1
Introduction
Metastable face-centered cubic (fcc) Ti1-xAlxN hard coatings deposited by physical
vapor deposition (PVD) on tungsten carbide substrates are extensively used in various
machining applications especially for dry or high speed cutting. Compared to Ti1-xCxN
or CrN, Ti1-xAlxN protective coatings exhibit better oxidation resistance, higher
hardness and thermal stability [1,2]. At elevated temperatures, nanometer-size fcc-AlN
and fcc-TiN domains are formed in the Ti1-xAlxN matrix by spinodal decomposition accompanied
by coating self-strengthening and hardness increase. Further heating results in the
formation of thermodynamically stable wurtzite (w) AlN and fcc-TiN phases [2–5].
Although Ti1-xAlxN coatings have been intensively investigated, only little is known
about the influence of the phase transformations and defect recovery effects on the
residual stress state in the coatings. Phase transformations in coatings are expected
to lead to very pronounced changes in residual stresses especially when there are
large differences in molar volumes of the involved phases. For example, in the case
of AlN the volume occupied per atom in the wurtzite crystal structure (w-AlN) is about
20% larger compared to the rocksalt structure (fcc-AlN) [6,7]. The transformation
from fcc-AlN to w-AlN results therefore in a large volume expansion.
Up to now, most of the temperature-dependent investigations on hard coatings have
been performed on compact layers heated homogeneously across large areas [8,9]. In
industrial applications, however, temperature induced phenomena occur locally in the
cutting zone at length scales of few tens of μm up to few mm, leading to complex three-dimensional
temperature distributions with maximal temperatures of more than 1000 °C reached within
milliseconds [10–13]. There are however very few experimental studies dealing with
the consequences of local thermal loads. As an important tool to simulate thermal
processes in hard coatings, cyclic laser thermal pulses have been applied [11–14].
In this study, a diode laser is used to simulate local thermal loads in a nano-crystalline
Ti0.52Al0.48N coating on hard metal. The peak temperature in the center of the laser
spot was set to 900 °C resulting in the formation of fcc-AlN and TiN precipitates
in the cubic Ti–Al–N matrix and small w-AlN domains detected by transmission electron
microscopy (TEM). By using a variety of experimental techniques, the aim of this work
is to determine local structure–property gradients across the heated zone and correlate
local phases, hardness and residual stresses. Complementary, residual stresses in
a Ti0.52Al0.48N coating heated homogenously in the temperature range of 25–1003 °C
are analyzed and correlated with the oscillating stress distribution across the laser
heated spot. The main aim of the paper is to understand how the local phase transformation
in Ti–Al–N influences residual stresses and hardness in the coatings.
2
Experimental details
2.1
Sample preparation
Ti0.52Al0.48N coatings with a thickness of 4.5 μm were grown on polished hard metal
substrates (WC–Co) in an Oerlikon Balzers INNOVA deposition system by reactive arc-evaporation
in N2 atmosphere by using four compound targets with a Ti/Al atomic ratio of 50/50.
During the deposition, the current was set to 200 A, the pulsed d.c. bias voltage
was − 70 V (25 kHz), the deposition temperature was 350 °C and the total pressure
of N2 was 3.2 Pa.
2.2
Laser treatment
In order to simulate local thermal loads and induce lateral structure–property gradients,
the as-deposited sample was locally heated with a commercial diode laser DL028Q (Rofin-Sinar
Laser GmbH). The sample was placed on a copper block, which was itself embedded in
a temperature controlled oil bath in order to keep the temperature at the bottom of
the substrate constant at 50 °C. To avoid oxidation the laser irradiation was performed
in vacuum conditions of about 10− 3 Pa. The temperature on the coating surface was
monitored by a calibrated pyrometer and regulated via the laser power. The sample
investigated in this study was heated by a laser beam with a diameter of 6 mm for
30 s with a peak temperature of 900 °C. After the laser irradiation the sample rapidly
cooled down back to 50 °C (cf. temperature profile reported in Fig. 1).
