Phase Transformations in Arc-Evaporated Ti-Si-N Thin Films
Axel Flink1, Manfred Beckers1, Jacob Sjölén2, Tommy Larsson2, Slavomir Braun3,
Lennart Karlsson2, and Lars Hultman1
1
Thin Film Physics Division, Department of Physics, Chemistry, and Biology, IFM,
Linköping University, SE-581 83 Linköping, Sweden
2
3
Seco Tools AB, SE-737 82 Fagersta, Sweden
Division of Surface Physics and Chemistry, Department of Physics, Chemistry, and
Biology, IFM, Linköping University, SE-581 83 Linköping, Sweden
Abstract
(Ti1-xSix)Ny (0≤x≤0.20) thin solid films have been deposited by arc evaporation onto
cemented carbide (WC-Co) and sintered c-BN substrates. X-ray diffraction and TEM
analyses show that low-Si content films up to x=0.09 exhibit a dense columnar singlephase cubic (Ti,Si)N solid solution while 0.14≤x≤0.20 films assumes a defect-rich,
feather-like nanostructure consisting of cubic nanometer size (Ti,Si)N crystallites and
amorphous SiNz:Ti phases. Correspondingly, the N content in the films measured by
elastic recoil detection analysis increases close to linear with increasing Si content from
y=0.99 (x=0) to y=1.13 (x=0.20). X-ray photo-electron spectroscopy reveals tetrahedral
SiNz binding configuration. The nanoindentation hardness of as-deposited films increases
with increasing Si content from 30 GPa (x=0) to 42 GPa (x=0.14). For higher Si contents,
up to x=0.20, the hardness decreases again to 38 GPa. Annealing experiments show that
films retain composition and hardness up to 1000 °C for 0.04≤x≤0.09 due to a
thermodynamically driven phase segregation to residual (Ti,Si)N and a coherent SiNz
grain boundary phase. For x<0.04, the films soften by thermally-induced recovery effects
1
within the film. For x>0.09, Si and N diffuse out of the films along boundaries
accompanied by Co and W interdiffusion from the substrate, yielding recrystallization,
grain boundary weakening and a corresponding decrease of hardness to <28 GPa. Sidepletion is also observed during annealing of films with x=0.17 and x=0.19 on thermally
more stable c-BN and Ta substrates, respectively. This is again accompanied by domain
recrystallization; the hardness is, however, almost retained at 38-40 GPa. At 1200 °C, the
films have fully recrystallized into nanocrystallites of TiN and porous grain boundaries
where Si and N out-diffusion took place. A residual Si content of ?? is still present in the
film.
Introduction
Materials science and advanced surface engineering are employed to develop wearresistant coatings for metal cutting tools with superior thermal properties, chemical
inertness, and mechanical strength [1,2] with the objective to enable higher cutting speeds
with correspondingly higher working temperatures. Such combination of material
properties is sought for in ternary or quaternary ceramic coatings deposited by physical
vapor deposition. In particular, the Ti-Si-N system exhibits interesting properties for
metal cutting applications. Adding Si to TiN results in hardening for TiN-Si3N4
nanocomposites [3,4,5], TiN/SiNx multilayers [6,7,8], and metastable (Ti1-xSix)N solid
solutions [9,10,11]. There is, however, no stable ternary phase reported in the Ti-Si-N
phase diagram [12].
The cubic (Ti1-xSix)N solid solutions are the least explored of the Ti-Si-N structures
mentioned above. Yet, the solid solution hardening is substantial. We previously showed
2
that (Ti1-xSix)N films deposited by arc evaporation at 500 °C exhibit an essentially linear
hardness increase with increasing Si from ~31 GPa in pure TiN to ~45 GPa for
Ti0.86Si0.14N [9]. Moreover, the hardening of the solid solutions is retained during
isothermal annealing at 900 °C in Ar for 2 h. By contrast, TiN softens significantly to ~25
GPa during such annealing because of recrystallization and residual stress recovery.
The upper limit of Si supersaturation in a cubic alloy with TiN is difficult to assess for
analysis reasons. For the possibility of a spinodal decomposition reaction, the range of
NaCl-structure Ti1-xSixN alloy compositions was shown by ab initio calculations to have
a lattice parameter which deviates less than 0.5 % from that calculated for TiN [13,14].
This makes it difficult to detect such phase separation and solid solution alloying by x-ray
diffraction. The strong segregation tendency of the species in a nitride solid solution,
however, will eventually lead to phase transformation into TiN and SiNx. In such a case,
the analysis challenge becomes one of detecting the smallest volumes of SiNx in a
crystalline matrix.
Here, we present a combination of high-resolution imaging and analytical techniques to
investigate both the details of phase transformations for different Si contents, and the
actual nitrogen content in the films, which we hereafter denote as (Ti1-xSix)Ny.
It is found that the forced solid solubility limit for Si in cubic (Ti1-xSix)Ny thin films
synthesized by arc evaporation is x≈0.1 (~5 at.%). As-deposited films with higher Si
content exhibit Si segregation and N uptake in the range 1.09≤y≤1.13 with a concomitant
tendency for the formation of an amorphous SiNz:Ti phase. Nanoindentation experiments
3
show that the hardness of as-deposited films reaches a maximum of 41.8±2.0 GPa for a
Si content of x=0.14. The solid solution films (0.04≤x≤0.09) are structurally stable with
retained hardness after annealing at 1000 °C in Ar for 2 h. Films with x≥0.14 in the asdeposited state, however, exhibit pronounced Si and N loss by out-diffusion and
desorption together with simultaneous recrystallization within the grains in the cubic state
together with Co and W interdiffusion from the cemented carbide (WC-Co) substrates.
Correspondingly, the post-annealing hardness decreases to <28 GPa. On the contrary, the
hardness of films with x=0.17 on thermally more stable sintered c-BN substrates is,
however, more or less retained at 38-40 GPa up to 1100 °C, despite local Si out-diffusion
and segregation, which results in local recrystallized areas.
