ArticlePDF Available

Correlation between fracture characteristics and valence electron concentration of sputtered Hf-C-N based thin films

Authors:

Abstract and Figures

Hard protective coating materials based on transition metal nitrides and carbides typically suffer from limited fracture tolerance. To further tune these properties non-metal alloying – substituting C with N – has been proven favorable for magnetron sputtered Hf-C-N based thin films. A theoretically predicted increase in valence electron concentration (from 8.0 to 9.0 e/f.u. from HfC to HfN) through nitrogen alloying lead to an increase in fracture toughness (KIC obtained during in-situ SEM cantilever bending) from 1.89 ± 0.15 to 2.33 ± 0.18 MPa·m1/2 for Hf0.43C0.57 to Hf0.35C0.30N0.35, respectively. The hardness remains close to the super-hard regime with values of 37.8 ± 2.1 to 39.9 ± 2.7 GPa for these specific compositions. Already the addition of small amounts of nitrogen, while sputtering a ceramic HfC target, leads to a drastic increase of nitrogen on the non-metallic sublattice for fcc single phased structured HfC1-xNx films, where x = N/(C + N). The here obtained results also provide experimental proof for the correlation between fracture characteristics and valence electron concentration.
Content may be subject to copyright.
Contents lists available at ScienceDirect
Surface & Coatings Technology
journal homepage: www.elsevier.com/locate/surfcoat
Correlation between fracture characteristics and valence electron
concentration of sputtered Hf-C-N based thin films
T. Glechner
a,
, S. Lang
a
, R. Hahn
a
, M. Alfreider
b
, V. Moraes
c
, D. Primetzhofer
d
, J. Ramm
e
,
S. Kolozsvári
f
, D. Kiener
b
, H. Riedl
a,c
a
Christian Doppler Laboratory for Surface Engineering of high-performance Components, TU Wien, Austria
b
Department of Materials Science, Montanuniversität Leoben, Austria
c
Institute of Materials Science and Technology, TU Wien, A-1060 Wien, Austria
d
Department of Physics and Astronomy, Uppsala University, SE-75120 Uppsala, Sweden
e
Oerlikon Balzers, Oerlikon Surface Solutions AG, 9496 Balzers, Liechtenstein
f
Plansee Composite Materials GmbH, D-86983 Lechbruck am See, Germany
ARTICLE INFO
Keywords:
Hf-C-N
Fracture resistance
Non-metal alloying
Valence electron concentration
Thermal stability
ABSTRACT
Hard protective coating materials based on transition metal nitrides and carbides typically suffer from limited
fracture tolerance. To further tune these properties non-metal alloying – substituting C with N – has been proven
favorable for magnetron sputtered Hf-C-N based thin films. A theoretically predicted increase in valence electron
concentration (from 8.0 to 9.0 e/f.u. from HfeC to HfeN) through nitrogen alloying lead to an increase in
fracture toughness (K
IC
obtained during in-situ SEM cantilever bending) from 1.89 ± 0.15 to
2.33 ± 0.18 MPa·m
1/2
for Hf
0.43
C
0.57
to Hf
0.35
C
0.30
N
0.35
, respectively. The hardness remains close to the super-
hard regime with values of 37.8 ± 2.1 to 39.9 ± 2.7 GPa for these specific compositions. Already the addition
of small amounts of nitrogen, while sputtering a ceramic HfeC target, leads to a drastic increase of nitrogen on
the non-metallic sublattice for fcc single phased structured HfC
1-x
N
x
films, where x = N/(C + N). The here
obtained results also provide experimental proof for the correlation between fracture characteristics and valence
electron concentration.
1. Introduction
Various challenges in the field of environmental sustainability as
well as emission reduction in general are closely linked to the usage of
ultra-stable materials. Especially, protective coatings of high-perfor-
mance components – such as blades or drive train systems in jet engines
or steam turbines – depict a key role to achieve further milestones
concerning efficiency and operating ranges [1,2]. For these extremely
harsh environments, involving highest temperatures accompanied by
abrasive and corrosive media, transition metal nitrides (TMN), and
carbides (TMC) are highly interesting coating materials due to their
excellent thermomechanical properties and chemical inertness [2].
Nevertheless, these ceramic compounds typically lack ductility com-
pared to common metals and metallic alloys [3–5]. Also, the formation
of continuous and dense oxide scales at the highest temperatures is a
strong limitation for wide usage [7,6].
Among diverse concepts [8], substitutional alloying of the non-
metallic sublattice by exchanging C with N atoms – forming
carbonitrides – depict an interesting approach to intrinsically tune these
coating properties. Through this substitutional exchange, a distinct
adaption of the prevalent bonding character and hence valence electron
concentration (VEC) is accomplished [9]. This fundamental approach
was initially proven for face-centered cubic (fcc) TiC
1-x
N
x
adjusting the
VEC to 8.4 e/f.u. exhibiting a hardness maximum due to the population
of a particular s band, located between the non-metal p and metal d
orbitals, being strongly resistive against shape changes and shearing
[10]. While for the hardness a distinct correlation with an ideal VEC for
various systems can be made, more or less positive behavior is sug-
gested for the toughness by increasing VEC up to 10 [11–15]. In our
previous studies, we could show, that the exchange of about 36% C
with N atoms in fcc-structured TaC
y
(VEC increasing from 8.24 to 9.36
e/f.u. compared to the binary TaC
0.81
) lead to a fracture toughness
increase from 1.8 to 2.9 MPa·m
-1/2
(while keeping the thin film mor-
phology similar, especially regarding their columnar growth), respec-
tively [4,16]. As already shown in previous studies [17–19], these
transition metal carbides and nitrides are strongly effected by structural
https://doi.org/10.1016/j.surfcoat.2020.126212
Received 11 March 2020; Received in revised form 2 July 2020; Accepted 16 July 2020
Corresponding author.
E-mail address: thomas.glechner@tuwien.ac.at (T. Glechner).
