the substrate and form small nuclei or islands which will grow and
form grains and act as the building blocks for the final coating. The phase transformation
during deposition is therefore a form of nucleation and growth. The final structure of the
coating will depend on the mobility and energy of the species arriving at the surface of the
substrate or the growing coating. One easy way to control the energy is by varying the
temperature. If the temperature is increased the diffusivity and mobility will increase. Hence
the atoms will be able to travel longer distances, which results in larger grains. In the case of
a lower temperature the atoms will be trapped in low energy lattice positions, which will
result in many small grains, and sometimes a porous structure [115, 116]. Arc evaporation
gives the possibility to vary the energy by applying a negative bias to the substrate, which
attracts the ions in the plasma. A higher bias will result in higher energy of the arriving
species and an increase of the penetration of ions into the growing structure. Hence, a higher
bias often results in larger compressive residual stresses and more defects [117]. An example
of this is seen in Figure 13, showing thermal responses of Ti0.50Al0.50N deposited with -20, 40 and -60 V substrate bias. The presented peaks correspond to stress relaxation and defect
24
Coating deposition
annihilation and is increasing in magnitude with a higher bias. The mechanism has an onset
around the substrate temperature during deposition, 500-600 °C. The graphs indicates that
the defect density increase when the bias is increased. The stress relaxation of TiAlN during
annealing is described in Ref. [26]. A higher defect density can also result in re-nucleation of
grains which reduced the total grain size [118].
Figure 13. Thermograms of stress relaxation in
Ti0.50Al0.50N deposited at different bias.
5.1.2.
Macroparticles
A side effect from the arc spot plasma generation is the ejection of macroparticles, also
called droplets. The macroparticles arise from the molten pool of cathode material and are
sprayed in a normal direction to the cathode as liquid droplets [112, 119]. The particles will
be incorporated in the growing film where they work as nucleation sites for new grains.
When growing a multilayer structure a macroparticle can result in a severe breakage of the
continuous multilayer stack, as seen in Figure 14. The image illustrates how the droplet
works as a nucleation site for a new column. The macroparticles will be present also at the
coating surface, which results in an increase of the surface roughness. The presence of
surface- and internal-macroparticles can be detrimental to the properties and the
performance of the coating. The generation of macroparticles and their properties is closely
related to cathode material and deposition condition. The amount of particles can be
reduced by e.g. low cathode current densities, cooling of the cathodes or using filtering. The
filtering is traditionally utilizing that the macroparticles have no net charge in contrast to the
plasma [120]. Hence, when the plasma flow is controlled and concentrated by a curved
25
Coating deposition
magnetic plasma duct, the inclusions are significantly reduced. A filter is, however, rarely
used in the cutting tool industry, due to the resulting decrease in efficiency.
Figure 14. Cross sectional STEM image of a macroparticle in a TiAlN/TiN
multilayer coating with TiAlN being the dark layer. The original size of the
shown macroparticle has most likely been decreased during sample polishing.
5.2. Multilayer growth
As mentioned in the introduction of this chapter there exist a vast number of deposition
techniques and all of them can usually be modified for multilayer deposition. The technique
of choice depends on the intended application for multilayer. In the case of for example
multilayer x-ray mirrors, where a high interface quality and low defect densities are wanted,
low energy sputtering are used. For the cutting tool industry where a high deposition rate is
the priority, cathodic arc evaporation is preferred. The deposition of a multilayer is
traditionally achieved by using two sources from which the alternate deposition is controlled
by either turning them on and off, or by using some kind of mechanism periodically shading
the sources. The shading can be performed by using a rotating substrate holder, often
referred to as planetary rotational system, or by shutters. The shutters found in the
deposition systems are usually computer controlled and has an opening/closing time of <0.1
s, allowing for a very controlled deposition of the individual layers. By controlling the
reactive gas during the deposition, it is possible to grow a multilayer stack consisting of e.g.
nitride-metal layers.
26
Coating deposition
In this work a full scale industrial cathodic arc evaporation system was used for the
multilayer growth. The alternating growth was performed using a rotating drum working
both as substrate holder and shading device, as seen in Figure 15. Cathodes of Ti and TixAl1-x
were located on opposite sides of the rotating sample fixture in such way that the rotation
speed was determining the individual layer thickness. As example, the use of 1, 2 and 4
revolutions per minute (rpm) in paper 2 resulted in three different Λ, of 25+50, 12+25 and
6+12 nm, respectively. By varying the displacement in height of the substrates relative to the
two different cathodes, resulted in passages of the substrates through regions of different
plasma flow. This allowed us to grow symmetric and asymmetric multilayers. The drawback
of this system is that it is harder to control the exact multilayer period but with the advantage
of high deposition rate.
Figure 15. Schematics of the deposition system used in this thesis.
Black squares indicate the substrates and the arrow the rotation of
the drum. The drum has a radius of ~0.25 m.
27
6. Characterization
This chapter describes the techniques which have been used to characterize the coatings in
this thesis. In the pursuit of a complete story and a proper explanation, a combination of
techniques typically has to be used. The information extracted from the coatings in this work
is thermal stability, mechanical properties, decomposition behavior, microstructure and
composition.
6.1. Nanoindentation
One of the most important physical properties of a protective coating is the hardness.
However, due to the small dimension of the coating and the resulting likelihood of influence
from the substrate during testing, ordinary hardness test methods, e.g. Vickers hardness test,
can not be used. Instead a nanoindentor, which only penetrates the coating, is the technique
of choice. In nanoindentation, the depth of penetration of a diamond indenter is measured
together with the prescribed loading curve. In this work the maximum load was in the range
of 5 - 25 mN. During an indent the load and displacement of the tip is logged. The resulting
load displacement response which typically shows an elastic-plastic loading is followed by
elastic unloading, as seen in Figure 16. The elastic equation of contact is then used in
conjunction with the unloading data to determine hardness of the specimen material.
29
Characterization
Figure 16. Load displacement curve obtained by nanoindentation. This
particular curve is a result from indentation on a monolithic Ti0.34Al0.66N
thin film.
All hardness values in this thesis have been calculated by the method by Oliver and Pharr
[121]. The method is based on determination of the area of contact between the tip and the
sample at the maximum load P. From the ratio between P and A, it is then possible to
determine the hardness as
H=
P
A
Eq. 3
Figure 17 shows a schematic cross section through an indentation with a Berkovich tip. The
contact depth hc, is given by
hc = hmax − ε
Pmax
dP dh
Eq. 4
where hmax is the total displacement of the surface at maximum load, Pmax the maximum load
and ε a geometric factor that depends on the tip shape, in this work ε = 0.75. The dP/dh
corresponds to the contact stiffness and is the gradient of the upper part of the unloading
curve as seen in Figure 16. Once the hc is known, the area of contact A can be calculated
with
A = 24.49hc2 .
30
Eq. 5
Characterization
This calculation is possible due to the known pyramidal shape of the used Berkovich
indenter. Now, knowing the applied force and the contact area, the hardness can be
calculated with Eq. 3. The arc evaporated coatings investigated in this work have high surface
roughness due to the droplets, which will affect the hardness measurements. Therefore the
coatings were polished before indentation with 1 µm diamond abrasive as the final step. The
indents, usually 20 – 30, were manually positioned on droplet free positions.
Figure 17. Schematic cross section of a Berkovich indent.