Experimental details
(Ti1-xSix)Ny coatings were deposited by arc evaporation in a Metaplas MZR323 system
with a base pressure of 5x10-6 mbar. 12x12x4 mm3 cemented carbide (WC-Co(6 wt.%)),
5x5x1 mm3 sintered c-BN-(50 vol.% TiC), and 15x15x1 mm3 metallic Ta were used as
substrates. The WC-Co substrates were ground and polished to a mirror-like finish with a
roughness Ra ≈ 0.01 µm. Both WC-Co and Ta were cleaned ultrasonically in an alkaline
degreasing agent prior to deposition. Nine samples were placed along three vertically
mounted 63 mm diameter targets of composition Ti, Ti0.8Si0.2, and Ti0.75Si0.25,
respectively, to obtain a Ti to Si film composition gradient. For the case of c-BN, the
deposition experiments were performed using 63 mm diameter Ti0.80Si0.20 targets in order
to obtain films of a constant high Si content. In both series, the arc evaporation was
operated in a reactive N2 atmosphere. The substrates were negatively biased to –50 V and
4
mounted on a rotating cylindrical fixture. The substrate temperature was kept at ~500 C
using a combination of plasma heating and a resistive heater.
The film compositions were analyzed by time-of-flight energy and elastic recoil detection
analysis (TOF-E ERDA). The measurements were performed using a 40 MeV
127 9+
I
ion
beam. The TOF-E setup, described in detail elsewhere [15], consists of a 437.5 mm long
telescope with two carbon foil time detectors, which is followed by a ion implanted
silicon charged energy detector with 10x10 mm² active area. The recoiled atoms are
collected at an angle of 45° relative to the incoming beam direction. The incidence angle
of the primary ions as well as the exit angle of the recoils is 67.5° with respect to the
substrate normal. The resulting time-of-flight over recoil energy spectra are transformed
to elemental depth profiles using the newly developed CONTES code [16]. The average
film compositions and statistical errors are obtained by normalizing over the measured
depth profiles over 100-300 nm sampling depth. The statistical error of the measurement
is 1-2 at.%.
The microstructure of the samples was investigated by x-ray diffraction (XRD),
transmission electron microscopy (TEM), scanning transmission electron microscopy
(STEM), and scanning electron microscopy (SEM). XRD was performed using a Bruker
AXS D8-advanced x-ray diffractometer with a line-focus Cu Kα x-ray source. -2 scans
were performed in the 2 range from 20 to 80. The residual stress was calculated using
the sin2ψ method based on the (Ti1-xSix)Ny 422 peak in the as-deposited films with
x≤0.14 and all annealed films. For as-deposited films with x≥0.14, the stress was
calculated based on the 002 peak due to dominant 002 preferred off-plane orientation. For
5
all calculations a The Poisson’s ratio (ν) of 0.22 and a Young’s modulus (E) of 450 GPa
was used, as were determined for TiN [17]. The indicated error bars originate from linear
fitting of the measured data points. The SEM analysis was performed on a LEO 1550
FEG-SEM. An analytical FEI Technai G2 UT FEG microscope equipped with an energy
dispersive x-ray spectrometer (EDX) operating at 200 kV was used for TEM, STEM, and
EDX elemental mapping. For the STEM analysis, a high angle annular dark field detector
(HAADF) was used and a camera length of 220 nm. The mapping was performed with a
pixel resolution of 2 nm over 150x150 nm2 for as-deposited (Ti0.81Si0.19)N1.13 and 1 nm
over 50x50 nm2 for (Ti0.81Si0.19)N1.13 annealed at 1000 °C for 2 h, respectively. Crosssectional TEM/STEM specimens of films deposited on WC-Co substrates were made by
clamping two small sample pieces film to film into a Ti-grid, followed by mechanical
grinding and polishing. Electron transparency was achieved by ion milling using a Gatan
Precision Ion Polishing System. The corresponding plan-view TEM specimens were
made by cutting out a disc followed by the similar steps of mechanical grinding,
polishing, and ion milling. Cross-sectional TEM/STEM specimens of films deposited on
c-BN and Ta substrates were prepared using a Zeiss 1540 EsB CrossBeam focused ion
beam with the so-called lift-out technique [18].
X-ray photoelectron spectroscopy (XPS) measurements were conducted using a Scienta
ESCA 200 with a monochromated Al (1486.6 eV) Kα beam. Survey scans were recorded
in the 0-800 eV energy range with a step size of 0.5 eV for each sample. For accurate
determination of the N1s, Si2p, Ti2p, O1s, C1s, Co2p, and W4f peak shapes,
respectively, local region scans were recorded with a step size of 0.1 eV. In order to
maintain sufficient charge neutralization for samples containing silicon, the use of a flood
6
gun was necessary. The binding energy scale of the narrow scans was calibrated by the
C1s line at 284.7 eV. Scans were recorded both prior to Ar-sputtering in order to avoid
sputter-induced chemical shifts, and after 20 min of Ar-sputtering at PAr =2.5x10-7 mbar
in order to rule out surface-oxide contaminations. The analysis of the spectra obtained
before and after sputtering, however, revealed the presence of sputter-induced chemical
shifts. Hence only scans before Ar-sputtering are discussed in the paper.
Isothermal annealing of the as-deposited films was carried out in a Sintevac Furnace from
GCA Vacuum Industries for a duration of 120 min in Ar at atmospheric pressure in order
to prevent oxidation of the sample surfaces. The individual annealing temperatures (Ta)
were 1000C (WC-Co and c-BN), 1100C (c-BN), and 1200C (c-BN), respectively. The
initial heating rate was 7 °C/min, which was reduced to 5 °C/min at 40 °C below the
respective Ta. After the annealing the samples were cooled in the furnace with typical
cooing times of 1.5 h from 1100 °C down to 500 °C and 4 h from 500 °C down to 100
°C.