Surface & Coatings Technology 399 (2020) 126212
Available online 19 July 2020
0257-8972/ © 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY license
(http://creativecommons.org/licenses/BY/4.0/).
T
defects, such as vacancies, influencing similarly the bonding states.
However, a distinct population of structural defects is for sure more
challenging compared to non-metal alloying, especially with respect to
the deposition process [17]. Furthermore, different theoretical studies
using ab initio molecular dynamics [20,21] also predict an increase of
the melting temperature of about 200 K for Hf
0.53
C
0.27
N
0.20
compared
to pure to HfeC – suggesting a new highest melting material. Next to
these fundamental aspects while exchanging C with N in transition
metal carbides, the selection of a proper transition metal is also highly
important. Based on their extremely high phase stability and the above-
mentioned facts, TaC
y
and HfC
y
are promising candidates [22]. Espe-
cially, the Hafnium based system is encouraging as one oxide is highly
stable – in its high-temperature modification (HfO
2
, fcc, #SG 225) even
up to 3083 K [23,24]. It can be synthesized by reactive cathodic arc
evaporation in an Al-Hf-O material system, but it is difficult to stabilize.
Studies on Hf-containing thin films [25] and bulk materials [26] sug-
gested high oxidation resistance and hence great potential (even above
1600 °C). However, only a few studies are available for Hf-C-N coatings
in general, but also for sputtered deposited coatings. These papers
mainly concentrate on the impact of deposition parameters (tempera-
ture, bias) on structure, chemical composition as well as tribological
performance [27,28].
Therefore, within this study, we combine theoretical Density
Functional Theory (DFT) calculations with an experimental validation
by magnetron sputtered HfeC and Hf-C-N thin films. Thereby, we
wanted to gain an in-depth understanding of the impact of the non-
metallic sublattice occupation (varying nitrogen content) on the ther-
momechanical properties, especially for fracture tolerance.
2. Experimental procedures
We used two laboratory-scaled magnetron sputtering systems
(modified Leybold Heraeus Z400 and AJA Orion 5) having different
target to substrate arrangements – a parallel face to face configuration
obtaining a target to substrate distance of 50 mm as well as a tilted set-
up with a normalized distance of 70 mm – to deposit HfeC and Hf-C-N
thin films. Both deposition systems were equipped with a 3-in. HfeC
compound target (99.6% purity, Plansee Composite Materials GmbH)
operated in DC-pulsed mode (pulse frequency 150 kHz, pulse width
2576 ns). The HfeC coatings were deposited in pure Argon atmosphere
(99.999% purity), while Hf-C-N thin films were sputtered reactively in
an Argon/Nitrogen gas mixture. All depositions were carried out at a
total pressure of 0.4 Pa, as well as a bias potential of −10 V. A base
pressure below 10
−4
Pa was ensured for all deposition runs. We want to
emphasize that all coatings were deposited using the same target.
In the tilted setup, a total gas flow of 20 sccm has been used. To
obtain different nitrogen concentrations within the Hf-C-N coatings, we
varied the amount of nitrogen in the total gas flow (f
[N2]norm
=
+
f
f f
N
N Ar
2
2
)
from 0.0 to 0.15, respectively. The heater temperature was fixed at
600 °C, which corresponds to T
sub
= 415 ± 15 °C at the substrate
surface. The confocally arranged cathode (tilted by 25° to the rotating
substrate holder) was powered by an MKS RPG-50 Pulsed DC plasma
generator.
For the parallel configuration (bottom to top set-up), all depositions
were carried out at a total gas flow of 34.1 sccm. The non-metal sub-
lattice occupation was again varied through the nitrogen to the total gas
flow rate from 0.0 to 0.50, respectively. In contrast to the tilted setup,
the working gas pressure was not fixed through a pressure valve but
adjusted by the gas flow and pumping speed. Therefore, it decreased
from 0.393 Pa in pure Argon to 0.376 Pa for f
[N2]norm
= 0.5. The heater
temperature was varied between 300 °C and 500 °C, which corresponds
to 250 ± 15 °C and 380 ± 15 °C at the substrate surface, respectively.
Based on the physical limits of the flow controller, our lowest ni-
trogen to total gas flow ratio (f
[N2]norm
) was 0.05 for both systems. All
thin films were deposited on single-crystalline Al
2
O
3
platelets (0001-
oriented, 10x10x0.53 mm
3
), single-crystalline Si stripes (100-oriented,
20x7x0.38 mm
3
), and polished austenitic stainless-steel platelets
(20x7x0.8 mm
3
). Selected coatings were also grown on 0.05 mm thick
steel foils, which were dissolved afterward in 20% hydrochloric acid
[22]. The residual Hf
1-y
C
y
and Hf
1-y-z
C
y
N
z
coating flakes were grounded
mechanically to fine-grained powders, which was used for substrate-
interference free X-Ray Diffraction (XRD) analysis. Preliminary, to the
coating process, the target, and all substrates were sputter-cleaned in a
pure Argon atmosphere at a total pressure of 7.0 Pa in the tilted and
1.3 Pa in the face-to-face set up, respectively.
Time-of-flight Elastic Recoil Detection Analysis (ERDA) using a
36 MeV I8+ primary ion beam and detecting recoils at an detection
angle of 45° by a segmented gas ionization chamber [29] was per-
formed at the tandem laboratory at Uppsala University to determine the
chemical composition of the coatings. Data evaluation is described in
more detail in [22]. The structure of all coatings has been analyzed by
XRD in Bragg Brentano configuration using a Panalytical Empyrean
diffractometer – equipped with a Cu-K
α
radiation source operating at
45 kV and 40 mA (wavelength λ = 1.54 Å). Scanning Electron Mi-
croscopy (SEM, FEI Quanta 250 FEGSEM operated at 15 kV) was used
to scale the coating thicknesses and survey the coating quality in gernal.