2.3
X-ray diffraction characterization
X-ray diffraction (XRD) characterization was performed at EDDI beamline of BESSY (Berlin,
Germany) using a white X-ray beam of the energy range of 20–100 keV and a diameter
of about 0.5 mm [15]. The relatively high penetration depth of the synchrotron radiation
allowed collecting the diffraction signal from the coating and the underlying substrate
during one energy-dispersive measurement. In energy-dispersive diffraction, the relation
between the lattice spacing d(hkl) of (hkl) crystallographic planes and the corresponding
diffraction line E(hkl) is given by
(1)
d
(
h
k
l
)
=
h
c
2
sin
θ
1
E
(
h
k
l
)
where h is the Planck's constant, c is the speed of light and θ is the diffraction
angle [15]. The position-resolved measurements were carried out in symmetric θ/2θ
configuration at a constant 2θ angle of 12° using a lateral step of 0.5 mm and acquisition
time of 100 s.
In order to assess the evolution of the Ti–Al–N 111 peak profile in the energy-dispersive
spectrum (cf. Section 3.2), the peaks were fitted using an exponentially modified
Gaussian function (EMG), which parameters a
0, a
1, a
2 and a
3 denote area, center, width and distortion [16], respectively:
(2)
y
(
x
)
=
a
0
2
a
3
exp
a
2
2
2
a
3
2
+
a
1
−
x
a
3
erf
x
−
a
1
2
a
2
−
a
2
2
a
3
+
a
3
a
3
Complementary, in-situ high-temperature X-ray diffraction (HT-XRD) was used to determine
X-ray elastic strains in an as-deposited Ti0.52Al0.48N sample which was repeatedly
heated in argon atmosphere in the ranges of 25–693, 25–904 and 25–1003 °C [17,18].
Before and during the strain characterization, the temperature was kept constant.
The exposure time per temperature step was about 15 min. The actual sample temperature
was monitored by a thermocouple attached to the coating surface with an accuracy of
± 2 °C.
The in-plane residual stresses in the coatings and in the underlying substrate were
evaluated from Ti–Al–N 111 and WC 100 diffraction lines by applying the sin2
ψ method [19]. The stress states of both coating and substrate were assumed to be
in-plane isotropic with σ
11 = σ
22 = σ. Single crystal elastic constants of TiN and WC were taken and converted into
X-ray elastic constants (XECs) presuming the Hill grain interaction model [19–22].
Since XECs of the characterized Ti–Al–N coatings may differ from those of TiN, X-ray
elastic strains are presented for a comparison.
2.4
Nanoindentation
The hardness across the laser treated area of the Ti–Al–N coating was determined position-resolved
by instrumented nanoindentation using an Ultra-Micro Indentation System (UMIS, Fischer-Cripps
Laboratories Pty Ltd, Sydney, Australia) equipped with a Berkovich diamond indenter
at a constant load of 15 mN, which was chosen so that the hardness of the layer is
not affected by the underlying softer substrate. The maximum indentation depth varied
between 120 and 150 nm depending on the varying mechanical properties of the heat-treated
layer. The nanoindentation data were analyzed according to the Oliver–Pharr method
[23] including an area function, which was determined experimentally using a silica
standard of known elastic modulus.
2.5
Raman scattering
Raman spectra collected across the heated spot were obtained in a micro-Raman spectrometer
(LabRAM HR800, Horiba Jobin Yvon, Villeneuve d'Ascq, France). Monochromatic laser
light provided by a 514.5 nm air-cooled Ar-ion laser was focused onto the specimen
surface by means of a 100 × microscope objective (numerical aperture (NA) = 0.80).
The laser nominal power was set to 41.5 mW at the source. Spectra were collected in
backscattering geometry with the aid of a Peltier-cooled CCD camera (spectral resolution
0.53 cm− 1/pixel) and analyzed in a commercial software environment (Labspec 4.02,
Horiba Jobin Yvon) with a collection of Gaussian–Lorentzian functions.