The film hardness was determined by nanoindentation using a Nanoindenter XP equipped
with a Berkovich diamond tip. Mechanical polishing on taper sections was performed
prior to the indents in order to reduce the surface roughness influence. 20 indents were
made in each sample with a maximum load of Pmax=25 mN. The indentation procedure
consisted of 5 segments: 1) load to Pmax, 2) hold 10 s, 3) unload to 10% of Pmax, 4) hold
60 s, and 5) unload. Any thermal drift was corrected for at hold segment 4. The average
hardness and its standard deviation were determined from the obtained data following the
method of Oliver and Pharr [19] by fitting 50% of the unloading curve with an
7
exponential function. The reference SiO2 sample hardness was 9.6±0.3 GPa and 9.89±0.3
GPa for the (Ti1-xSix)Ny films deposited onto WC-Co and c-BN substrates, respectively.
Results and Discussion
Composition of the as-deposited films
Table 1 shows the composition of the nine as-deposited of the as-deposited (Ti1-xSix)Ny
films on WC-Co substrates as determined by ERDA. The Si content ranges from x=0 (0.1
at.%) for the sample facing the Ti target, to x=0.20 (9.1 at.%) for the film facing the
Ti0.75Si0.25 target. The corresponding N content increases close to linear with increasing
Si content from y=0.99 to y=1.13. The impurity levels for all as-deposited films were
below 1 at.% with a small [O] increase from ~0.45 at.% to ~0.8 at.% with increasing Si
content.
Table 1 reveals a substantial reduction of Si content in the films compared to the
composition of the targets for the geometry of the deposition experiments. For example,
the film facing the Ti0.75Si0.25 target shows a Si content of only x=0.20. This can be
explained by the different average ionization for the Ti and Si species during arc
evaporation, which are typically +2.1 and +1.4, respectively [20]. Using a substrate bias
of -50 V, the Ti ions will hence impinge on the sample surface with a correspondingly
higher energy causing persisting Ti sub-plantation, while the comparably surface-near Si
will be preferentially resputtered. Hence, the film will have a reduced concentration of Si.
Microstructure of the as-deposited thin films
8
Error! Reference source not found. shows x-ray diffractograms obtained from the asdeposited films on WC-Co substrates. All films have a single-phase NaCl-structure with
a lattice parameter of ~4.24 Å, i.e., more or less identical to stoichiometric TiN (4.24 Å
[21]). For the pure TiN film (x=0), a random crystallographic orientation is observed. The
uptake of Si induces peak broadening and a <002> out-of-plane preferred orientation,
which is nearly complete for x≥0.09.
The residual stress (σ) of films deposited on WC-Co substrates is presented in Error!
Reference source not found.(a). All as-deposited films are under substantial
compressive stress, which gradually increases from x=0 at σ=-2.7±0.2 GPa, up to σ≈-5.5
GPa for x=0.19. One has to keep in mind that a linear fit of sin2ψ method data is only
exact for isotropic materials. The strong preferred orientation in films with x≥0.09 may
therefore reduce the accuracy of the given stress values. The same applies to the use of
the low angle 002 peak, for which ψ-tilt induces smaller peak shifts compared to the high
angle 422 peak. A direct comparison between the 200 and 422 peaks for the as-deposited
x=0.14 film, however, resulted in similar stress values, σ=-4.83±0.28 GPa and σ=4.55±0.53 GPa, respectively, which implies that the results obtained from 002 and 422
are comparable. The residual stress values for the annealed samples will be presented
below.
Error! Reference source not found.(a,b) shows detailed x-ray diffractograms around
the 002 peak of the as-deposited and annealed films on WC-Co. In the as-deposited state,
the full width at half maximum (FWHM) values remain close to constant at ~0.5° from
x=0 up to x=0.04 (Error! Reference source not found.a). Also the peak positions of
9
films with 0.01≤x≤0.09 show no peak shift when compared to TiN (x=0). Recent
calculations by Alling et al. [14] predict a lattice parameter change from 4.255 Å for x=0
to 4.250 Å for x=0.125 when the Si atoms are fixed at NaCl-lattice Ti positions. If the
atoms within the unit cell are allowed to relax, however, the lattice parameter for x=0.125
is predicted to be 4.231 Å. Similar results were obtained with other ab initio methods in,
e.g., [14,22]. These lattice parameter changes correspond to a shift of 0.25° and 0.05° to
higher angles. As the observed FWHMs are bigger than the predicted peak shifts, any
alloying of Si and TiN will be virtually impossible to detect, and hence, no conclusion
whether Si is dissolved in a solid solution or segregated into a-SiNz or other grain
boundary tissue phases can be drawn. As for the resulting residual stress, the data
obtained after annealing will be discussed later.
For Si content x≥0.09, the 002 peak FWHM values increase substantially from ~0.7° to
~1.8° indicating extensive grain refinement. For x≥0.14, the 002 peak position gradually
shifts slightly towards lower 2θ angles, corresponding to an out-of plane lattice parameter
of 4.27 Å for x=0.19. The shift coincides with the increase of compressive stress,
compare Error! Reference source not found.. This could find its explanation in the
formation of a Si3N4 tissue phase with or without Ti, since Si3N4 has a substantially larger
molar volume compared to the NaCl-structure solid solution.
Fractured cross-sections of the as-deposited films were investigated by scanning electron
microscopy. From these, the film thickness was measured to ~1.5-2.5 µm. On this
relatively coarse scale the film microstructures were dense and columnar with
10
incorporated macro particles from the arc evaporation process. Films with x≥0.14 appear
as fine-grained.
Cross-sectional transmission electron micrographs (XTEM) and selected area electron
diffraction (SAED) patterns of as-deposited TiN and (Ti0.91Si0.09)N1.04 films on WC-Co
substrates are presented in Error! Reference source not found. and Error! Reference
source not found.. The images of both films display a dense columnar structure, while
the SAED patterns suggest a more coarse grained substructure for TiN in comparison to
(Ti0.91Si0.09)N1.04. The difference in structure is expected, as Si acts as a grain refiner and
induces point defects. The SAED patterns also verify the random crystallographic
orientation for TiN as well as the <002> preferred out-of-plane orientation for x=0.09
observed in XRD. The high-resolution (HR) image in Error! Reference source not
found.(b) shows parts of two overlapping grains with corresponding moiré fringes. The
low-angle grain boundary is defined by dislocations and (semi) coherency strain. It is
difficult, however, to rule out the presence of any monolayer-thick amorphous SiNz
phase.