6.2. Thermal analysis and calorimetry
Thermal analysis (TA) is the analysis of the change in sample properties related to a forced
alteration of the temperature. In this work sample properties refer to microstructure,
chemical composition and thermodynamic properties. An alteration in temperature refers to
a predetermined sequence of temperatures with respect to time.
Here all thermal analysis was performed using differential scanning calorimetry (DSC) in
combination with thermo gravimetry (TG) and masspectrometry. From this setup it was
possible to investigate changes in mass, decomposition temperatures and the chemical
composition of evolving gases.
Figure 18. Schematic of a DSC setup, S
indicating the sample and R the reference.
31
Characterization
6.2.1.
Differential scanning calorimetry and thermogravimetry
A schematic of the used DSC setup can be seen in Figure 18. In this technique the change of
the difference in heat flow rate to the sample and the reference, i.e. TR and TS, is analyzed
while they are subjected to a temperature program, T(t). The relative changes in heat flow
are used to investigate at what specific temperatures thermo mechanical phenomenon occurs.
In a coating, changes in heat flow are attributed to phase transformations, release of a light
element or stress relaxations. Positive heat flows are assigned to exothermic effects, and the
corresponding peaks here point in the positive direction in the thermograms. To avoid
disturbance from oxidation on the phase transformations, the DSC measurements are
performed in a protective atmosphere. Here a helium or argon flow of 50 ml/min was
typically used, but measurements in vacuum are also a possibility. Figure 19 shows two
typical heat flow responses from the thermal analysis in this work. The graphs consist of 5
exothermic peaks. It is obvious that there is alteration of thermal stability between the two
analyzed samples. Peak T4 is located at a significantly higher temperature in the thermogram
corresponding to the multilayer, compared to the one of the monolith. The differences are
further discussed in paper 2.
Figure 19. Differential scanning calorimetry measurements. This particular
case shows a monolithic TiAlN and a TiAlN/TiN multilayer with Λ=25/50
nm.
Thermogravemetric analysis (TGA) or TG is a technique in which the mass change of the
sample is recorded while it is subjected to a temperature program. In combination with DSC
the TG is powerful to understand the processes in the thin film at elevated temperature. The
32
Characterization
mass change may be connected to release of material, as for example nitrogen from TiAlCrN
[33].
In this work, 50-70 mg of coating powder was used for each anneal, and before starting
the heat treatment the sample was out-gassed for 12 h at 250 °C. A run consisted of heating
the samples to the maximum temperature with a constant heating rate of 20 °C/min directly
followed by cooling to R.T. Immediately after the first heating/cooling cycle an identical
cycle was performed, which was used for the baseline correction. It should be noted that this
temperature program and baseline correction is only appropriate for coatings which are in a
stable state after the first heat treatment.
6.2.2.
DSC sample preparation of hard coatings
For the experiments in this thesis only the thermal flow in the coating and not the substrate
was investigated. This made it necessary to separate the coatings from the substrate, which
was performed in a combined mechanical and chemical way. For this purpose Fe foils were
used as substrates. After deposition most of the foil is removed mechanically by grinding
after which the remaining substrate is dissolved in 64% hydrochloric acid for 48 h. The
separated as-deposited film, now in the shape of millimeter sized flakes, was collected and
cleaned in distilled water, acetone and ethanol and ground to a fine powder.
6.3. Atom probe tomography
The atom probe tomography technique originates from the field ion microscope which was
first successfully used for imaging atoms in 1955 [122]. By evolution of this technique, the
first successful atom probe tomography experiment was performed in 1967 by Müller et al.
[123]. However, only recently the technique has evolved in such way that a broader range of
materials, such as the coatings in this work, can be characterized. Even though the sample
preparation and time of data collection has been significantly improved over the last few
years, the technique is still considered to be advanced. The main drawbacks with the
technique, compared to e.g. atom resolved transmission electron micrograph (TEM), are the
lack of crystallographic information and that it is destructive. On the other hand it allows for
a precise quantification of the elements and the difficulties with overlapping artifacts, always
present in an image resulting from transmission, is of course not an issue. The number of
publications where the technique is used on materials related to this thesis are still limited.
Some atom probe tomography studies on the decomposition of TiAlN coatings has been
performed by Rachbauer et al. [27, 97, 124] and Johnson et al. [96]. The sample preparation
of coatings is performed with a focused ion beam (FIB) using a beam energy of 30 to 2 keV.
The method consists of a lift out of a sample piece and several milling and sharpening steps
33
Characterization
resulting in a symmetric and sharp tip. A sample tip during milling is seen in Figure 20 (a).
For more details on the sample preparation used in this thesis see paper 5 and Ref. [125].
Figure 20. (a) SEM image Ti0.34Al0.66N/TiN of a multilayer coating during APT
sample preparation and (b) a reconstruction of the same coatings as in (a) using 3%
of the collected ions.
During the APT measurements an applied voltage causes the atoms in the sample, in shape
of the sharp tip, to field evaporate atom by atom. The evaporated ions are accelerated
towards an area detector at which the time of flight and the position are recorded. The
collected data allows for a three dimensional atom resolved reconstruction of the sample tip.
A typical reconstructed tip of a multilayer is presented in Figure 20 (b). In this work the
technique was used in paper 5 to investigate the early stage spinodal decomposition in
TiAlN/TiN multilayers.
6.4. Scanning electron microscopy
Scanning electron microscope (SEM) is, apart from the optical microscope, probably the
most commonly used microscopy techniques. The popularity comes from the simple sample
preparation requirements, the usually well developed and user-friendly interface, and the
straight forward interpretation of the images. The SEM has much better resolution than the
optical microscope but significantly lower compared to a TEM. The SEM is basically an
electron probe scanning over the surface. When the electrons interact with the sample
34
Characterization
several interactions will occur. The three most important for SEM are secondary electrons,
elastically back scattered electrons and the production of x-rays. The secondary electrons are
the ones used for topographic imaging. The number of elastically back-scattered electrons
depends on the atomic number and is thus used for getting elemental contrast, also called Zcontrast imaging. The x-rays are used to determine the chemical composition with EDS
which is further discussed in paragraph 6.6. An SEM was used to get a topographic overview
of the wear after the performed metal cutting test in paper 3.
6.5. Transmission electron microscopy
Transmission electron microscopy is a versatile characterization technique which was
invented already in 1931. In this thesis, the technique was used in all the appended papers.
TEM is invaluable for the analysis of thin films, where it can give information on e.g. the
microstructure, crystal structure, interfaces, defects, binding type and elemental composition.
Since the sample to be analyzed has to be electron transparent, rather time consuming
sample preparation is needed, at least for the coatings in this work. A sample thickness of
<100 nm is desired, and the main technique used here to achieve this is through mechanical
polishing in several steps followed by ion beam milling. For paper 3 and 4, where a limited
sample size and/or a specific location were of interest, FIB was used for the sample
preparation. The FIB instrument can be compared to a SEM, but instead of electrons,
gallium ions are emitted making it possible to mill material at a reasonable rate. For a detailed
description of FIB sample preparation see Ref. [126].
The basics of TEM is rather simple, an electron beam shines through a sample which
results in a projected images on a CCD or on fluorescent screen. In reality the technique is
very sophisticated with apertures, electromagnetic lenses, correctors etc. controlling the
electron beam. The most common imaging mode is bright field imaging, which is performed
with a direct transmitted parallel beam, where the contrast in the image is a result from massthickness and diffraction contrast.