2.6
Transmission electron microscopy
Transmission electron microscopy (TEM) specimens from irradiated and not irradiated
sample areas were prepared by mechanical polishing and subsequent ion milling. A TEM/STEM
JEOL 2100 F, operated at 200 kV and equipped with an image-side Cs-corrector (CEOS)
and energy filter (Gatan, Tridiem), was used. The atomic resolution of this microscope
at 200 kV is better than 1.4 Å. The high-resolution transmission electron microscopy
(HR-TEM) images presented here were recorded on a 2 k × 4 k pixel CCD camera at a
magnification of 800.000 using an acquisition time of 1.0 s for each image and under
a negative Cs imaging condition.
3
Results
3.1
Damage inspection and chemical analysis
A radial symmetric color gradient can be observed on the surface of the laser treated
coating (Fig. 2), which could suggest the presence of phase gradients. Outside the
irradiated zone, the coating exhibits a gray color (corresponding to the as-deposited
state) turning into yellowish-brown in the center of the laser spot. A quantitative
representation of the coating color across the irradiated spot is presented in Fig. 3
in terms of red, green and blue values, derived from the digital image (Fig. 2). On
the investigated sample, no cracks or signs of damage of the coating surface have
been observed by light optical microscopy and scanning electron microscopy. Moreover,
three-dimensional surface topography characterization performed using confocal optical
microscopy revealed that the surface roughness of the laser heated sample was not
affected by the laser. Energy-dispersive X-ray spectroscopy (EDX) showed that the
coating had a composition of Ti0.52Al0.48N. Position-resolved EDX analysis across
the heated spot (Fig. 2) did not indicate any lateral gradients of elemental composition.
3.2
Microstructural changes
In Fig. 4a, the diffraction pattern collected from the not irradiated area of the
laser heated sample revealed 111 and 200 peaks corresponding to the as-deposited fcc-Ti0.52Al0.48N
with NaCl-type structure.
In Fig. 4b, the diffraction pattern acquired from the center of the heated spot is
depicted. It can be seen that the 111 and 200 peaks are broader compared to the as-deposited
state indicating the presence of fcc-AlN and fcc-TiN nanocrystalline domains, which
are known to form in an early stage of the decomposition of metastable Ti–Al–N [2–4,6].
Since the peak positions of the newly formed domains and the matrix peaks are within
a narrow energy range in the energy-dispersive diffraction pattern, the peaks from
the individual phases overlap leading to the observed broad diffraction peaks in Fig. 4b.
The presence of fcc-AlN nanocrystalline domains can be identified e.g. by the slightly
higher intensity at the right-hand shoulder of the matrix diffraction peaks. The contribution
of fcc-TiN can be clearly seen by the pronounced left shoulder of the 111 and 200
cumulative peaks. Small domains of w-AlN, which were detected by plane-view HR-TEM
investigations in the center of the heated spot (cf. Section 3.7) were not resolved
in the diffraction patterns.
In Fig. 5, a compositional image compiled of energy scans collected at sin2
ψ = 0.4 (that corresponds approximately to the Ti0.52Al0.48N unstrained direction)
documents the development of the fcc-Ti–Al–N 111 peak across the heated spot. The
peak appears broader in the sample center at x = 0 with a pronounced left shoulder
corresponding to fcc-TiN. Moreover, Fig. 5 indicates an increase of the lattice parameter
(shift to lower energy values) in the center and a slight lattice relaxation (shift
to higher energy values) at about x = 2 and − 2 mm.
In order to assess the structural changes across the heated spot, Ti–Al–N 111 peaks
were fitted using the EMG function and the peak asymmetry (expressed by the parameter
a
3) and the integral breadth were evaluated (Eq. (2)). The results in Fig. 6a and b
indicate that the peak asymmetry and breadth start to increase at x = ± 2 mm off the
laser spot center and reach maxima in the center, which is accompanied by the coating
color change (Fig. 3). As mentioned above, the increase of the integral breadth towards
the center of the heated spot indicates the formation of fcc-AlN and fcc-TiN domains.
Also strains of second and third order associated with the decomposition may contribute
to the peak broadening.