More detailed analysis on the film morphology was done by Trans-
mission Electron Microscopy (TEM FEI TECNAI G20, an acceleration
voltage of 200 kV) – only for selected coatings.
Thermal treatments on sapphire substrates were done in a high-
temperature vacuum furnace (Centorr Vacuum Industries Series
LF22–2000) up to 1200 °C. The heating rate was 20 K·min
−1
followed
by an isothermal annealing time of 10 min and subsequent passively
cooling.
Mechanical properties such as Young's modulus and hardness were
examined on sapphire substrates by nanoindentation using an Ultra-
Micro-Indentation System (UMIS) equipped with a Berkovich diamond
tip. For every sample measurement, at least 25 indents with different
loads, 3 to 45 mN, were performed and their load-displacement curves
were analyzed after Oliver and Pharr [30]. The load range was chosen
to identify the ideal range where we have no influence of the substrate
and the surface roughness on our obtained values. It should be noted
that the indentation depths were below 10% of the film thickness,
which is reported to be sufficient to extinguish substrate interference
[31]. The Young's modulus was calculated using the Poisson's ratio for
HfC
1-x
N
x
obtained by the DFT calculations. In addition, we did not see
any clear dependency of the hardness and moduli on the load. All the
results presented in this study are averages of hardness and modulus
data obtained using loads from 3 to 45 mN (step size varied gradullay
from 2 to 0.5 to smaller loads). Residual stresses within the HfeC and
Hf-C-N coatings were investigated by using the modified Stoney Eq.
[32]. The curvature of the coated silicon substrates was analyzed by an
optical profilometer (Nanovea PS50).
All micro-mechanical tests were done on free-standing cantilevers,
which have been prepared by FIB machining. Before FIB-milling (using
an FEI Quanta 200 3D DualBeam-FIB), the silicon substrates were
etched using an aqueous 40 wt% KOH solution at 70 °C to obtain free-
standing film material. After that, the geometry was machined using a
beam current of 1.0 nA for coarse milling and 0.5 nA for fine milling at
an acceleration voltage of 30 keV. The initial notch was milled with 50
pA. The micromechanical experiments were conducted inside an LEO-
1540 FE-SEM/FIB dual-beam workstation with a PI 85 SEM pi-
coindenter mounted. The indenter was equipped with a diamond wedge
tip. All experiments were performed in displacement-controlled mode
(5 nm∙s
−1
). The calculation of the fracture toughness was then per-
formed following guidelines given by [33–35]. A load-deflection curve
and an image of the fractured cantilever can be found in the supple-
mentary material.
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
2
3. Calculational methodology
Density Functional Theory (DFT) calculations were performed with
the VASP code (Vienna Ab Initio Simulation Package) [36], using the
projector-augmented plane-wave (PAW) pseudopotentials [37]. To gain
a deeper insight into the effect of a chosen exchange-correlation (xc)
potential, we compared two different standard approximations namely,
Local Density Approximation (LDA) [38] and Perdew-Burke-Ernzerhof
Generalized Gradient Approximation (GGA-PBE) [39]. We used the fcc,
Fm
3
m, #SG 225 crystal structure creating a 2 × 2 × 2 supercell with
64 atoms from the conventional 8-a.u. cell. Various HfC
1−x
N
x
compo-
sitions were obtained by substituting carbon atoms with nitrogen uti-
lizing the Special Quasi-random Structure (SQS) method [40]. The in-
fluence of structural defects was investigated by the systematic removal
of either metallic or non-metallic species, also applying the SQS code.
The Brillouin zone of the cubic supercell was sampled with
6 × 6 × 6 k-points, and a plane wave cut-off energy of 500 eV was
chosen based on convergence tests guaranteeing a total energy accuracy
of at least 1 meV/at. Equilibrium lattice parameters and ground-state
energies were obtained by relaxing the supercell volumes, shapes, and
atomic positions (ISIF = 3 tag in VASP). Elastic constants of the
HfC
1−x
N
x
supercells were calculated using the Universal Linear-In-
dependent Coupling Strain (ULICS) method [41].
Polycrystalline bulk and shear modulus (B and G), Poisson ratio (ν),
and Young's modulus, E, were calculated applying Hill's average [42] of
the Reuss and Voigt bulk moduli (B
R
and B
V
) and shear moduli (G
R
and
G
V
), respectively.
4. Results and discussion
The angular distribution of metal and non-metal species (due to
their diverging atomic masses of Hf = 178.49 u compared to C = 12.01
u) during sputter deposition (e.g. HfeC) was the main driving force for
comparing two different target to substrate alignments [43]. Based on
the geometrically differences for these two setups, distinct target cur-
rents (0.5 and 0.8 A, respectively), as well as substrate heater tem-
peratures were used to gain similar coating morphologies. The chemical
compositions of all coatings (Hf
1-y
C
y
and Hf
1-y-z
C
y
N
z
) were analyzed by
TOF-ERDA. During these measurements, Hafnium (Hf) Carbon (C), and
Nitrogen (N), as well as Argon (Ar) and Oxygen (O) were detected.
Expected systematic uncertainties in the Hf-C-N system are typically
found to be at most 5 to 10% of the detected concentrations for absolute
measurements free from standards. Nevertheless, much higher preci-
sion is obtained for inter-sample comparison, as performed within this
study [43]. For the binary coatings deposited with the parallel and
tilted configuration, we obtained Hf
0.6
C
0.4
and Hf
0.43
C
0.57
, respectively.