Error! Reference source not found. shows TEM images as well as SAED and fast
Fourier transform (FFT) patterns from the as-deposited (Ti0.81Si0.19)N1.13 film on WC-Co
substrate. The film exhibits a dense, extremely defect-rich, fine grained, feather-like
structure as can be seen in the cross-sectional image in Error! Reference source not
found.(a). The combination of the cross-sectional SAED and plan-view FFT in Error!
Reference source not found.(a,d) state an <002> fiber texture. The 002 dark-field
image in Error! Reference source not found.(b) reveals that individual feathers consist
11
of bundles of elongated nm-size crystallites or nanocolumns. A HRTEM cross-sectional
image of a typical bundle is shown in Error! Reference source not found.(c). It
illustrates that the crystallites have low-angle grain boundaries in-between themselves.
Error! Reference source not found.(d) depicts a plan-view image with an apparent
phase separation into a crystalline cubic and an amorphous phase according to the lack of
diffraction contrast in the latter, also after tilting of the sample. The cubic phase builds up
the bundles of nanocolumns seen in Error! Reference source not found.(b,c). An EDX
map across 20x20 nm2 with a pixel resolution of 0.33 x 0.33 nm2, not shown here,
indicates that the crystallites consist of Ti-rich cubic (Ti,Si)N crystallites and the
amorphous areas of Si-rich SiNz:Ti.
The higher magnification image in Error! Reference source not found.(e) reveals that,
the (Ti,Si)N crystallites in bundles are composed of ~2x2 nm subgrains (or basic
structural units), which are delineated by dislocations in the form of semi-coherent grain
boundaries. Consequently, individual subgrains are slightly tilted with respect to each
other, causing an in-plane trans-rotation with a radius of ~100 nm (as derived from the
crystal plane trace indicated by the white line in Error! Reference source not
found.(e)). We propose that the {002} lattice planes are rotating in a self-organized
manner due to the segregation of Si to the grain boundaries. The corresponding FFT in
Error! Reference source not found.(e) from an area as small as 12x12 nm2 verifies the
tilting between the subgrains. The lattice defect density assessed from counting
dislocations in the sub-grains is ~1x1014 cm-2, which corresponds to values found for
heavily cold work hardened material. The observed phase separation with a concomitant
grain refinement agrees with the XRD analysis in which a substantial peak broadening
12
was observed for x≥0.14, c.f., Error! Reference source not found.. On a larger scale,
the dark-field micrograph in Error! Reference source not found.(b) illustrates that the
(Ti,Si)N-crystallites and the amorphous phase grows in a vertical lamellar mode where
the grain width is determined by the rates of deposition and element partitioning on the
surface for the two phases. While a self-organized nano-columnar structure has also been
observed for B-superstoichiometric Ti-B thin films [23] and for Zr-Si-N films in [24], the
present observation includes the in-plane trans-rotation within the columns, which is
original.
Chemical bonding in the as-deposited films
X-ray photoelectron spectroscopy on the
as-deposited TiN,
(Ti0.91Si0.09)N1.04,
(Ti0.86Si0.14)N1.09, and (Ti0.81Si0.19)N1.13 films on WC-Co substrates was carried out in
order to investigate the chemical bonds of the solid solution and two-phase structures,
respectively. The XPS scans of Ti2p, N1s and Si2p core levels are presented in Error!
Reference source not found.. The seemingly high oxygen-related XPS signals can be
ascribed to surface contaminations as they disappear after clean sputtering, and hence, do
not contradict the overall low oxygen content determined by ERDA (compare Table I).
The Ti2p signal displays three distinct peaks, associated with chemically shifted Ti2p3/2
core levels. The high binding energy feature at 458.25 eV is attributed to TiO2 [25], the
other peaks at 456.7 eV and 455.2 eV correspond to Ti-O-N [26] and TiN [25],
respectively. The film with highest Si content exhibits shifts for the TiO2 and TiN
contributions with respect to the other films. These shifts can be tentatively assigned to
the change of chemical environment in the vicinity of TiO2 due to presence of Si in the
film. A change in the shape of the N1s line is observed depending on the Si content in the
13
sample. The Si contribution in the N1s signal can be observed at 398.0 eV and is assigned
SiNz [27]. The contribution from Ti-N is found at 397.3 eV [25,27]. The peaks at 396.0
eV and ~396.4 eV are attributed to surface contamination such as Ti-O-N [27], since they
disappear after Ar-sputtering. As observed from the normalized Si2p spectra, the
recorded signal exhibits a more or less retained shape at 101.8 eV independently of the Si
content in the sample. There is, however, a small peak broadening of ~0.2 eV for x=0.19
compared to x=0.09, which may be assigned to the formation of amorphous Si3N4:Ti
phase. Furthermore, no SiOx peaks were detected in this analysis, which indicates that Si
in the films is not subject to oxidation.
Since there is no XPS data available in literature from cubic Si-N phases, we measured
on a dc magnetron sputtered c-SiNx/TiN multilayer thin film from Ref. [8] where cSiNx(001) is epitaxially stabilized between TiN(001) layers. The layer thicknesses are 20
Å and 5 Å for TiN and SiNx, respectively, and are thus thin enough to provide XPS signal
from several layers. Whether the c-SiNx has octahedral or tetrahedral Si-N bonds is still
under discussion [8]
The N1s and Si2p spectra from the c-SiNx/TiN multilayer film are included in Error!
Reference source not found.. The N1s line has very similar appearance in comparison to
the x=0.09 film, with corresponding peaks assigned to TiN, SiNx as well as Ti-O-N
surface contaminations. The Si2p peak is observed at 102.0 eV, which corresponds to a
shift of only ~0.2 eV compared to the binding energy of 101.8 eV for the as-deposited
x=0.09 film. The reason for the shift may be the absence of Ti as second nearest
neighbor.