3.3
Residual stresses across the laser heated sample
The in-plane residual stresses σ and the corresponding X-ray elastic strains ε
111 for the Ti–Al–N 111 reflection, determined position-resolved across the laser
treated coating, are presented in Fig. 6c. The stress profile can be divided into
three segments. In zone I (representing the sample edge which was not directly irradiated
by the laser), the compressive residual stress of about − 1000 MPa represents a superposition
of growth stress (originating primarily from the ion bombardment during the coating
growth) and thermal stress (formed during cooling down from deposition temperature
T
D of 350 °C to room temperature T
RT).
The thermal stress σ
Thermal formed during the cooling can be expressed as follows
(3)
σ
Thermal
=
E
TiAlN
1
−
ν
TiAlN
α
WC
−
α
TiAlN
T
RT
−
T
D
where E
TiAlN and ν
TiAlN represent the elastic modulus and the Poisson's ratio of the coating, respectively.
Considering that the coefficient of thermal expansion (CTE) of TiN (used as an approximation
of CTE of Ti–Al–N) α
TiN = 9.35 × 10− 6 K− 1 is larger than CTE of WC α
WC = 4.8 × 10− 6 K− 1
[24], tensile thermal stress formed in the coating during cooling. The observed compressive
stress in the not directly irradiated coating at room temperature (Fig. 6c) indicates
a dominant contribution of the compressive growth stress to the total stress.
In zone II (Fig. 6c) from x = − 4 to − 2 and 2 to 4 mm (representing the edges of
the irradiated area), a decrease of the total compressive stress can be observed,
which is associated with thermally activated recovery processes (cf. [2,6]). The compressive
stress relaxation involves annihilation and migration of deposition-induced defects
and other mechanisms resulting in the reduction of the coating volume (cf. our previous
reports on CrN [25,26]). Interestingly, the lattice recovery did not result in a significant
decrease of the integral breadth (Fig. 6b) indicating an interplay between recovery
processes decreasing the integral breadth and the formation of nm-sized precipitations
in the early stage of the decomposition increasing the integral breadth.
In the zone III (Fig. 6c, from x = − 2 to 2 mm), one can expect that the compressive
residual stress increase was induced by the formation of nanocrystalline domains of
w-AlN (as detected by HR-TEM, cf. Section 3.7). The 20% larger volume occupied per
atom of the wurtzite structure compared to the rocksalt structure resulted in an increase
of the coating volume. Since the coating was firmly attached to a stiff substrate
and thus constrained from relieving the accumulated stress, compressive stress developed.
The growth of the w-AlN domains on the expense of fcc-AlN, corresponding to a more
progressed state of the decomposition (towards the center of the heated spot), further
increased the compressive residual stress.
The formation of defects coming along with the film decomposition may as well result
in further increase of the compressive residual stress, presuming that the sum of
all defects causes a coating volume expansion. For instance, strains and dislocation
networks can be clearly seen in the plane-view HR-TEM image recorded in the center
of the heated spot (cf. Section 3.7).
Due to the relatively large X-ray beam penetration depth τ, it was possible to determine
the residual stress as a function of τ in the underlying WC–Co substrate (Fig. 7).
The initial slightly compressive stress in the non-irradiated sample zone of about
− 200 MPa (introduced by mechanical polishing of the substrate before the deposition
process) turns into tensile stress of about 500 MPa in the center of the laser irradiated
zone. The tensile stress development can be explained in the following way (cf. [13]).
During the laser pulse, the substrate center faces a higher temperature than the border.
The thermal expansion of the material in the center is suppressed by the colder surrounding
leading to the formation of compressive stress and a subsequent plastic deformation
in the heated area. During cooling down to room temperature, the elastically deformed
surrounding material contracts and induces tensile stresses in the plastically deformed
sample center. The data in Fig. 7 document that the plastic deformation of the substrate
caused the formation of tensile stress at depths of more than 6 μm.
It is obvious that the substrate plastic deformation influenced the residual stress
state in the coating (Fig. 6c). It can be expected that the observed tensile stresses
caused a relaxation of compressive stresses in the center of the irradiated spot.