The difference in stoichiometry is related to the angular distribution
during sputtering. In Fig. 1, an isothermal section of the ternary Hf-C-N
system is presented including the δ-Hf(C, N)
1-x
(light blue) as well as the
α-Hf phase-field as suggested by S. Binder [44] for an experimentally
validated isothermal section at 1150 °C (the temperature has no direct
relation to the synthesis or growth process and is only based on the
limited data set for isothermal sections within the ternary Hf-C-N
system). All coatings deposited are plotted in half-filled cubes and open
diamonds for the parallel and tilted set-up, respectively. In agreement
with a previous study [22], are binary HfeC compounds most likely off-
stoichiometric phases (Hf
0.43
C
0.57
and Hf
0.6
C
0.4
), which is typically
observed when using TMC compound targets in Ar atmospheres (in
relation to the angular distribution) [45]. Furthermore, already rela-
tively low f
[N2]norm
results in high nitrogen contents on the non-metallic
sublattice, x (x = N/(C + N)). This behavior is also highlighted in
Fig. 2, where the non-metallic sublattice occupation is plotted against
the nitrogen flow rate and underlines that we were unable to obtain
compositions below x = 0.50. The trends are similar for both target
configurations and show a big increase in, x, from f
[N2]norm
= 0 to 0.05,
followed by a slight rise for higher nitrogen flow rates (from 0.05 to
0.20). The occupation of nitrogen on the non-metallic sublattice is ac-
companied not only by a decrease in Carbon but also by a decrease in
Hafnium, resulting in Hafnium sub-stoichiometric thin films, deviating
from a hypothetical HfC-HfN tie line – see the right border of the δ-Hf
(C, N)
1-x
phase-field indicated in Fig. 1. This behavior is similar to the
Ta-C-N system [16], where it is less pronounced. The reduction of the
deposition temperature (hence surface mobility of incoming species)
enhances this effect further, see open cube (T
sub
= 250 °C) in Fig. 1,
whereby the nitrogen content increases on the cost of the Hf content.
The substrate surface temperature also influences the residual stress
state within the coatings, increasing the compressive character for de-
creased surface temperatures from −2.2 GPa to −2.8 GPa for 380 °C
and 250 °C, respectively. As the net stress state is influenced by several
factors, e.g. temperature and thickness (after R. Abermann [46]), a
higher growth rate for T
sub
= 380 °C (30 nm/min compared to 22 nm/
min) most likely dominates the compressive influence. The difference in
growth rate is mostly related to the varied nitrogen flow rate ratio for
these specific coatings, respectively. Nevertheless, lower surface tem-
peratures, as well as target currents accompanied by higher f
[N2]norm
might promote the introduction of reactive working gas nitrogen onto
the non-metallic sublattice.
Fig. 1. Isothermal section (1150 °C [44]) of the ternary Hf-C-N system with
indicated compositions of the parallel and tilted target configuration obtained
by ERDA. Open diamonds refer to the tilted set up (T
sub
= 415 °C), whereas
half-filled cubes indicate coatings obtained from the parallel configuration
(T
sub
= 380 °C). The open cube refers to the composition at a lower substrate
temperature (T
sub
= 250 °C) in the parallel set up with f
[N2]norm
=0.05.
Fig. 2. Influence of the nitrogen to total gas flow ratio f
[N2]norm
on the nitrogen
sublattice occupation x = N/(C + N). The tilted set up is indicated by open
diamonds, whereas half-filled squares refer to the parallel configuration.
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
3
Furthermore, the effect of slightly different angular sputter dis-
tributions was also obvious in the colors of the coatings. In general,
HfeC coatings appear dark grey up to black, whereas Hf-C-N coatings
have a coppery color. Hf-C-N coatings deposited with the face to face
arrangement, obtain a color-gradient, from coppery to black to coppery
over a distance of 20 mm (substrate size) being centered with respect to
the race track. This gradient is contributed to the angular sputter dis-
tribution of Hf and C as well as the relatively short substrate to target
distance (50 mm). The black regions exhibit about 7 at.% more carbon
on the extend of Hf, whereas the nitrogen content remains the same
throughout the substrate width. However, no gradient is observed in
the coatings deposited using the tilted arrangement. Due to this fact, all
further mechanical characterizations and annealing experiments are
done on coatings deposited in the tilted setup.
Fig. 3 a and b show the structural evolution with increasing nitrogen
content for both experimental set-ups. All coatings, Hf
1-y
C
y
as well as
Hf
1-y-z
C
y
N
z
, exhibit face-centered cubic (fcc) structures. Furthermore,
we observe a broadening of the peaks with increasing f
[N
2
]norm
in-
dicating a decrease in crystallite size as well as being a hint for higher
defect densities within the columns. The incorporation of nitrogen in
the thin films leads to a shift of the 2Θ values to higher angles towards
HfeN, which indicates a smaller lattice parameter [47,48] – visible for
pure Hf
1-y
C
y
compared to nitrogen alloyed coatings.
To further evaluate the obtained structural results, we performed
TEM investigations shown in Fig. 4 for Hf
0.43
C
0.57
(a to c) and
Hf
0.35
C
0.24
N
0.41
(d to f), respectively. Films on Si substrate have been
prepared for these TEM anlysis. Here a focus is to prove the presence of
amorphous phases, as a clear picture about the over stoichiometric
HfeC and Hf-C-N (especially on the non-metal sublattice) compositions
is not obvious. Both coatings show a dense columnar growth mor-
phology, while the addition of nitrogen leads to a decreasing column
size for the Hf
0.35
C
0.24
N
0.41
coatings – as also seen in the XRD spectra.
For Hf
0.43
C
0.57
the SAED analysis reavealed distinct spots on a ring-
pattern clearly corresponding to fcc structures – see Fig. 4b. These spots
get more blurred around the diffraction spots for Hf
0.35
C
0.24
N
0.41
due to
smaller diffracting domains (e.g. grains and/or crystallites) – see
Fig. 4e. As the dark-field image in Fig. 4e consist out of nice shaped
columns, also a highly crystalline state is assumed for the
Hf
0.35
C
0.24
N
0.41
coating. The high-resolution image of Hf
0.43
C
0.57
(see
Fig. 4c) give a detailed insight on the crystal structure and the mor-
phology at the column boundaries, which are marked with white da-
shed lines. In addition FFT analysis of the marked areas have been
conducted to prove the high crystallinity of these areas. The FFT clearly
revealed dot-like patterns, suggesting also the crystalline character on
this small length scale. For the Hf
0.35
C
0.24
N
0.41
coating, also no obvious
amorphous grain boundary phases are visible at a higher resolution (see
Fig. 4f). This is an indication, that carbon and nitrogen are sharing the
octohetral sites on the non-metal sublattice, and may minor amorphous
CNx regions are formed as suggested by further studies [45] on TMC
based caotings (reporting access carbon accumalated on grain bound-
aries in nano-composite morphologies). However, with respect to the
access non-metal species in the observed chemistry – and limited ex-
perimental proof – we have to admit, that detailed analysis, e.g. using
XPS, could draw a more clear picture about the grain boundary mor-
phology (e.g. very thin CN
x
interface phases) at this nano-scale level.