TEM also allows for direct imaging of individual atoms, so called high resolution
transmission electron microcopy (HRTEM). The image contrast is in this case is an outcome
from interference of the electron waves with the sample, which results in a phase shift. By
resolving the crystal lattice planes it is possible to study structural features such as grain
boundaries, defects and multilayer interfaces.
Since the electrons act as waves they will interfere with the atoms in a similar way as the xrays described in paragraph 6.7. It is thus possible to determine the texture of the coating
analyzed in the TEM. In contrast to the x-ray diffraction it is possible to investigate a very
small area of interest with the electron diffraction. This is executed by introducing a selected
35
Characterization
area electron diffraction (SAED) aperture, which limits the beam to a nanometer-sized area,
from which microstructural information is collected.
A modern TEM typically also allows for scanning transmission microscopy (STEM),
where the beam is focused to a probe which is scanned across the sample. The information
from the scattered beam is then collected by a detector allowing Z-contrast and diffraction
contrast imaging. The detector used in this work is a high angle annular dark field (HAADF)
detector. STEM is typically applied together with spectroscopic methods such as EDS and
electron energy loss spectroscopy (EELS) from which the distribution of elements can be
revealed.
6.6. Energy dispersive spectroscopy
When the electron beam in the microscope bombards and interacts with the sample x-ray
photons will be emitted. The emitted x-ray has a characteristic wavelength and energy
depending on the element due to its unique atomic structure. Thus, by collecting the x-rays it
is possible to determine the local or overall composition of the sample. The method is
known as EDS and is especially applicable in the SEM and STEM mode because the prefocused and nano-sized probe. The quantification of light elements, such as nitrogen, is
uncertain because of the emitted x-rays has low energy and will absorbed before they reach
the detector. For the measurements of such elements the characterization is often
complemented with other techniques such as EELS or elastic recoil detection analysis
(ERDA). In this work EDS was used in the STEM to acquire elemental maps or line-profiles
in paper 1 and 4.
6.7. Wide angle x-ray scattering
Diffraction occurs as waves interact with periodic structures with a repeated distance about
the same as the wavelength. X-rays have wavelengths (λ) in the order of a few Å, i.e. the
same as the inter-atomic distances in most crystalline solids. This result in that x-rays can
diffract constructively from solids which have regularly repeating atomic structures.
Accordingly, solids with no long range orders, such as amorphous materials, can usually not
be
characterized.
36
Characterization
X-ray source
detector
θ
2θ
d
Figure 21. Schematic illustration of the θ-2θ setup.
Dashed lines represent the x-rays and the dots the atoms.
The wide angle x-ray scattering (WAXS) is probably the most commonly used technique,
both in academia and industry, to determine the crystal structure of solids. This is due to that
it requires little or non sample preparation and it is non-destructive. Most of the x-ray
scattering in this work was performed on a laboratory scale θ-2θ setup, schematically
illustrated in Figure 21. This method is based on measuring the scattered intensity as
function of the scattering angle 2θ. By applying Braggs law, Eq. 6, the distance (d) between
the lattice planes can be determined. The diffraction pattern, consisting of intensity peaks,
will thus be a fingerprint of the specific crystal structure, where each peak will correspond to
a specific lattice plane. The θ-2θ setup only probe lattice planes which are parallel to the
surface, which means that by investigating the intensity difference between the peaks, the
preferred growth orientation of a coating can be determined.
nλ = 2d sin θ
6.7.1.
Eq. 6
Residual stress measurements
One of the most commonly used methods for measuring the residual stresses in coatings [41,
127], is the non destructive x-ray diffraction sin2ψ method [128]. Depending on the tilt angle
ψ , the stress in the film will give rise to in a shift of the Bragg reflection. From this shift it is
possible to estimate the stress in the film. In the method, a biaxial stress state is assumed, i.e.
a system with in plane stresses only is considered. The strain, for such model is, assuming an
isotropic system (σ φ = σ xx = σ xy ) expressed as
ε=
(1 + v)
σ φ sin 2 ψ ,
E
Eq. 7
37
Characterization
where E is Young’s modulus, v the Poisson’s ratio and is seen to depend only on the tilt
angle. Since the strain is expressed as
ε=
d φψ − d φ 0
dφ 0
,
Eq. 8
where d φψ is the measured lattice parameter at tilt angle ψ and d φ 0 is the measured lattice
parameter at ψ = 0 . From Eq. 7 together with Eq. 8 d φψ is now to solve as
d φψ =
(1 + v)
σ φ d φ 0 ⋅ sin 2 ψ + d φ 0 .
E
Eq. 9
Now, measuring the d φψ at different ψ and plotting it versus sin 2 ψ will result in a linear
relationship where the gradient, looking at Eq. 10, can be expressed as
m=
(1 + v)
σ φ dφ 0 .
E
Eq. 10
From this equation the stress can be extracted if the Young’s modulus and the Poisson’s
ratio from the literature are inserted. For this method, preferable a high angle Bragg
reflection is selected, as it will result in a higher accuracy for the measurement due to larger
peak shifts induced from the strain. For the TiN based coatings in this work the (422) peak
was scanned. The measurement is usually performed in a high resolution θ-2θ mode with
varying ψ angle, starting with ψ = 0 .
6.8. Small angle x-ray scattering
The wide angle x-ray scattering, discussed in the previous paragraph, gives information about
the electron density on an atomic scale. If the measured angles are decreased to very small
values (0-1°) ordered electron density inhomogenities in the nanometer region will instead be
recorded. This is due to that any scattering process comply with the reciprocity law i.e. an
inverse relationship between the scattering angle and the size of the scattering species. The
fundamentals of the technique can be explained with a simple example, using Braggs law, Eq.
6. Let the wavelength of the x-rays be 1 Å. For crystallographic planes, with d from 1 – 2 Å,
the scattering angle 2θ will thus be around ~20 - 40°. If we instead consider domains in an
ordered arrangement, surrounded by vacuum and with a distance d between the domains of
100 Å, the scattering peak will be located at 0.5°, i.e. in the small angle region.
38
Characterization
The recorded inhomogenites can be a result from e.g. precipitation of particles [129, 130],
voids [131] or phase changes [132, 133]. Because SAXS requires very little sample
preparation and is usually non-destructive, it is used for investigations of a broad range of
material such as polymers [133], metals [129], proteins [134], and ceramics [132] in the form
of solids and liquids. For hard coatings it has so far been applied successfully to follow
decomposition and phase changes in TiAlN [132] and ZrAlN [135]. The x-ray source used
for SAXS in this thesis is a synchrotron, but the technique can be applied also on small
laboratory sources.
Table 1. Scattering densities and intensities of the materials relevant for this work.