3.4
In-situ high-temperature residual stress characterization
In order to understand the residual stress profile across the laser heated spot (Fig. 6c),
a further sample with the same Ti0.52Al0.48N coating on hard metal was repeatedly
thermally cycled and in-plane X-ray elastic strains as well as residual stresses were
evaluated by analyzing the Ti–Al–N 111 peaks (Fig. 8). During the first cycle in the
range of 25–693 °C (circles in Fig. 8), the as-deposited residual stress of − 900 MPa
changed to -50 MPa after cooling to 25 °C. The stress relaxation is associated with
defect recovery as observed in zone II in Fig. 6c.
The subsequent heating in the range of 25–904 °C (squares in Fig. 8) induced the spinodal
decomposition and a compressive stress increase to about − 3500 MPa, which decreased
thermo-elastically down to − 2642 MPa after cooling down. The compressive stress increase
from − 2059 to − 3429 MPa was observed even after the temperature decreased from 904
to 837 °C. This documents that the spinodal decomposition was not completed at 904 °C
to a full extent, as expected [3,6]. During the third cycle in the range of 25–1003 °C
(triangles in Fig. 8), the stresses in the coating changed exclusively thermo-elastically
without a significant change (hysteresis) after cooling.
In the diffraction data collected during the second and third temperature cycles,
not only peaks from cubic phase were observed but also very weak and broad peaks indicating
the presence of w-AlN were detected at approximately 22 keV, explaining the high compressive
residual stress formation in the coating.
3.5
Hardness across the laser heated sample
In Fig. 6d the hardness profile across the laser-affected zone is presented. The hardness
of the not directly irradiated coating is about 35 GPa. At x = ± 2 mm the hardness
starts to increase reaching a maximum of 47 GPa in the center of the laser spot. The
hardness increase can be related to age hardening effects coming along with the formation
of fcc-AlN and fcc-TiN domains and in particular to coherency strains between the
matrix and coherent precipitations [3]. The large strains of first order (cf. Section 3.4)
and lattice distortions (cf. Section 3.7) observed in the center of the heated spot
additionally favor the hardness increase, being dominant over mechanisms that are
known to decrease the hardness of the coating such as the formation of w-AlN and/or
the introduction of misfit dislocations, relaxing the coherency strains between the
precipitations and the matrix [3]. In zone II (x = − 4 to − 2 and 2 to 4 mm) the stress
relaxation did not induce any significant hardness change in the coating indicating
that there is an interplay between recovery processes causing coating softening and
effects originating from the very early stage of the decomposition increasing the
hardness.
3.6
Raman scattering across the laser heated sample
The aim of the Raman investigations was to obtain more insights into the Ti0.52Al0.48N
decomposition and residual stress formation in the coating. In Ti–Al–N, the presence
of both heavy metal ion and nitrogen ion vacancies reduces the local symmetry of the
crystal, causing atomic displacements having non-zero first order polarizability derivatives,
resulting in Raman activity [27,28]. Representative Raman spectra from the border
and the center of the irradiated coating (Fig. 9) can be discussed in terms of two
main features: (i) vibrations in the 150–300 cm− 1 range are longitudinal and transverse
acoustic modes (LA and TA) determined by Ti and Al atoms, whereas (ii) the modes in
the 400–650 cm− 1 range are longitudinal and transverse optical modes (LO and TO)
associated with vibrations of the N atoms [27–29]. Other low-intensity modes are acoustic
and optic overtones (cf. description of the modes in Table 1).
The normalized Raman spectra in Fig. 10 report the result of a line-scan performed
across the irradiated sample indicating continuous changes especially for TO and LO
modes in the 500–650 cm− 1 range. The spectra from Fig. 10 were fitted and the peak
positions for the two main first-order acoustic and optical modes were evaluated.