The obtained experimental findings on the structural evolution of
Hf-C-N were also verified by DFT calculations applying VASP. Fig. 5a
shows that the Energy of formation (E
f
) lowers with increasing nitrogen
content on the non-metallic sublattice, x (HfC
1-x
N
x
), suggesting stabi-
lization of HfeN based structures (using PBE potentials). Compared to
calculations for the TaC
1-x
N
x
system [16] the decrease in E
f
and hence
the driving force to form nitrogen-rich compounds is much stronger for
HfC
1-x
N
x
. This might be the reason why we observed such a high affi-
nity of incorporating nitrogen during the reactive sputter process. Va-
cancies (vacancy concentration of 6% on the corresponding sublattice)
only slightly influence the phase formation of HfC
1-x
N
x
structures as the
defect-free cells are preferred over the full compositional range – please
compare open with half-filled and full squares in Fig. 4a, respectively.
Nevertheless, Carbon-rich compositions are more prone to non-metal
vacancies, whereas HfN prefers metal vacancies. Fig. 5b presents the
lattice parameter with respect to the sublattice occupation, x, applying
LDA and PBE exchange-correlation potential (xc). The lattice parameter
decreases with increasing nitrogen content as it was also observed in
the XRD analysis of the thin films. Both potentials result in the same
trend but with different absolute values. The calculated lattice para-
meters are compared with lattice parameters determined by XRD (open
diamonds, green) using a LaB
6
standard powder as a reference. Based
on simplifications, we used the sublattice occupation x = N/(C + N)
for the experimental compositions to compare the calculated with the
experimental lattice constants. Nevertheless, the experimental trend
matches the PBE lattice parameter quite well. The lattice constants of
the XRD references, 4.63 Å [48] and 4.51 Å [47] for HfC and HfN,
respectively, are exactly in-between the values obtained by LDA (HfC
a
c
= 4.58 Å, HfN a
c
= 4.46 Å) and PBE (HfC a
c
= 4.65 Å, HfN
a
c
= 4.54 Å).
As already mentioned above, the mechanical properties, especially
hardness and elastic modulus, of transition metal carbides and nitrides
are strongly influenced by their VEC – obtaining a maximum at VEC of
8.4 e/f.u. This correlation was proven for TiC
1-x
N
x
, but still, some
Fig. 3. Structure evolution with increasing nitrogen content for tilted (a) and
parallel (b) target to substrate configuration. Spectra are obtained from pow-
dered coatings. The reference fcc-structured HfeC and HfeN pattern are taken
from [47,48].
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
4
questions arise for specific systems with respect to the way how the VEC
(e.g. through vacancies) is achieved [13,10,17]. However, a correlation
between the fracture tolerance and VEC, and specifically the nitrogen
sublattice occupation in transition metal carbide-based materials, is
rare. Therefore, this work gives additional information on how the
sublattice occupation, hence the VEC, affects the ductility in HfC
1-x
N
x
.
To develop this correlation, we compared the semi-empirical criteria B/
G ratio (Pugh-ratio) [49] as well as Pettifor's Cauchy pressure (C
12
-C
44
)
[50] with experimentally evaluated fracture characteristics. In addi-
tion, Niu et al. proposed recently a universal ductile to brittle criterion
((C
12
-C
44
)/E) [51]. Fig. 6a gives the Bulk, B (open squares), and Shear
modulus, G (half-filled squares) – on which base the Pugh-ratio is cal-
culated – in comparison to literature data (open stars) [13]. The bulk
modulus is in good agreement with the endpoints provided by the lit-
erature. Furthermore, the increase in B accompanied by a decrease in G
suggests higher Pugh-ratios. In Fig. 6b C
12
-C
44
and (C
12
-C
44
)/E are
plotted over the VEC, whereby more positive values (above the Pettifor
criterion) indicate a more ductile character. The criteria result in the
same trend with a small decrease in ductility for VEC = 8.25 e/f.u.
(x = 0.25), followed by a subsequent increase with increasing nitrogen
content (VEC). However, for all compositions, we are below the Pettifor
criterion and hence in the so-called brittle regime.
Comparing the Pugh-ratio of the group IV TM‑carbonitride, HfC
1-
x
N
x
, with the group V TM‑carbonitride, TaC
1-x
N
x
[16], with experi-
mentally obtained K
IC
values emphasizes an almost continuous increase
in ductility with increasing VEC, see Fig. 7. Here the fracture toughness,
K
IC
, is plotted on the left y-axis in relation to the VEC and compared
with the B/G ratio plotted on the right y-axis for HfC
1-x
N
x
(open
squares) and TaC
1-x
N
x
(open circles), respectively. This is in excellent
agreement with calculations from Ref. [13] indicating a positive effect
on the fracture characteristics for VEC up to 10. All indicated toughness
Fig. 4. Cross sectional bright images of Hf
0.43
C
0.57
(a) as well as a dark field micrograph of Hf
0.35
C
0.24
N
0.41
(d), respectively (films depsoted on Si substrates). The
corresponding SAED analysis to Hf
0.43
C
0.57
and Hf
0.35
C
0.24
N
0.41
are given in (b) and (e), respectively. Bright field images with higher resolution (c) and (f) provide a
closer look at the column boundaries of the respective coatings. In (c) column boundaries are marked with dashed lines and the insets show the FFTs of the
corresponding columns.