Phase
Scattering density
(1010 cm-2)
c-Ti0.50Al0.50N
c-Ti0.34Al0.66N
c-AlN
c-TiN
h-AlN
38.79
37.52
33.82
42.51
26.82
Scattering intensity
(1020cm-4)
matrix
24.70
16.24
143.28
matrix
13.69
24.90
114.49
The technique was used for an investigation of the decomposition of monolithic TiAlN
which, as discussed in paragraph 2.2, results in domains rich of c-AlN and c-TiN. This study
was possible only because the domains have a difference in their electron density i.e. the
scattering density. In the example in the introduction, the particles were surrounded by
vacuum, which is of course not the case inside the TiAlN coating. In this case one instead
has to consider the difference in electron density of the domains ( ρ domain ) and a matrix
( ρ matrix ). The scattering densities of the domains and the original matrix, presented in Table
1, were calculated using the scattering contrast calculator in Igor Pro and the Irena package
[136]. The effective electron density difference is expressed as
∆ρ = ρ matrix − ρ domain
Eq. 11
and the contrast, or more correct, the scattering intensity, as
(∆ρ ) 2 = ( ρ matrix − ρ domain ) 2
Eq. 12
The scattering intensities, calculated with Eq. 12 for the TiAlN coatings, are found in Table
1. According to the table the isostructual decomposition of the TiAlN coating will result in
39
Characterization
differences in electron contrast, which give rise to small angle scattering. An example of this
is seen in Figure 22 [132]. At 849 °C a diffuse donut is present which is due to the scattering
from the domains. The radius of the donut clearly decreases with increasing temperature,
which is due to the evolving and growing domains. The upper limit of the measured domain
size is limited by the central spot and the beam stopper, as seen at 997 °C, where it interfere
with the SAXS pattern.
Figure 22. Evolution of a SAXS pattern resulting decomposition of TiAlN during annealing [132].
To compare different annealing temperatures and composition quantitative data has to be
extracted from the SAXS patterns. The fist step to do this is to plot one dimensional lineouts.
An example of this is seen in Figure 23, showing lineouts from Ti0.50Al0.50N in as-deposited
state and annealed at 900 °C for 1, 20, 35 and 52 minutes. The figure shows a decrease of the
peak position which corresponds to the donuts-radius decrease. The second step consists of
using a model to extract quantitative data. Two models have been used on TiAlN to extract
an average domain radius, the Unified-fit and maximum entropy (MaxEnt), the method used
in paper 5. A brief description of the method based on the article by Jemain et al. [137] is
given below.
40
Characterization
Figure 23. One dimensional line outs of SAXS patterns
corresponding to Ti0.50Al0.50N coating isothermally annealed
at 900 °C, from paper 4.
The method of using maximum entropy for extracting a size distribution from small angle
neutron scattering (SANS) and SAXS data was originally developed by Potton et al. [138].
The model was later implemented and adapted for Igor Pro and the Irena package, and was
first used successfully for characterization of 9Cr-1MoVnb steel [137]. The method is based
on comparing an intensity of a calculated size distribution I cal , with an experimentally
collected intensity, Iexp. If a spherical particle shape is assumed with a diameter D the
scattering from the total specimen can be expressed as
∞
I exp = ∫ G (q, D) S ( D)dD
Eq. 13
0
where G(q,D) is the scattering function, S(D) the size distribution of particles with diameter
D and q the scattering vector. If N is the number of particles per unit volume having a
diameter between D and dD and M is the number of measurements the calculated intensity
can be expressed in a similar way
N
I cal = ∑ G (q j , Di ) S ( Di )∆Di , j = 1,...., M
i =1
Eq. 14
41
Characterization
The maximum entropy method then compares the M number of different measured
intensities (Eq. 13), with the corresponding M calculated intensities (Eq. 14), from domains
distributed at N bins. This comparison is made trough χ 2 statistics given by
 I exp j − I cal j
χ 2 = ∑

σj
j =1 
M




2
,
Eq. 15
where σ j is the standard deviation of the measured intensity. The overall size distribution of
a sample can then be calculated by
N
 S ( Di )∆Di 
S = ∑ ( S ( Di )∆Di )log

b


i =1
Eq.16
where b is a small constant [139]. In paper 4, where this model is used, the size distribution is
fitted using a Gaussian and the peak value is plotted versus time. This allow for a straight
forward comparisons between the long time annealing experiments with many data points.
42
7. Phase-field simulations
Phase-field simulations are used to model and simulate microstructure evolution in materials.
The method originates from the pioneer works of Cahn and Hilliard [140] and Allen and
Cahn [141] dealing with free energy of non-uniform systems and anti-phase boundaries,
respectively. Their works have resulted in a number of phase-field methods for a variety of
microstructural mechanisms e.g. grain growth, wetting, solidification and decomposition.
The methods are valuable for predicting industrial processes such as e.g. sintering in powder
metallurgy [142] or phase transformation in steels [143]. In this work the simulations were
performed using the Cahn-Hilliard model. The aim was to study the spinodal decomposition
of both monolithic and multilayer TiAlN coatings. The technique is well suited for modeling
of spinodal decomposition because of the resulting diffuse interfaces.
7.1. The Cahn-Hilliard phase-field model
The driving force for the simulated decomposition is the minimization of Gibbs’ free energy,
∆G , which is given by
∆G =
1
Vm
∫ (∆G
)
( xk ) + κ ∇xk + ∆Eel dΩ
2
m
Ω
Eq. 17
Where Vm is the molar (m) volume, Gm the free energy of mixing per mole, xk the molar
fraction of element k, κ the gradient energy coefficient and E el the elastic energy per mole.
The integration of the equation is performed over the whole volume denoted Ω . The
∆G m is expressed as
∆Gm = ∆H mix − T∆S mix
Eq. 18
43
Phase-field simulations
where T is the temperature, ∆H mix the enthalpy of mixing and ∆S mix the entropy of mixing.
The enthalpy of mixing is given by
n
∆H mix = x A x B ∑ ( x A − xb ) n Ln ,
Eq. 19
i =0
where A and B denotes the two elements and Ln is a third order Redlich-Kister polynomial
extracted from density functional theory (DFT) data[144]. ∆S mix is given by
∆S mix = − R ( x A ln x A + x B ln x B ),
Eq. 20
where R is the molar gas constant. The second term of the Gibbs’ free energy is the gradient
energy which is expressed as
b2
κ=
2
n
∑ (x
i =0
A
− xb ) n Ln ,
Eq. 21
where b is the inter-atomic distance. The last term in Eq. 17, Eel, i.e. the contribution from
the elastic energy, is calculated using compositional dependent elastic stiffness constants[26]
from Ref.[145] using DFT. To perform a simulation a box with a certain number of nodes is
designed and for each node the variation of composition in time is calculated with the CahnHilliard equation[140] given by

 ∆G m
∂x k
= ∇  M ′∇
− κ ∇x k
∂t
 ∆x k

2
+
∆E el
∆x k

 .

Eq. 22
The original equation has been altered to include elastic energy. M´ is the mobility of the
elements and is given by
M ′ = x A xB
( x A DB + x B D A )
RT
where D is the self diffusivity for the different elements.
44
Eq. 23
Phase-field simulations
7.2. Microstructure evolution of monolithic and multilayer Ti1-xAlxN
To simulate the microstructure of decomposing monolithic TiAlN and TiAlN/TiN
multilayers the method described above were used. All simulations were performed in a
specific crystallographic orientation with constant homogenous nitrogen content throughout
the two-dimensional box. The diffusivity of the moving species, Ti and Al, was
approximated to be equal. To simulate the thermal fluctuations present in an as-deposited
TiAlN coating a random compositional fluctuation was set. This is seen as the noise
presented in Figure 24 (a) showing the starting condition of the box for monolithic
Ti0.34Al0.66N. The box is viewed in the [001] direction and have a size of 50x50 nm2. When
the temperature is increased to 850 and 900 °C the coating decomposes, as seen in Figure 24
(b) and (e). In the initial stage, discussed in paragraph 4.2 and paper 4, the domains are
constant in size but increase in elemental intensity for both the annealing temperatures. The
sequent coarsening stage, discussed in paragraph 4.2.1, where the blue TiN domains are
growing in size, is considerable faster at 900 °C compared to 850 °C. Similar results where
also seen in the experimental results in paper 4. The simulations were essential for paper 4 in
order to understand the radius evolution of the domains extracted from SAXS-data.