14
Our interpretation of the XPS result is that the Si-N bonds have tetrahedral coordination,
but the presence of octahedral coordination can, however, not be ruled out, since there are
no reference values in literature. In the context whether Si is located in the amorphous or
crystalline phase in the samples, XPS does not provide unambiguous evidence since the
observed shifts are in the order of the energy resolution of the instrument. Hence, other
techniques with higher energy resolution are suggested to reveal the bond nature of SiNz.
As a final remark, we note that Ar-sputtering of the films induces peak shift for the Si2p
peak, this was not taken into account in our previous paper where the shift of the Si2p
peak position to 100.9 eV was most likely sputter-induced [9].
To conclude the analysis of the as-deposited (Ti1-xSix)Ny films grown by arc-evaporation
onto WC-Co substrates we note that when x increases from 0 to 0.20, also y increases
from 0.98 to 1.13. Films with 0.01≤x≤0.09 exhibit a metastable cubic solid solution phase
with a lattice parameter of 4.24 Å, where Si acts as a grain refiner. Higher Si contents
initiates a phase separation into metastable c-(Ti1-xSix)Ny nanocrystallites and an
amorphous SiNz:Ti. All films are under compressive residual stress, which increases from
σ=-2.7 GPa to σ=-5.5 GPa with increasing Si. XPS suggests the presence of tetrahedral
binding configuration between Si-N for at least Si rich films. The location of this Si to the
amorphous or cubic crystalline Ti-Si-N phases (or both), however, has not been
elucidated this far.
In the next section we will consider the thermal stability of the (Ti1-xSix)Ny films on WCCo after isothermal annealing at 1000 °C for 2 h. This is followed by a study on
15
thermally more stable sintered c-BN and Ta substrates annealed at 1000, 1100, and 1200
°C.
Composition and microstructure of annealed (Ti1-xSix)Ny films on WC-Co
In the discussion below, the films are labeled according to the as-deposited composition.
To study the influence of annealing on the film composition, the nine films deposited on
WC-Co were analyzed by ERDA. The corresponding compositions after isothermal
annealing for 2 h at 1000 °C are presented in Table II. The solid-solution films with
x≤0.09 exhibit no Si loss, while the N content decreases slightly. On the contrary, the
two-phase films with x≥0.14 exhibit an increasing loss of both Si and N. It is noteworthy
that regardless of the initial Si content within these films, the retained Si content after
annealing is ~5 at.% (x≈0.10). This value, in fact corresponds to the effective solid
solution limit found for the as-deposited films. Furthermore, besides the Si and N
outward diffusion, we also observe a Co and W release from the substrate and diffusion
into the films. Each of these effects which also increases with increasing Si content of the
as-deposited films. In contrast, the O impurities are more or less retained for all films,
whereas the C impurities increase minutely.
Error! Reference source not found. shows the x-ray diffractograms from the
corresponding post-annealed samples. The 1000 °C heat treatment did not induce any
phase transformation as the diffractograms appear similar to those from the as-deposited
films, (c.f., Error! Reference source not found.). All films have thus also preserved
their texture, i.e., films with x≤0.04 exhibit a preserved random crystallographic
16
orientation, while films with x≥0.09 maintain their strong <002> out-of-plane preferred
orientation. Moreover, the FWHM decreases slightly for x≤0.09 films and more
prominent for x≥0.14.
Results for the film residual stress states after annealing are presented in Error!
Reference source not found.(b). The stress in the TiN film has decreased to around zero
due to defect annihilation and grain coarsening, which are known to operate at an
annealing temperature of 1000 °C [28]. Films with Si content of 0.01≤x≤0.09, exhibit a
reduced, but still relatively high compressive stress of between σ=-0.9 GPa for x=0.01
and σ=-2.5 GPa for x=0.09. Structurally, this may be correlated to migration of ionbombardment induced Si defects in the NaCl lattice and the formation of an x-ray
amorphous SiNz grain boundary phase. All films with higher Si contents are virtually
fully relaxed, but the relaxation mechanisms is complicated by the superimposed Si and
N out-diffusion and Co and W interdiffusion presented above.
For a more detailed analysis on the relaxation processes we consider the XRD data
presented in Error! Reference source not found.(a,b), and the FWHMs and positions of
the 002 peak depending on the as-deposited Si content x. We start with the single-phase
(x≤0.09) films in Error! Reference source not found.(a). TiN exhibits a small peak shift
of to higher angle of ~0.1° after the annealing. This can be attributed to both the decrease
of the N-content from TiN0.99 to TiN0.86 and the stress reduction. An N-substoichiometry
for TiN effectively reduces the lattice parameter [29] and also a decreased in compressive
stress will reduce the out-of-plane lattice parameter, and hence yield a higher 2θ-angle.
Furthermore, the FWHM has decreased from ~0.55° to ~0.40°, which corroborates the
17
occurrence of point defect annihilation and grain coarsening as introduced above. The
same applies to films with 0.01≤x≤0.09, but to a lower extent, correspondingly the asdeposited peak positions are retained after annealing, and the FWHM is decreased only
slightly, less than 0.05°. On the other hand, all films with x≥0.14 exhibit a substantial
peak shift after annealing to higher angles, close to the position of annealed TiN at 42.6°
as a consequence of the virtually complete stress relaxation. Also, the FWHM decreases
considerably from ~1.2-1.8° to ~0.7-0.8°, which again is due to grain coarsening,
intensified by the W and Co interdiffusion.
To further characterize the W and Co interdiffusion, we studied the elemental distribution
in the annealed films by analytical STEM as displayed in Error! Reference source not
found.. The STEM image of the 1000 °C annealed (Ti0.81Si0.19)N1.13 film in Error!
Reference source not found.(a) reveals a film recrystallization into a cellular structure,
with – according to the Z contrast – heavier elements gathered at the grain boundaries.