Fig. 5. (a) Trend of the calculated E
f
over the nitrogen sublattice occupation, x,
for perfect and defected supercells. Defect-free cells possess the lowest E
f
and
are therefore most preferred. (b) Comparison of the two used xc potentials LDA
(open cubes) and PBE (open circles) based on the lattice parameter, a
c
. The
calculated values are furthermore compared with the experimentally de-
termined ones (open diamonds) from XRD and references (star symbols)
[47,48].
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
5
values, K
IC
, have been obtained by equal single cantilever beam
bending tests (values obtained for at least 5 or more beams). Values for
the Ta-C-N system are taken from a previous study [4]. Here the authors
want to highlight, that the VEC of all compounds denotes the valence
electrons per unit cell as commonly reported in the literature for these
compounds [10] only taking into account the carbon substitution with
nitrogen in the metal-normalized notation (HfC
1-x
N
x
and TaC
1-x
N
x
).
Therefore, sub-stoichiometries are more or less neglected within this
depiction. For both systems, the introduction of nitrogen increases the
fracture toughness following the theoretical predictions. The fracture
toughness of the HfC
1-x
N
x
increases from 1.89 ± 0.15 to
2.33 ± 0.18 MPa·m
1/2
for Hf
0.43
C
0.57
and Hf
0.35
C
0.30
N
0.35
respectively.
The experimentally observed slight increase for HfC
1-x
N
x
also fits the B/
G prediction, as the ratio is growing slower with VEC compared to TaC
1-
x
N
x
. However, the effect of grain size and sub-stoichiometry is neglected
here and therefore an absolute comparison is difficult.
To further validate the quality of the theoretically evaluated elastic
constants (e.g. B or G), we compare the calculated Young's Modulus
with the experimentally determined mechanical properties, especially
Indentation Modulus, of the deposited thin films in Fig. 8a. Hardness
and indentation modulus were measured on sapphire substrates
whereas residual stresses were evaluated on silicon substrates. The
predicted Young's Modulus is in good agreement with the measured
Indentation Modulus values – see Fig. 8a. For Carbon-rich composition,
the slight deviation – also compared to literature data [13,22,52,53] –
may be related to the deviations in sub-stoichiometry for the binary Hf
1-
y
C
y
system. Nevertheless, the trend follows the already observed results,
suggesting moderately reduced Young's modulus with increasing ni-
trogen contents. In contrast to that, the hardness, see Fig. 8b, starts with
a slight increase from 37.8 ± 2.1 to 39.9 ± 2.7 GPa for Hf
0.43
C
0.57
to
Hf
0.35
C
0.30
N
0.35
, respectively. When the nitrogen content is further in-
creased the hardness drops to 33.6 ± 1.9 GPa for Hf
0.32
C
0.20
N
0.48
. The
residual stress – as plotted in Fig. 8b right y-axis – decreases from
−3.2 GPa to −2.1 GPa for x = 0 to x = 0.70, respectively. It follows a
different trend than the hardness, which we suggest is due to the
bonding nature and not the residual stress, assuming that the residual
stress state follows the same trend for the films grown on both sapphire
and on silicon substrates. However, we have to mention that: (i) the
residual stress states may follow a different trend in the samples grown
on sapphire (from which hardness data is extracted) from those grown
on silicon substrates and (ii) even if the stress states were to follow the
same trend for the films on both sapphire and silicon, differences in
grain sizes and film textures on sapphire and silicon can lead to dif-
ferent trends in hardness of films on sapphire and silicon substrates.
Neverthelss, the higher compressive stresses are also in good agreement
to the increased deposition rates for Nitrogen poor gas mixture.
As already mentioned, Nitrogen alloying in the HfeC system is
predicted to increase the melting temperature [20]. To investigate the
potential change in the mechanical properties of the HfeC and Hf-C-N
coatings due to thermal loading, we performed different annealing tests
Fig. 7. Empirical criterion for ductility, Pugh ratio (B/G), for Group IV (Hf) and
Group V (Ta) TM-Carbonitrides with respect to the VEC [4]. These calculations
are compared with experimentally determined K
IC
values – open square denotes
to Hf and open circles to the Ta based systems, respectively. All experimental
data analysis were done on Si substrates, with growth temperatures of 415 °C
for the Hf-based and 450 °C for the Ta-based films, respectively.
Fig. 8. (a) Experimentally determined Indentation Modulus in relation to the
sublattice occupation in comparison to DFT based Young's Modulus values and
literature data for HfC and HfN [13,22,51,52] (star symbols). (b) Hardness
(open squares) and residual stress (open diamonds) as a function of x (N/
(C + N)) in comparison with hardness values from literature [22,51,52] (star
symbols). The indentation analysis was performed on Si substrates deposited in
the tilted set-up at 415 °C.
Fig. 6. (a) Bulk modulus (B) and shear modulus (G) as a function of the ni-
trogen sublattice occupation, x, and hence VEC. (b) Cauchy pressure (C
12
-C
44
)
and (C
12
-C
44
)/E in relation to the sublattice occupation with indicated Pettifor
criterion – (C
12
-C
44
) > 0 is expected to be ductile. Star symbols represent
literature [13] values for B and C
12
-C
44
.