Figure 24. Simulated microstructure evolution of Ti0.34Al0.66N (50x50 nm2 sized box), with red
representing Al and blue Ti seen in the [001] direction, after (a) RT and annealing at 850 °C for (b) 1
min, (c) 5 min, and (d) 24 min and at 900 °C (e - f) for the same times. The simulations were
performed by J. Ullbrand.
45
Phase-field simulations
It is also possible to simulate more complex cases, such as quaternaries or multilayer
structures. In paper 5, the spinodal decomposition in a multilayer consisting of TiAlN/TiN
was simulated to understand the experimentally observed evolving microstructure and heat
response. The simulations explained how the interfaces, initial fluctuations and coherency
stresses influenced the presence of SDSD and the resulting microstructure. For an example
of multilayer phase-field simulation see Figure 12.
46
8. Metal cutting
This chapter elucidates the metal cutting operation by coated cutting inserts. The
embodiment of metal cutting is in theory rather simple; material from the work piece is
removed until the desired product dimensions are reached. In reality, the process is far more
complicated and numerous parameters have to be considered to achieve a satisfactory endresult. To improve the cutting performance of the coated tools the involved mechanisms
and the coating behavior has to be understood. Here I will first go through the prevailing
conditions at the cutting insert during operation and then the most common wear
mechanism at a macroscopic level. This is followed by a description, based on the literature,
of the cutting performance of TiAlN and multilayer coatings in general.
8.1. Conditions during cutting
What is apparent is that the coated inserts, in comparison to uncoated, generally have an
enhanced protection against thermal and mechanical loads, i.e. the coating decreases the
physical and chemical interactions between the insert and the work piece. For instance, it has
been discovered that the stress encountered at the edge of uncoated tool can be almost 2
GPa higher compared to a coated tool during operation [5]. In addition, the temperature of
the coated tool has been observed to be considerable lower, >300 °C, for the same cutting
parameters [5]. The lower temperature is primarily attributed to the higher amount of
thermal energy transferred to the chip during operation.
A lot of efforts have been devoted to ascertain the conditions at the cutting insert and in
the coating during operation. It has been particularly important to establish the temperature,
and there are numerous reports on different methods to how to estimate this, e.g. by
theoretical modeling [146], thermocouples [147] or, more modern, thermography [148, 149].
The most common method used today is by utilizing an IR-CCD camera which provides a
temperature map offering relatively good resolution [149]. One should, however, be aware
that even the modern thermograph methods suffer from uncertainties. The generation of
47
Metal cutting
heat has been attributed to the shearing of the work material and the sliding of the chip, and
is thus closely related to the cutting speed, i.e. a higher cutting speed increases the
temperature [150]. Temperatures in the range of 700-1000 °C is usually measured at the hot
spot during modern metal machining operations [95]. Furthermore, it has been revealed that
the temperature is reached rapidly after which a steady temperature is observed [149, 151].
These annealing conditions, particularly the heating rates, are intricate to imitate in a labfurnace but the setup applied in paper 4 is relatively close.
The high temperatures of the cutting insert during operation results in a higher diffusivity
and increased chemical interaction between the coating and the substrate and the coating
and the work piece. An effect of the substrate-coating interaction is frequently seen as a peak
corresponding to Co in x-ray diffractograms of annealed coatings deposited on WC [7, 13].
A proof of this is illustrated in Figure 25, showing how Co has diffused in the grain
boundary into the coating after heat treatments. A higher diffusion rate of Co in the grain
boundaries is expected, considering the local higher defect density. The figure also gives an
example how a multilayer structure can alter the chemical interaction.
Another factor to consider during cutting is the stress distribution. High stresses have been
revealed both by calculations [152, 153] and experiments [95] to influence the
thermodynamics of the cubic solid solution TiAlN. The determination of the stresses
prevailing at the cutting edge is however more complex compared to the temperature
distribution. A model where chip thickness, contact lengths, shear strength of the work piece,
and cutting forces are considered has been used successfully [146]. The stresses during metal
machining has by this model been estimated to 2-6 GPa [95, 154].
8.2. Wear mechanisms
It is not an easy task to make a complete description of the wear during metal machining.
Often a combination of theoretical modeling, materials science, chemistry, heat transfer,
mechanics, and tribology has to be applied to give a complete story. The wear types
presented here are in the dimensions that they are easily to study with the most common
methods for characterization available in the industry, i.e. optical microscopy (OM) and
SEM. There are however more wear mechanisms at a microscopic level, such as micro
cracking and dislocation motions, requiring advanced characterization methods such as FIB
and analytical TEM.
The three most common wear mechanisms discussed in the literature are crater wear, flank
wear, and notch wear. The wear types are illustrated schematically in paper 3. The material
removed from the work piece during metal cutting is called the chip. When the chip slides
along the rake face of the cutting insert, it will cause a significant increase of the temperature.
Due to the increased chemical interactions and the abrasive wear a crater will be formed. An
48
Metal cutting
example of crater wear, from paper 3, is seen in Figure 26. Kramer et al. [155] early suggested
that such wear behavior depends mainly on the solubility of the tool material into the work
piece material i.e. a higher solubility result in increased crater wear. The crater wear is,
logically, also dependent on the mechanical properties of the work piece material and the
combination of adhesion and abrasion. At time of writing, it is not clear exactly which
mechanisms dominate the crater wear in relation to cutting parameters. The break down of
the cutting edge, attributed to crater wear, occurs first when the crater reaches the edge. The
crater wear has consequently no detrimental effect on cutting performance of the insert until
this occurs.
The flank wear is, in contrast to the crater wear, a continuous wear of the cutting edge.
This wear mechanism is generally related to the constant abrasive wear from hard second
phases such as inclusions of carbides and oxides present in the work material.
When machining a work piece material that has been strain-hardened from previous
cutting, e.g. stainless steels, especially in combination with high temperature, notch wear is
likely to occur at the depth of cut. Notch wear primarily depends on the insert geometry and
the oxidation properties of the coating and is not considered in this work.
Figure 25. STEM micrograph and EDS elemental map of Co, Al and
Ti of a heat treated TiAlN coating, showing grain boundary diffusion
of Co. The Co containing substrate is located below the image.
49
Metal cutting
8.3. Cutting performance of TiAlN coatings
When the TiAlN coating was introduced to the hard coating community the improved
cutting performance, compared to TiN, was mainly attributed to the fundamental advantage
that it forms a dense, well adhered and protective Al2O3 film on its surface, and an inner
TiO2 when exposed to high temperatures [15, 156, 157]. The oxides were reported to
prevent diffusion of oxygen into the coating material and thus improving the performance
[15, 156, 157]. It was, however, not clear if the coating actually was subjected to an oxidizing
atmosphere at the insert/work piece contact area. Another early reported advantage for
TiAlN was attributed to its relatively low thermal conductivity, allowing for more heat to
dissipate through the chip removal, resulting in a lower thermal load on the insert [158]. A
feasible explanation of the improved cutting performance were however lacking in the
literature for a long time. In 2003 Mayrhofer et al. made a breakthrough by showing that the
hardness of the TiAlN coating increases upon annealing [6]. The hardness increase was
attributed to the decomposition of the TiAlN into coherent c-AlN and c-TiN domains.