Also the EDX maps in Error! Reference source not found.(b-f) state that Ti and N are
accumulated within the grains, while W and Co are mainly situated at the grain
boundaries. Contrary, the low Si content in combination with the overlap of Si Kα and W
Mα lines does not provide evidence whether Si is located within the grains and/or at the
grain boundary, see Error! Reference source not found.(c).
Chemical bonding of annealed (Ti1-xSix)Ny films
XPS on annealed (Ti0.81Si0.19N)1.13 at 1000 °C confirmed the presence of W and Co in the
film. Scans over the Si2p peak, however, revealed a shift from 101.8 eV to ~102.3 eV.
We attribute the peak shift to the formation of Si-O-N [30], due concurrent segregation of
18
Si and O to the boundary. Any Ti2Si formation due to the out-diffusion of N has binding
energy at ~98 eV and can be excluded for our films. For Co-Si, which would be expected
at the grain boundaries, a peak should be present at ~101 eV [31] both for Si2p and Co3s
bindings. No Co-Si signal was, however, resolved.
Mechanical properties (Ti1-xSix)Ny films on WC-Co
The hardness of the as-deposited and 1000 °C annealed films on WC-Co substrates are
presented in Error! Reference source not found.. The as-deposited TiN film exhibits a
hardness of 29.8±1.5 GPa. The addition of Si to TiN induces a hardness increase to a
maximum of 41.8±2.0 GPa for a Si content close to x=0.14. Further increase in the Si
content results in a small hardness reduction. The hardness of the as-deposited TiN film
is in the same order as for other arc-evaporated TiN [9,32], but significantly higher than
monolithic TiN(001) of 20 GPa [33]. The difference to the single-crystal material is
attributed to defect hardening from the high density of lattice point defects common to
arc evaporation as discussed in [35]. This is also expressed here in the compressive stress
state of the films. The hardness enhancement for the as-deposited films with 0.01≤x≤0.09
can be explained by a combination of several hardening mechanisms as solid solution,
grain size, point defects (residual stresses). Solid solution hardening originating from
atomic and elastic modulus mismatch is explained in the Fleischer model [34]. A smaller
grain size induces hardening according to the Hall-Petch relation [35,36] and point
defects impede dislocation movements by inducing local lattice distortions. Any
thermally induce strain between WC-Co and the film (TiN) would yield tensile stress in
the order or ~1 GPa for a deposition temperature of 500 °C [37]. The film with x=0.14
exhibits all above hardening mechanisms as well as coherency strain hardening from the
19
phase separation into a two-phase structure. The films with Si-content x>0.14 possess
slightly lower hardness than 41.8 GPa because of weakening of the grain boundaries due
to thickening of the SiNz:Ti phase as the Si content increase. Formation and thickening of
the a-SiNz phase above a few monolayers is known to decrease the hardness due to
amorphization and lost coherency as seen in for instance [4] for nanocomposites and [6,7]
for nanolaminates.
After annealing at 1000 °C for 2 h the hardness for TiN decreases to 23.5±1.5 GPa. This
decrease is commonly reported in literature, e.g., [28], and consistent with the XRD
diffractograms and stress results presented above, and again, is an effect of defect
annihilation and stress relaxation. The films with Si content 0.04≤x≤0.09, however, retain
the hardness at 30.2±1.6 GPa and 38.8±2.0 GPa, respectively. This implies that the SiNz
phase segregated to the grain boundary constitutes a strong coherent interface with the
grains. Furthermore, the hardness drops to <28 GPa for films with x≥0.14 (in the asdeposited state),. This is due to recrystallization within in the cubic state due to Si and N
out-diffusion as well as Co and W interdiffusion from the substrate as observed by EDXSTEM and ERDA. It can be assumed that these elements possibly together with O form
tissue phases, including CoSi2, TiO2 or SiO2, which together with the decrease in residual
stress weaken the structure.
Phase stability of the films on thermally more stable substrates
In order to clarify whether the observed Si and N losses at temperatures above 900 °C for
films with an as-deposited Si content of x≥0.14 are inherent, or induced by the WC-Co
interdiffusion, more temperature stable substrates have to be employed. Furthermore,
20
thermally more stable substrates are required also for the high temperatures needed to
study any additional phase transformations for the solid solution films with 0.01≤x≤0.09.
We use both sintered polycrystalline c-BN, a commonly used substrate in metal cutting
applications, and Ta plates.
Table III presents the compositions measured by ERDA of the as-deposited films on cBN substrates and their composition after isothermal annealing at 1000 °C, 1100 °C, and
1200 °C. The as-deposited films on c-BN substrate have a (Ti0.83Si0.17)N0.9 composition.
The films on c-BN substrate were stable up to 1000 °C. At 1100 °C, however, Si and N
out-diffusion and desorption took place, which changed the composition to
(Ti0.85Si0.15)N0.85. When annealed at 1200 °C, the film content changed drastically with an
up-take of 11.4 at.% C. The source for this C is the binding phase of TiC in the substrate.
Hence, the c-BN substrates offered effective stability of 1100 °C, which is up to 200 °C
more than for the WC-Co substrate, but the interdiffusion of C hinders the assessment on
the inherent stability of the film.
The two films investigated on Ta substrate have the compositions of (Ti0.91Si0.09)N1.04 and
(Ti0.81Si0.19)N1.13 as measured by ERDA (see Table III) to. The Si content in the
(Ti0.91Si0.09)N1.04 film on Ta substrate is sustained up to 1000 °C. At 1100 °C, however,
local Si and N out-diffusion via the grain boundaries and eventual desorption is initiated,
which cause a composition changes to (Ti0.92Si0.08)N. This process of Si and N is
enhanced at 1200 °C and the remaining film consists of (Ti0.98Si0.02)N.
21
The (Ti0.81Si0.19)N1.13 film on Ta substrate is compositionally stable up to 1000 °C, at
1100 °C Si and N out-diffusion occurs resulting in a, (Ti0.??Si0.??)Ny composition. At 1200
°C the Si depletion is more or less complete and only ?? at.% remain.