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
6
and subsequent mechanical characterization, see Fig. 9a and b. We
annealed the thin films deposited on sapphire in a vacuum furnace at
600 °C, 800 °C, 1000 °C, and 1200 °C. We observed no visible change in
the color of the thin films after annealing up to 800 °C. Beyond this
temperature, we observed a change in color towards more greenish or
blueish. Furthermore, delamination occurred for the Hf
0.43
C
0.57
coating
at 1200 °C. However, we assume that in this case, the delamination took
place because of the reduced adhesion to the substrate, assumed to be
due to high compressive stresses. For all composition tested, the hard-
ness remains almost constant until 800 °C followed by a decrease at
higher temperatures. With increasing nitrogen content, the slope for
decreasing hardness levels out, and HfC
1-x
N
x
exhibits a higher hardness
after annealing. For example, the initial hardness value for Hf
0.43
C
0.57
is
37.8 ± 2.1 GPa and after annealing at 1000 °C, it reduces to
25.2 ± 1.3 GPa, whereas for Hf
0.35
C
0.24
N
0.41
the hardness is
37.9 ± 2.4 GPa and after annealing to 1000 °C it settles at
33.0 ± 1.7 GPa, respectively. For all Hf-C-N coatings, the Young's
modulus stays approximately constant over the annealing temperature.
5. Conclusion
We investigated the influence of nitrogen alloying on the synthesis,
structure, and thermomechanical properties of magnetron sputtered
HfeC thin films. The influence of the non-metallic sublattice occupation
x = N/(C + N) – hence the variation of the valance electron con-
centration – on structural and mechanical properties, was studied both
theoretically as well as experimentally. The thin films were deposited
by non-reactive and reactive magnetron sputtering from a ceramic
HfeC target. Already at low nitrogen partial pressures, we obtained a
high nitrogen content within the coatings, which was even intensified
by a lower substrate temperature as well as power density. Therefore,
only compositions with x above 0.50 could be deposited. All HfeC and
Hf-C-N coatings exhibit a single-phase fcc structure with dense co-
lumnar morphologies. DFT based calculations on the lattice constants –
decrease from a
c
= 4.65 Å for HfC to a
c
= 4.54 Å for HfN – are in
excellent agreement with experimental values a
c
= 4.67 Å for
Hf
0.43
C
0.57
and a
c
= 4.61 Å for Hf
0.35
C
0.30
N
0.35
(x = 0.54,
HfC
0.46
N
0.54
), respectively. In addition, the predicted elastic constants –
B/G ratio as well as Cauchy pressure – also suggest a more ductile
character with increasing nitrogen content, implying an increased VEC.
This prediction could be also validated experimentally with an increase
in fracture toughness from 1.89 ± 0.15 to 2.33 ± 0.18 MPa·m
1/2
for
Hf
0.43
C
0.57
to Hf
0.35
C
0.30
N
0.35
, respectively. As the VEC ranges between
8.0 e/f.u. to 9.0 e/f.u. for HfC
1-x
N
x
based composition the increase is
moderate, as an optimum VEC is suggested slightly below 10. The
measured hardness exhibits a maximum of 39.9 ± 2.7 GPa for
Hf
0.35
C
0.30
N
0.35
. The Young's modulus ranges from 532 to 426 GPa for
Hf
0.43
C
0.57
and Hf
0.32
C
0.20
N
0.48
, respectively. The compressive stress
state decreases (−3.2 to −2.1 GPa for x = 0 up to 0.70) with in-
creasing nitrogen content. Nevertheless, we want to mention, that these
mechanical enhancements need to be seen in correlation also with
morphological changes – which were not clearly observable within this
study – and therefore related to the VEC. Furthermore, the non-metal
alloying of nitrogen increases the hardness of Hf-C-N coatings after
annealing to a maximum of 25.4 ± 1.7 GPa for Hf
0.35
C
0.24
N
0.41
at
1200 °C. In summary, the obtained results indicate the correlation of
fracture characteristics and valence electron concentration for transi-
tion metal ceramics – such as HfeC or Hf-C-N. Furthermore, the great
potential of Hf-C-N coatings – for well-defined deposition conditions to
reach the desired C to N ratios could be validated for its thermo-
mechanical properties.
Supplementary data to this article can be found online at https://
doi.org/10.1016/j.surfcoat.2020.126212.
CRediT authorship contribution statement
T. Glechner:Conceptualization, Software, Investigation, Writing -
original draft.S. Lang:Investigation.R. Hahn:Investigation, Writing -
review & editing.M. Alfreider:Investigation.V. Moraes:Investigation.D.
Primetzhofer:Investigation.J. Ramm:Writing - review & editing.S.
Kolozsvári:Resources.D. Kiener:Investigation.H. Riedl:Supervision,
Conceptualization, Writing - review & editing, Project administration.
Declaration of competing interest
The authors declare that they have no known competing financial
interests or personal relationships that could have appeared to influ-
ence the work reported in this paper.
Acknowledgments
The financial support by the Austrian Federal Ministry for Digital
and Economic Affairs and the National Foundation for Research,
Technology, and Development is gratefully acknowledged (Christian
Doppler Laboratory “Surface Engineering of high-performance
Components”). We also thank for the financial support of Plansee SE,
Plansee Composite Materials GmbH, and Oerlikon Balzers, Oerlikon
Surface Solutions AG. In addition, we want to thank the X-ray center
(XRC) of TU Wien for beam time as well as the electron microscopy
center - USTEM TU Wien - for using the SEM and TEM facilities. The
computational results presented have been achieved using the Vienna
Scientific Cluster (VSC). The authors acknowledge the TU Wien
Bibliothek for financial support through its Open Access Funding
Program.
References
[1] W.G. Fahrenholtz, G.E. Hilmas, Oxidation of ultra-high temperature transition
metal diboride ceramics, Int. Mater. Rev. 57 (2012) 61–72.
[2] W.G. Fahrenholtz, E.J. Wuchina, W.E. Lee, Y. Zhou, Ultra-high Temperature
Ceramics: Materials for Extreme Environment Applications, John Wiley & Sons,
2014.
[3] S. Kiani, J. Yang, S. Kodambaka, Nanomechanics of refractory transition-metal
carbides: a path to discovering plasticity in hard ceramics, J. Am. Ceram. Soc. 98
(2015) 2313–2323.