Later comprehensive investigations of the cutting performance and decomposition behavior
were performed where the improvements were assigned to the age hardening [7, 8, 30].
Evidence of an active decomposition of TiAlN during metal cutting was first shown in paper
3 in this thesis, however in a multilayer structure. Recent work by Norrby et al. [95] confirms
that the decomposition is present also in a monolithic Ti0.60Al0.40N after metal cutting.
Furthermore, it has been shown that the high stresses prevailing during cutting can affect
cutting performance of TiAlN by an alteration of the thermal stability. Alling et al. [153]
showed theoretical that the favorable spinodal decomposition will occur earlier. This was
also recently confirmed experimentally by Norrby et al. [95].
Figure 26. Crater wear of (a) a monolithic Ti0.34Al0.66N and (b) a Ti0.34Al0.66N /TiN
multilayer coating (Λ=6+12 nm) from paper 3.
50
Metal cutting
8.4. Cutting performance of multilayer coatings
The improved wear mechanism of multilayers has been explained by, for example, increased
hardness which decreases the abrasive wear resistance [159, 160], altered friction coefficient
[161, 162], and decreased tool/coating interactions [163]. TiAlN based multilayer coating
found in the literature showing improved wear properties compared to monoliths are, for
example, TiAlN/TiAlCN [164], TiAlN/TiNbN [165], TiAlN/TiN [13, 31, 32, 161],
AlN/TiN/TiAlN [166], TiAlN/CrN [49, 167, 168] and TiAlN/Mo [169, 170]. However, in
the majority of those publications the conclusions are based on results from tribological test
methods, such as the pin on disk, which are not directly comparable to metal machining. In a
metal machining operation the coatings are continuously subjected to virgin material, in
contrast to most tribological test methods. However, there exist publications showing an
alteration of the crack mechanisms, during metal machining, due to the interfaces. For
example, in a TiN/TiCN multilayer coating, used in interrupted-cut machining, both the
crack formation and propagation was reported to be suppressed by the layered structure
compared to the monoliths of its constituents [171]. Furthermore, Prengel et al. [16] tested a
multilayer, consisting of layers of TiAlN with different Al content, for a milling operation of
ductile and gray cast iron, with and with out cooling. For the high speed dry milling the
coating was reported to perform significantly better compared to the monolithic TiAlN, due
to its ability to resist micro chipping. The TiAlN-based multilayers in this thesis were tested
with in continuous turning operation on AISI 316L stainless steel. Stainless steel is generally
considered to be a complex material to machine, due to that the chip has a strong tendency
to weld to the flank face of the cutting insert [159]. We showed that the cutting performance
is closely related to the multilayer period, i.e. when the period is decreased both the flank and
crater wear are decreased.
To summarize, the wear mechanisms of a coating during metal machining are very
complex, especially for a multilayer structure. Hence, much research remains to be done on a
microstructural level for both multilayer and monolithic coatings.
51
9. Stabilization of c-Ti0.25Al0.75N
The results presented in this chapter are unpublished and not part of the appended papers
and report on the structure of as-deposited monolithic Ti0.25Al0.75N and Ti0.25Al0.75N/TiN
multilayer coatings. The solid solution c-Ti1-xAlxN can be deposited for x<67 at.% using arc
evaporation [7, 23]. A higher Al content result in deposition of a hexagonal phase or a
mixture of amorphous, hexagonal and cubic phase depending on the deposition technique [7,
23, 24]. Here we investigate if it is possible to deposit cubic layers of Ti0.25Al0.75N in a
multilayer coating with TiN as the second layer type, utilizing the epitaxial stabilization effect.
It has been shown that structures not allowed by the phase diagram can be grown by this
technique [53, 54]. To our knowledge there is no work in the literature doing this by
cathodic arc evaporation for the presented multilayer system. Such study is of interest
because epitaxial stabilization in coatings deposited with industrial arc evaporation systems is
relatively unexplored, but can result in attractive properties [172]. The drawback using
cathodic arc evaporation is the presence of macro particles (see paragraph 5.1.2 for a more
detailed description) which breaks the periodic layer growth and acts as nucleation points
[112] for growth of a not cubic Ti0.25Al0.75N. The study is also attractive since Tantardini et al.
[173] showed that a non-isostructural c-Ti0.7Al0.3N/h-Ti0.3Al0.7N multilayer coatings deposited
by unbalanced dc magnetron sputter, exhibited a hardness of almost 50 GPa.
9.1. Deposition conditions
Coatings were deposited using the Sulzer/Metaplas MZR-323 reactive cathodic arc
evaporation system operating in a N2 atmosphere of 2 Pa, a base pressure of 0.5 mPa and a
substrate bias of -40 V. For the growth of the monolithic coating three 63 mm compound
cathodes of Ti0.25Al0.75 were used. For the multilayer growth the Ti0.25Al0.75 cathodes were
placed opposite to three cathodes of Ti. Cleaned cemented carbide pieces, polished to a
mirror like surface, were used as substrates. To achieve the desired stabilization effect a
multilayer consisting of ~6 nm thick Ti0.25Al0.75N layers were deposited. This was made with
53
Stabilization of c-Ti0.25Al0.75N
a drum rotation of 4 revolutions per minute based on the growth rate of the multilayer
periods seen in paper 2. To control the initial growth to a cubic structure, a ~50 nm thick
layer of TiN was deposited before starting the multilayer deposition.
Figure 27. X-ray diffractograms of monolithic and multilayer
Ti0.25Al0.75N in as-deposited state.
9.2. Microstructure
Figure 27 shows the x-ray diffractograms of the monolithic and multilayer Ti0.25Al0.75N
coatings. The monolithic Ti0.25Al0.75N only shows peaks corresponding to the substrate. This
is what can be expected for an arc evaporated Ti0.25Al0.75N and similar to what Hörling et al.
[7] observed. The multilayer coating show a higher intensity peak between the TiN and cTi0.25Al0.75N both at the 111 and 200 planes. This is similar to what was observed for the
Ti0.34Al0.66N/TiN multilayer with shortest period in paper 2 and is assigned to super lattice
reflections.
54
Stabilization of c-Ti0.25Al0.75N
Figure 28. TEM cross sectional images of (a) monolithic Ti0.25Al0.75N and (b)
multilayer Ti0.25Al0.75N/TiN. Both images are acquired in the same magnifications.
Figure 28 shows cross sectional TEM images (a) of monolithic and (b) multilayer coatings.
The monolithic Ti0.25Al0.75N exhibits a dense fine grained microstructure. The multilayer, on
the other hand, shows a columnar structure. Figure 29 shows a HR-TEM image of a
Ti0.25Al0.75N layer, the neighboring TiN layers, and the corresponding fast Fourier transform
(FFT). The thinner lines to the left indicate the positions of the interfaces between the TiN
and Ti0.25Al0.75N layers. The image reveals coherency across the layers and confirms the
epitaxial growth of a cubic structure expected from the x-ray diffractograms, Figure 27.