Error! Reference source not found. displays a cross-sectional TEM from the asdeposited (Ti0.83Si0.17)N0.9 film on c-BN substrate, which exhibits a dense, very defectrich segregated two-phase structure. Error! Reference source not found.(b) shows the
XTEM image from the area close to the free surface of the same film annealed at 1100 °C
for 2 h for which no substrate interdiffusion was observed in ERDA. The film volume
closest to the free surface (see Error! Reference source not found.(b)), however, has
locally recrystallized into ~10-50 nm grains that are elongated along the substrate normal.
XTEM analysis of the 1200 °C-film on c-BN substrate (not shown) reveals that the film
has recrystallized throughout its thickness to grain sizes of >100 nm. Moreover, the
grains are facetted, typical for a relaxed, annealed structure.
Error! Reference source not found. (observera att figurordningen mellan 12 och 11(ce) ska ändras) shows a STEM image together with EDX elemental maps of the
(Ti0.83Si0.17)N0.9 film on c-BN substrate annealed at 1100 °C. The Si depleted grains
appear with a dark contrast in the STEM micrograph in Error! Reference source not
found.(a). The maps display that the structure has recrystallized due to Si out-diffusion.
At this annealing temperature also segregation of Si within the film is present into Si-rich
precipitate as shown in Error! Reference source not found.(b-d).
22
Error! Reference source not found.(c) displays the as-deposited (Ti0.81Si0.19)N1.13 film
on Ta substrate. It reveals a dense, very defect-rich segregated two-phase feather-like
structure, similar to the (Ti0.81Si0.19)N1.13 film deposited on WC-Co, c.f., Error!
Reference source not found.(a). Error! Reference source not found.(d and e) shows
bright-field (d) and dark-field (e) XTEM images from the (Ti0.81Si0.19)N1.13 film on Ta
substrate annealed at 1200 °C. The microstructure has recrystallized from the feather-like
two-phase structure into a nanocrystalline structure with nm-sized polyhedral grains
elongated in the film growth direction. The film consists essentially of TiN any O
Manfred?? as revealed by ERDA. The bright contrast at the grain boundaries in (d)
indicates the presence of a porous structure, which is a consequence from the outdiffusion of Si and N. As a result, Ti and Si have diffused into the Ta substrate and
formed a 1-µm-thick Ta-Si layer as revealed by EDX-STEM. Within this layer, Ti has
together with O from the substrate constitute TiO2 precipitates, which appear as bright
regions in the Ta-Si layer in Error! Reference source not found.(d). In-between the TaSi layer and the film, there is a ~100 nm broad recrystallized TiN band. The film
thickness has decreased from 2.0 µm to 1.5 µm during annealing at 1200 °C because of
the diffusion of Ti, Si, and N from the film into the substrate.
The hardness of the films on c-BN substrates measured with nanoindentation is shown in
Error! Reference source not found.. As-deposited films had hardness a of 39.5±2.1
GPa. In difference to the film annealed on WC-Co substrate, the hardness was retained at
1000 °C. At 1100 °C the hardness was slightly reduced to 38.6±2.3 GPa, as a
consequence of the Si depletion confined mainly close to the film surface. At 1200 °C the
23
hardness has dropped to 15.8±1.5 GPa because of a fully recrystallized, porous structure
strongly influenced by interdiffusion of mainly C from the TiC binder phase.
Conclusions
(Ti1-xSix)Ny films with 0≤x≤0.20 were deposited by arc-evaporation onto WC-Co
substrates. As-deposited films with 0.01≤x≤0.09 exhibit a dense, defect-rich,
polycrystalline, metastable (Ti1-xSix)Ny solid solution with a lattice parameter of ~4.24 Å.
For x≥0.14, the films attain a <002> fiber-textured two-phase structure, which consists of
defect-rich Ti-rich c-(Ti,Si)N crystallites and Si-rich a-SiNz:Ti. The N-content increased
close to linear with increasing Si content from y=0.99 to y=1.13. The nanoindentation
hardness increased from 30 GPa for TiN to 42 GPa for x=0.14. For higher Si contents,
the hardness was reduced slightly to 38 GPa due to formation amorphous SiNz:Ti phase.
We propose that Si exists in the films in four different forms:
A) Si is substituted for Ti in the NaCl-type lattice in octahedral coordination.
B) Si point defects of induced by the metal ion-bombardment generates tetrahedrally
coordinated SiNz defect clusters.
C) Crystalline SiNz grain boundary phase which forms a coherent interface with (Ti,Si)N
grains.
24
D) Amorphous SiNz grain boundary phase, which is semi or incoherently bonded to the
environment.
XPS analysis could, however, not reveal any significant distinction between the different
forms of Si coordinations. As a reference investigation, TiN/SiNx superlattice with cSiNx epitaxially stabilized between TiN(001) layers did not reveal any clear peak shifts
of the Si2p peak compared to the as-deposited films.
Isothermal annealing experiments of the solid solution films with x≤0.09 in the asdeposited state result in a defect annihilation process and an inherent thermodynamically
driven segregation of type B Si to the grain boundaries to form a coherent x-ray
amorphous SiNz grain boundary phase (type C) with thickness of a few monolayers. The
SiNz phase effectively inhibits diffusion across grains, which limits further
recrystallization as well as stress relaxation. As additional Si reach the grain boundary,
the SiNz phase thickens and eventually amorphize (type D). At ~1100 °C, out-diffusion
of amorphous Si and N is initiated.
Similar annealing experiments were performed on the two-phase films with x≥0.14 and
result in recrystallization into polyhedral grains and increased Si-N confinement to the
grain boundaries. These films are compositionally stable up to 1000 °C, before local Si
depletion occurs which consequently result in residual porous grain boundary structures.
Finally, nanoindentation experiments reveal that the hardness of 38 GPa for the x=0.09
film on WC-Co substrates is retained at 1000 °C as a result of the recrystallization and
25
formation of the strong coherent interface phase of SiNz. For x≥0.14, W and Co
interdiffusion decrease the hardness to <28 GPa, due to formations of weaker grain
boundary phases. The nanoindentation hardness of (Ti0.83Si0.17)N0.9 on sintered c-BN
substrates was despite the local Si and N out-diffusion at 1100 °C almost retained at ~3840 GPa.