In Fig. 11, a consistent shift can be seen in the peak position of both modes with
respect to the irradiated area in the center of the specimen. In general, in Ti–Al–N
the mode shift to higher frequencies is related to compressive stress [29,30]. The
shift shown in Fig. 11 could thus be precipitously associated with the presence of
tensile stresses in the irradiated coating. However, the measured shift (up to ~ 40 cm− 1)
would imply a residual stress level that is unrealistic (in Ti–Al–N, a typical value
of biaxial stress-induced shift is about − 2.2 cm− 1/GPa [28]). Consequently, in order
to evaluate the residual stress magnitudes from the collected Raman spectra (Fig. 10),
values of biaxial stress-induced shifts would have to be considered for every ratio
of fcc-Ti–Al–N, fcc-TiN and fcc-AlN. It is obvious that the phase changes represent
here a dominant contribution in the spectral changes. Fig. 10 indicates that a consistent
sharpening of the peaks appears together with a splitting of the TA/LA peak into two
components at ~ 215 cm− 1 and ~ 330 cm− 1 in the irradiated center. Being Raman peak
width sensitive to the degree of disorder in the lattice, those effects suggest a
higher fraction of crystalline domains and are in accordance with the increasing presence
of the precipitations in the irradiated sample center.
3.7
Transmission electron microscopy
Low magnification TEM images showed that the grain size ranged from hundreds of nanometers
to about 1 μm. A typical Cs-corrected HR-TEM image (Fig. 12) recorded on the as-deposited
sample within one big grain shows homogeneous intensity and a perfect cubic-lattice
atomic structure of the Ti0.52Al0.48N coating. The corresponding indexed fast Fourier
transformation (FFT) image further indicates the single cubic phase present in the
non-irradiated sample area what is in agreement with the XRD data from Fig. 4a.
In contrast, a Cs-corrected HR-TEM acquired from the annealed sample area shows structural
fluctuations (Fig. 13), where some stripes and contrast variations due to the large
strains are clearly visible. The FFT images document the presence of cubic phases
with lattice parameters similar to the cubic Ti0.52Al0.48N. The resolution of the
FFT images does not allow distinguishing between both phases. Some additional tiny
spots appear in the FFT owing to the presence of a structure modulation along a certain
orientation. The analysis shows that the tiny spots originate from w-AlN domains (see
Fig. 13, where a w-AlN lattice is labeled), which construct a semicoherent relationship
with the matrix.
There are numerous dislocations present in the matrix. To visualize the defects, a
displacement map using the Bragg reflection (200) is available, which clearly shows
several dislocations existing in the selected area. The relatively large amount of
such lattice defects observed in the HR-TEM image (Fig. 13) indicates that not all
precipitates are longer coherent and the strains of third order are partly relaxed
by forming dislocation networks at the interface between the precipitates and the
matrix.
4
Discussion
Lateral gradients of residual stress, phase and hardness observed across the laser
treated Ti0.52Al0.48N coating (Fig. 2) reveal the complex nature of thermo-mechanical
processes in locally heated Ti–Al–N on hard metal. Heating above 900 °C results in
the formation of fcc-AlN, fcc-TiN and w-AlN. According to the HR-TEM data (Fig. 13),
the precipitates are not always coherent and there is a significant amount of dislocations
in the annealed area. It is obvious that coherency strains between the coherent precipitations
and the matrix, the local modulations of strain and composition, the presence of atomistic
structural defects as well as relatively large matrix strains of the first order are
responsible for the hardness increase in the irradiated sample center (Fig. 6d). The
TEM data demonstrate that even the relatively short laser pulse (Fig. 1) results also
in the formation of w-AlN phase (Fig. 13) strongly influencing the residual stress
state in the coating (Fig. 6c).
The results in Fig. 7 demonstrate that the local heating caused a local plastic deformation
in the substrate with tensile stresses in the substrate interface area. This effect
can significantly deteriorate the mechanical integrity and the performance of the
coating/substrate composite. One can expect that thermal laser pulses with higher
surface temperatures and/or applied repeatedly would lead to the formation of relatively
large tensile stresses in the substrate/coating composite and cause the formation
of cracks. Consequently, the tensile stresses in the substrate observed in Fig. 7,
as well as our data collected from the CrN/steel system [25], indicate that the substrate
thermo-mechanical behavior and its dimensional changes resulting from the inhomogeneous
heat treatment must be considered when analyzing thermo-mechanical behavior of locally
heated hard coatings. It is not enough to concentrate only on the processes in the
coating, like obstacle controlled plasticity, recovery effects and spinodal decomposition.