9.3. Mechanical properties
Figure 30 shows hardness of the monolithic and multilayer Ti0.25Al0.75N coatings measured
with nanoindentation with a 25 mN load. The data was analyzed by the method of Oliver
and Pharr [121]. The hardness values of cubic Ti0.50Al0.50N and Ti0.50Al0.50N/TiN are inserted
to be used as references. A low hardness of the monolithic Ti0.25Al0.75N is expected since
Tantardini et al. [173] reported a hardness of 16.4 GPa of Ti0.30Al0.70N produced by
unbalanced DC magnetron sputtering. Poor mechanical properties of this coating
composition is also reported from cutting test by Hörling et al. [8] showing less than half of
the tool life time (7 min) compared to the c- Ti0.34Al0.66N (20 min). When the Ti0.25Al0.75N is
layered with TiN an increase of ~5 GPa in hardness is obtained, as seen in Figure 30. This
hardness increase can be attributed to several effects. The first effect arises from the fact that
we have a multilayer coating with several hundred of interfaces acting as crack deflectors and
dislocation barriers. One can expect a large difference in E-modulus between the two layers
[145] i.e. fulfilling a requirement for Koehler hardening [56]. The multilayer effects are
described in more details in paragraph 3.2. There is also a effect from coherence between the
55
Stabilization of c-Ti0.25Al0.75N
layers, which will increase the hardness [12, 26]. The last effect is the contribution from the
overall polycrystalline cubic structure present in the multilayer but not in the monolith. The
cubic phase of Ti1-xAlxN is well known to be harder than the one with a present hexagonal
phase. [6, 13, 27, 30].
Figure 29. HR-TEM image of a Ti0.25Al0.75N and the neighboring TiN
layers showing coherence across the layers. The thinner lines to the
left indicate the positions of the interfaces between the TiN and
Ti0.25Al0.75N layers. Inset shows corresponding FFT pattern with
zone axis [110].
To summarize, in this chapter we show that isostructual c-Ti0.25Al0.75N/TiN multilayers can
be grown by cathodic arc evaporation using the epitaxial stabilization effect. This is
confirmed by both x-ray diffraction, showing peaks corresponding to the cubic phase, and
HR-TEM showing a coherent growth with the c-TiN layer. The FFT of the layers further
confirm this showing a (110) cubic pattern. We also reveal that the hardness of the cTi0.25Al0.75N/TiN multilayer is comparable to monolithic c- Ti0.50Al0.50N i.e. ~5 GPa higher
than monolithic Ti0.25Al0.75N.
56
Stabilization of c-Ti0.25Al0.75N
Figure 30. Hardness of monolithic and multilayer Ti0.25Al0.75N
measured with nanoindentation. Hardness values of isostructural
Ti0.50Al0.50N, multilayer and monolithic, are added as reference.
57
10. Summary of papers and contribution to the field
This chapter gives a summary of the included papers and my opinion how the results may
contribute to the community.
10.1.
Paper 1
Direct observations of the decomposed microstructure of Ti0.34Al0.66N, in terms of elemental
contrast, were lacking in the literature before this publication. Cubic metastable
Ti0.34Al0.66N/TiN multilayers with layer thicknesses of 25 and 50 nm, respectively, were
grown by reactive arc evaporation using Ti0.33Al0.67 and Ti cathodes in a N2-atmosphere.
XRD and TEM revealed that the metastable c-Ti0.34Al0.66N layers decompose into c-TiN rich
and c-AlN rich domains with retained lattice coherency after annealing at 900 °C for 2 h.
Elemental mapping by EDS showed a homogenous distribution of Ti and Al in the asdeposited 25 nm Ti0.34Al0.66N layers. In the annealed specimen the Ti0.34Al0.67N had
decomposed into domains of high Al content surrounded by areas of low Al and high Ti
content. The resolution of the STEM/EDS image is sufficient to expose chemical diffuse
boundaries from an expected spinodal decomposition process. However, in these
experiments possible projection of overlapping particles contributing to the diffuse
boundaries could not be ruled out. Thus, in the investigations of the interfaces by EDS and
HR-TEM there was nothing that contradicted the presence of spinodal decomposition.
The results in this paper showed that the TiAlN-layer decompose to well defined AlN and
TiN domains. This gave an estimation of the size and shape of the domains after 2 hour of
annealing. The observation motivated time resolved studies of the microstructure evolution
of Ti0.34Al0.66N, but also investigations on how the decomposition is affected by the
multilayer interfaces.
59
Summary of the papers and contribution to the field
10.2. Paper 2
There exist numerous publications showing how the mechanical properties of a coating are
altered with a multilayer structure, see paragraph 3.2 and 8.4. However, there are few
investigating how the multilayer structure and the period length influence the thermal
stability and age hardening of the coatings. To investigate this, cubic monoliths of
Ti0.34Al0.66N and multilayers of Ti0.34Al0.66N /TiN with three different periods were grown by
reactive arc evaporation. The multilayers were synthesized by mounting the substrates on a
single axis rotating drum set to rotate 1, 2 and 4 times per minute. This resulted in multilayer
periods of 25/50, 12/25 and 6/12 nm with the thinner layer being the Ti0.34Al0.66N.
X-ray diffraction revealed that the Ti0.34Al0.66N in the multilayer decomposes in the same
two steps seen in the monolith i.e. first to c-AlN and c-TiN followed by a transformation to
h-AlN [6, 7, 26]. DSC showed that the first step of decomposition in the multilayers is
shifted towards lower temperatures. The multilayer coatings further showed, in contrary to
the monolith, increasing h-AlN diffraction peak intensity between the diffractograms of
films heat treated at 1000 and 1100 °C. This suggested that the transformation occurred later,
or slower, in the multilayers. The DSC measurements confirmed the XRD data, showing
that the phase change was shifted to higher temperatures compared to the monolithic
Ti0.34Al0.66N. It was also shown that the hardness drop occurred at higher temperature in the
multilayer coatings, which was in line with the measured heat responses. STEM showed that
h-AlN domains in the multilayers are confined by the TiN layers, i.e. the growth was stopped
in the direction perpendicular to the multilayer interfaces. With this study we showed that
the age hardening and decomposition behavior of Ti0.34Al0.66N can be significantly affected
by a multilayer structure.
10.3. Paper 3
The aim of this study was to investigate how the change in thermal stability and age
hardening, seen in paper 2, affects the cutting performance of the Ti0.34Al0.66N/TiN coating
compared to monoliths of Ti0.34Al0.66N and TiN. The multilayer structures of the coatings as
investigated in paper 2, were deposited on pressed and sintered WC-Co milling inserts
(geometry CNMG120408-MR3). The cutting performance of the inserts was evaluated with
continuous turning of AISI 316L stainless steel with a cutting speed of 250 m/min, feed of
0.15 mm/rev and with a 2 mm depth of cut. TEM specimens from the coating on the worn
cutting insert were prepared by a FIB.
A decrease of multilayer period resulted in both improved resistance to flank and crater
wear. The multilayer with period Λ=6+12 nm showed similar flank wear resistance as a
monolithic Ti0.34Al0.66N coating deposited under identical deposition conditions. All the
60
Summary of the papers and contribution to the field
multilayers, regardless of multilayers period, showed improved crater wear resistance to the
Ti0.34Al0.66N monolith. TEM studies revealed a retained multilayer structure with a varied
defected density in the coating exposed to 15 min of continues wear. HR-TEM showed local
coherency over the multilayer interfaces both in as-deposited state and after the continuous
turning. Further, both coherency and incoherency inside the Ti0.34Al0.66N after the cutting test
was observed. STEM imaging and EDS mapping revealed that the layer has decomposed to
Al-rich and Ti-rich areas.
With our study we showed that there is a connection between the multilayer period and
the cutting performance. Furthermore it revealed that there is a stress relaxation and
decomposition of Ti0.34Al0.66N active during metal cutting. The results are important for the
increased understanding of the cutting behaviour of the widely used Ti0.34Al0.66N coatings.