Acknowledgements
The Swedish Research Council (VR) and the Swedish Foundation for Strategic Research
(SSF) MS2E program are acknowledged for financial support. Jens Jensen (Tandem
Laboratory), Urban Wiklund, and Mattias Lindquist at Uppsala University are
acknowledged for assistance with ERDA measurements and nanoindentation,
respectively. Mats P. Johansson at Seco Tools AB is acknowledged for the residual stress
measurements and preparation of one the FIB TEM specimens. Björn Alling at
Linköping University is acknowledged for fruitful theoretical discussion. Hans Söderberg
at Swerea Kimab AB (earlier at Luleå University of Technology) is acknowledged for
providing the c-TiN/SiNx multilayer sample.
References
26
Table I. Elemental composition of the as-deposited films on WC-Co substrates as
measured with ERDA. The values are obtained from averaging over the measured depth
profiles (~100-300 nm), and have been normalized to 100 at.%.
Before annealing
Sample
Ti
Si
N
O
C
Others
(at.%) (at.%) (at.%) (at.%)
(at.%)
(at.%)
1
49.34
0.12
49.20
0.45
0.77
0.12
2
49.33
0.64
48.86
0.56
0.51
0.1
3
47.07
1.86
49.96
0.60
0.38
0.13
4
44.12
4.20
50.69
0.43
0.42
0.14
5
40.84
6.52
51.42
0.59
0.51
0.12
6
39.36
7.33
52.08
0.65
0.44
0.14
7
38.46
8.25
52.08
0.70
0.41
0.1
8
37.33
9.14
52.26
0.76
0.41
0.1
9
37.29
9.02
52.26
0.81
0.52
0.1
Composition
TiN0.99
(Ti0.99Si0.01)N0.98
(Ti0.96Si0.04)N1.02
(Ti0.91Si0.09)N1.04
(Ti0.86Si0.14)N1.09
(Ti0.84Si0.16)N1.12
(Ti0.92Si0.18)N1.11
(Ti0.80Si0.20)N1.12
(Ti0.81Si0.19)N1.13
Table II. Elemental composition of the post annealed films at 1000 °C on WC-Co
substrates as measured with ERDA. The values are obtained from averaging over the
measured depth profiles (~100-300 nm), and have been normalized to 100 at.%.
After annealing
Sample
Ti
Si
N
O
C
Co
W
Composition
(at.%) (at.%) (at.%) (at.%)
(at.%)
(at.%) (at.%)
1
52.84
0.17
45.55
0.46
0.97
0.00
0.00
TiN0.86
2
51.68
0.58
46.60
0.45
0.68
0.00
0.00
(Ti0.99Si0.01)N0.89
3
48.11
1.80
47.80
0.77
1.23
0.00
0.00
(Ti0.96Si0.04)N0.96
4
45.27
4.10
48.26
0.95
1.20
0.00
0.00
(Ti0.92Si0.08)N0.98
5
43.22
4.65
48.23
0.66
0.83
2.06
0.34
(Ti0.90Si0.10)N:(Co,W)
6
42.56
5.07
45.34
0.87
0.42
5.18
0.55
(Ti0.89Si0.11)N:(Co,W)
7
40.04
4.62
44.61
0.87
1.33
7.75
0.77
(Ti0.90Si0.10)N:(Co,W)
8
38.95
5.39
44.08
0.98
0.95
8.86
0.67
(Ti0.88Si0.12)N:(Co,W)
9
42.77
3.46
43.32
0.68
1.08
6.81
1.78
(Ti0.93Si0.07)N:(Co,W)
Table III. Elemental composition of the as-deposited and post annealed films at 1000 °C,
1100 °C, and 1200 °C on c-BN and Ta substrates, respectively, as measured with ERDA.
27
The values are obtained from averaging over the measured depth profiles (~100-300 nm),
and have been normalized to 100 at.%.
c-BN
substr.
as-dep.
1000°C
1100°C
1200°C
Ta
substr.
as-dep.
1000°C
1100°C
1200°C
Ta
substr.
as-dep.
1000°C
1100°C 1*
1100°C 2*
1200°C
Ti
(at.%)
43.5
43.2
45.4
51.1
Ti
(at.%)
43.1
42.6
42.9
48.3
Ti
(at.%)
36.9
37.2
42.8
39.8
45.9
Si
(at.%)
8.6
9.0
7.9
3.2
Si
(at.%)
5.0
4.9
3.2
1.8
Si
(at.%)
9.3
9.2
1.4
7.8
1.2
N
(at.%)
47.3
47.0
45.5
34.0
N
(at.%)
50.5
49.9
46.2
44.1
N
(at.%)
52.3
52.0
41.1
50.4
37.7
C
(at.%)
0.6
0.7
1.4
11.4
C
(at.%)
0.5
1.8
6.5
4.1
C
(at.%)
0.4
0.7
11.6
1.8
12.8
O
(at.%)
O
(at.%)
0.8
0.8
1.0
1.6
O
(at.%)
1.0
0.9
3.0
0.2
2.4
others
(at.%)
others
(at.%)
0.1
0.0
0.2
0.1
others
(at.%)
0.1
0.0
0.1
0.0
0.0
Composition
(Ti0.83Si0.17)N0.90
(Ti0.83Si0.17)N0.90
(Ti0.85Si0.15)N0.85
(Ti0.94Si0.06)N0.63
Composition
(Ti0.90Si0.10)N1.05
(Ti0.90Si0.10)N1.05
(Ti0.93Si0.07)N1.00
(Ti0.96Si0.04)N0.88
Composition
(Ti0.80Si0.20)N1.13
(Ti0.80Si0.20)N1.12
(Ti0.98Si0.02)N0.93
(Ti0.84Si0.16)N1.06
(Ti0.97Si0.03)N0.80
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