In order to correlate the local phases (Fig. 6b), residual stresses (Fig. 6c), hardness
(Fig. 6d) and the temperature applied during the laser treatment, the finite element
method (FEM) was used to model the temperature profile on the surface of the irradiated
specimen using a transient heat transfer simulation. The temperature dependent material
data for the Ti–Al–N coating [12] and the WC–Co substrate [31,32] were taken from
literature. The intensity distribution of the laser was assumed to be constant across
the cross section of the beam with a diameter of 6 mm. This assumption led to a good
agreement of the simulated lateral temperature distribution with measurements performed
on another material [33]. The laser intensity was adjusted in such a way that the
maximal measured temperature reaches 900 °C and is held constant for 30 s. In Fig. 14,
the experimentally determined integral breadth (Fig. 6b), in-plane X-ray elastic strain
(Fig. 6c) and hardness (Fig. 6d) are correlated with the temperature distribution
across the sample as obtained from FEM. The results indicate that the hardness increase
correlates well with the coating phase changes (fcc-AlN and fcc-TiN domains increase
the integral breadth) and reveal the residual stress evolution with temperature. The
results in Fig. 14 also demonstrate that it is possible to perform fast screening
of structure–property relationships in complex materials just by analyzing one sample.
By considering the results from Figs. 6c, 7 and 8, the following phenomena can be
in general assumed to predetermine the volume-averaged residual stress σ
RT of the Ti–Al–N coating at room temperature after the heat treatment: (i) growth
stress σ
Growth which depends predominantly on the growth conditions such as applied bias voltage
and deposition rate, (ii) thermal stress σ
Thermal discussed in Section 3.3, (iii) stress relaxation caused by structure recovery
σ
Recovery at temperatures above T
D, (iv) stress changes caused by phase transformation in the coating σ
Decomposition and (v) stress changes caused by substrate dimensional changes σSubstrate
like local plastic deformation.
From the practical point of view, however, it is more important to know what kinds
of stresses contribute to the actual stress magnitude during the coating service.
The stress–temperature dependencies in Fig. 8 indicate that the coating phase transformation
represent a dominant contribution to the compressive stresses. Similar (in this specific
case) compressive and growth stresses can be considered as favorable. On the other
hand, the recovery effects and substrate plastic deformation can significantly deteriorate
the coating performance by generating tensile stresses. In the extreme case, one can
suppose that the substrate plastic deformation (resulting possible in the formation
of cracks [13]) can induce crack formation in the coating. In the case of cutting
tools, one should also note that cutting forces during the operation influence the
actual residual stress state in the coatings and finally also their lifetime.
From the methodological point of view, by coupling the local (Figs. 4–7) and the in-situ
(Fig. 8) XRD characterization it was possible to interpret the oscillating distribution
of residual stresses (Fig. 6c) across the laser heated spot and understand the role
of the substrate in the thermal treatment. This complex approach of tailoring the
thermo-mechanical behavior of both coating and substrate represents a progress in
contrast to pure modeling approaches where many experimental parameters are usually
not considered.
In the case of nano-crystalline coatings like Ti–Al–N, depth gradients of microstructure
and residual stresses influence decisively the coating function [34,35]. It is obvious
that in the investigated Ti0.52Al0.48N coating, parameters like residual stresses,
phases and microstructure change as a function of the coating thickness. Further research
is therefore necessary to reveal the relationships between the coating structural
depth gradients and mechanical properties.
5
Conclusions
Local lateral temperature gradients induced by laser irradiation of a Ti0.52Al0.48N
coating on hard metal resulted in the formation of pronounced gradients of phases,
hardness and residual stresses which were correlated using scanning characterization
techniques. The results open the possibility to perform post-deposition residual stress
engineering in self-hardened Ti–Al–N coatings by choosing dedicated annealing conditions.
The present study documents that local laser heating is a powerful tool to simulate
local thermal loads in hard coatings. Moreover, the approach demonstrates that by
combining laser heating to synthesize new material modifications and appropriate position-resolved
experimental techniques, it is possible to perform fast screening of structure–property
relationships in complex materials just by analyzing a relatively small sample area.