10.4. Paper 4
The aim of this study was to increase the understanding of the microstructural evolution
during the isostructural decomposition of TiAlN and how it is influenced by composition
and isothermal annealing. Two compositions, Ti0.33Al0.67N and Ti0.50Al0.50N, were studied by
in-situ small angle x-ray scattering (SAXS) using a synchrotron source. Phase-field simulations
were used to understand the experimental results. We showed that the isostructural
decomposition occurs in two stages; spinodal decomposition (initial stage) and coarsening
(latter stage). During the initial stage, spinodal decomposition, of the Ti0.50Al0.50N alloy, the
phase separation proceeded with a constant compositional wavelength of ~2.8 nm of the
AlN- and TiN-rich domains. The time of the initial stage depended on the temperature as
well as the composition, and was shorter for the Ti0.33Al0.67N coating. Following the initial
stage, the AlN- and TiN-rich domains coarsened. The coarsening process is kinetically
limited by the diffusion, which allowed us to estimate of the diffusivity constant and the
activation energy for the metals in the coatings.
From an application point of view, these findings are important because they imply that
already after a short time of metal cutting, considering that the temperatures may reach
above 900 °C [147], the microstructure of the coating is in a coarsening stage.
10.5. Paper 5
In this study, the presence of surface directed spinodal decomposition in arc evaporated Ti1xAlxN/TiN multilayers, with two compositions, x=0.67 and x=0.50, was investigated using a
combination of experiments and phase-field simulations. Such study is of interest since
simulations shows that the kinetics of the spinodal decomposition and the resulting evolving
microstructure can be significantly affected by the presence of an interface or a surface. The
characteristics of interface controlled decomposition are the formation of a layered
61
Summary of the papers and contribution to the field
microstructure parallel to the interface i.e. a dominant wave vector directed normal to the
surface.
DSC revealed that the isostructural spinodal decomposition to c-AlN and c-TiN in the
multilayers occur at the same temperature regardless of composition. The onset was located
at a lower temperature compared to the monolithic coatings. Z-contrast STEM imaging
confirmed this by showing a decomposed structure of the multilayers at a temperature where
it was not present in the monoliths. Furthermore, the thermograms show that the
decomposition occurs over a larger temperature range in the multilayers, in comparison to
the monoliths. This is in accord with the phase-field simulations showing longer
decomposition time of the multilayers. 3D atom probe measurements revealed an AlN rich
layer followed by an enriched TiN-layer at the interface in the decomposed Ti0.34Al0.66N/TiN
multilayer, which is in close agreement with the simulated microstructure using large
elemental fluctuations in the initial stage.
The results in this work propose an underlying mechanism for the altered thermal stability
of the multilayer coatings. Since it has been shown that microstructural features such as grain
boundaries might initiate SDSD [104], the understanding of the decomposition type is
important also when considering monolithic TiAlN.
62
11. Future work
This chapter gives an outlook of the possibilities for future work, based on the results
presented within this thesis.
11.1. In-situ decomposition studies
The results in paper 4 show that the coarsening rate of the domains in TiAlN resulting from
spinodal decomposition is significantly increased with temperature. It is further observed,
that the Ti0.50Al0.50N has a period of time of the spinodal decomposition with a constant
compositional wavelength. What is lacking in the literature at the moment is in-situ imaging
of the decomposition, i.e. a motion picture of the evolving microstructure. A modern STEM
equipped with a high temperature sample holder, can provide this. Paper 4 and Figure 11 in
chapter 4, is of great importance for such study, since they allow for selection of appropriate
temperatures and magnifications. Furthermore, an in-situ STEM study of the decomposition
in the multilayers could possibly resolve the evolving SDSD nanostructure, discussed in
paper 5.
11.2. Wear behavior
The wear behavior of TiAlN/TiN multilayers with different periods was investigated in
paper 3. A more detailed study of the microstructure after cutting should be performed, to
increase the understanding of the cutting behavior of multilayers. Such study should contain
investigations of multilayer coated cutting tools, exposed to a series of much shorter
machining times compared to the ones used in paper 3. This is based on the results in paper
4 and 5, showing that the decomposition of the Ti0.34Al0.66N, especially in multilayers, occur
at very short annealing times. This is in line with the results of Norrby et al. [95] showing a
coarsened decomposed microstructure of Ti0.40Al0.60N after only 10 minutes of continuous
cutting. In such study, also the chemical interaction between the cutting insert and the work
piece with the coatings should be considered and investigated. The motivation for this is that
63
Future work
multilayer structures can work as diffusion barriers as discussed in chapter 8. Such
investigation will give a more detailed explanation of the improved crater wear resistance of
the multilayer coatings.
11.3. Mechanical properties
In paper 2 we showed that the age hardening of Ti0.34Al0.66N/TiN multilayers was more
pronounced than monolithic Ti0.34Al0.66N, i.e. it occurred over a wider temperature range and
the relative hardness increase was larger. A study should be performed using FIB sample
preparation and TEM on an indent. Cross sections of indents allow for investigation of the
contact induced deformation mechanisms of coatings. Recently Verma et al. [174] showed
that columnar TiAlN/TiN multilayers, similar to the ones investigated in this work, provides
a more distributed columnar sliding, which reduced the shear cracking. Furthermore, they
showed that interfacial dislocations provide a stress relief mechanism by enabling lateral
movement of material. It has also been shown that at higher loads the main fracture
mechanism consists of crack propagation along the columns while lower loads results in
plastic yielding of the top layers [175]. A comparative study, of as-deposited and
decomposed multilayers, investigating the crack propagation and micro mechanisms during
contact deformation, can give a more detailed explanation of the improved mechanical
properties upon annealing. A similar study on an age hardened Ti0.34Al0.66N monolith, i.e. an
investigation of the crack behavior after annealing, is also interesting and lacking in the
literature.
11.4. Surface directed spinodal decomposition
Paper 5 investigates the presence of SDSD in TiAlN / TiN multilayers. A layer rich in AlN
was observed at the multilayer interfaces. The throughout periodicity which has been
observed in simulations and some experimental results of other material system undergoing
SDSD was, however, not present. This is due to the high initial elemental fluctuations and
high defect density introduced during growth. A similar study should be performed on
TiAlN/TiN multilayers with lower as-deposited elemental fluctuations and dislocations
density. A more homogenous coating can be grown with changed deposition parameters,
such as bias and substrate temperature. An alternative is to use reactive sputtering, allowing
for growth of coatings with much lower defect densities and better interface quality
compared to the ones investigated in paper 5. Such study will explore to what extend the
decomposing structure can be influenced by an interface.
64
Future work
11.5. Improved thermal stability by alloying
It was shown in paper 2 that the unfavorable transformation from c-AlN to h-AlN is
suppressed in the multilayer coatings compared to the monolithic coating. A similar
alteration has been observed in TiAlN alloyed with Cr [176]. Furthermore, other studies
have shown that the spinodal decomposition can be significantly influenced by alloying [3438]. Based on these publications a study of a TiAlXN/TiN multilayer should be performed
to investigate if there is a possibility for cumulative attractive properties from the multilayer
structure and the alloying elements. Another approach is to replace the TiN layer which have
poor mechanical properties and low oxidation resistance. The multilayer should, based on
the results in paper 2 and Ref. [33, 176], have a period of ~15 nm and a relatively low
percent of the X element. The characterization should be performed using nanoindentation,
DSC and STEM investigation.
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