[4] T. Glechner, R. Hahn, T. Wojcik, D. Holec, S. Kolozsvári, H. Zaid, S. Kodambaka,
P.H. Mayrhofer, H. Riedl, Assessment of ductile character in superhard Ta-C-N thin
films, ta Mater Ac (2019).
[5] H. Kindlund, D.G. Sangiovanni, I. Petrov, J.E. Greene, L. Hultman, A review of the
intrinsic ductility and toughness of hard transition-metal nitride alloy thin films,
Fig. 9. Hardness (a) and Indentation Modulus (b) of selected coatings after
vacuum annealing at 600, 800, 1000, and 1200 °C, respectively. The indenta-
tion analysis was performed on films grown (T
sub
= 415 °C) on sapphire sub-
strates.
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
7
Thin Solid Films 688 (2019) 137479.
[6] R. Hollerweger, H. Riedl, M. Arndt, S. Kolozsvari, S. Primig, P.H. Mayrhofer,
Guidelines for increasing the oxidation resistance of Ti-Al-N based coatings, Thin
Solid Films 688 (2019) 137290.
[7] D.E. Hajas, M. to Baben, B. Hallstedt, R. Iskandar, J. Mayer, J.M. Schneider,
Oxidation of Cr2AlC coatings in the temperature range of 1230 to 1410°C, Surf.
Coat. Technol. 206 (2011) 591–598.
[8] P.H. Mayrhofer, R. Rachbauer, D. Holec, F. Rovere, J.M. Schneider, 4.14 - protec-
tive transition metal nitride coatings, in: V.K. Sarin, L. Llanes, D. Mari, C.E. Nebel
(Eds.), Comprehensive Materials Processing, Elsevier, Oxford, 2014, pp. 355–388.
[9] H. Holleck, Material selection for hard coatings, J. Vac. Sci. Technol. A 4 (1986)
2661–2669.
[10] S.-H. Jhi, J. Ihm, S.G. Louie, M.L. Cohen, Electronic mechanism of hardness en-
hancement in transition-metal carbonitrides, Nature 399 (1999) 132.
[11] D.G. Sangiovanni, L. Hultman, V. Chirita, Supertoughening in B1 transition metal
nitride alloys by increased valence electron concentration, Acta Mater. 59 (2011)
2121–2134.
[12] D.G. Sangiovanni, V. Chirita, L. Hultman, Electronic mechanism for toughness en-
hancement in TixM1-xN (M=Mo and W), Phys. Rev. B: Condens. Matter Mater.
Phys. 81 (2010) 104107.
[13] K. Balasubramanian, S.V. Khare, D. Gall, Valence electron concentration as an in-
dicator for mechanical properties in rocksalt structure nitrides, carbides and car-
bonitrides, Acta Mater. 152 (2018) 175–185.
[14] X.-X. Yu, G.B. Thompson, C.R. Weinberger, Influence of carbon vacancy formation
on the elastic constants and hardening mechanisms in transition metal carbides, J.
Eur. Ceram. Soc. 35 (2015) 95–103.
[15] N. Koutná, D. Holec, O. Svoboda, F.F. Klimashin, P.H. Mayrhofer, Point defects
stabilise cubic Mo-N and Ta-N, J. Phys. D. Appl. Phys. 49 (2016) 375303.
[16] T. Glechner, P.H. Mayrhofer, D. Holec, S. Fritze, E. Lewin, V. Paneta,
D. Primetzhofer, S. Kolozsvári, H. Riedl, Tuning structure and mechanical proper-
ties of Ta-C coatings by N-alloying and vacancy population, Sci. Rep. 8 (2018)
17669.
[17] H. Riedl, T. Glechner, T. Wojcik, N. Koutná, S. Kolozsvári, V. Paneta, D. Holec, D.
Primetzhofer, P.H. Mayrhofer, Influence of carbon deficiency on phase formation
and thermal stability of super-hard TaCy thin films, Scr. Mater. 149 (2018/5)
150–154.
[18] M. Raza, D. Cornil, J. Cornil, S. Lucas, R. Snyders, S. Konstantinidis, Oxygen va-
cancy stabilized zirconia (OVSZ); a joint experimental and theoretical study, Scr.
Mater. 124 (2016) 26–29.
[19] M. to Baben, M. Hans, D. Primetzhofer, S. Evertz, H. Ruess, J.M. Schneider,
Unprecedented thermal stability of inherently metastable titanium aluminum ni-
tride by point defect engineering, Materials Research Letters 5 (2017) 158–169.
[20] Q.-J. Hong, A. van de Walle, Prediction of the material with highest known melting
point from ab initio molecular dynamics calculations, Phys. Rev. B Condens. Matter
92 (2015) 020104.
[21] Q.-J. Hong, Methods for melting temperature calculation, phd, California Institute
of Technology, http://resolver.caltech.edu/CaltechTHESIS:11092014-074023936,
(2015) , Accessed date: 9 August 2016.
[22] H. Lasfargues, T. Glechner, C.M. Koller, V. Paneta, D. Primetzhofer, S. Kolozsvári,
D. Holec, H. Riedl, P.H. Mayrhofer, Non-reactively sputtered ultra-high temperature
Hf-C and Ta-C coatings, Surf. Coat. Technol. 309 (2017) 436–444.
[23] D.T. Livey, The high temperature stability of oxides and sulphides, Journal of the
Less Common Metals 1 (1959) 145–151.
[24] P. Zeman, Š. Zuzjaková, P. Mareš, R. Čerstvý, M. Zhang, J. Jiang, E.I. Meletis,
J. Vlček, Superior high-temperature oxidation resistance of magnetron sputtered
Hf–B–Si–C–N film, Ceram. Int. 42 (2016) 4853–4859.
[25] Y. Shen, J.C. Jiang, P. Zeman, V. Šímová, J. Vlček, E.I. Meletis, Microstructure
evolution in amorphous Hf-B-Si-C-N high temperature resistant coatings after an-
nealing to 1500 °C in air, Sci. Rep. 9 (2019) 3603.
[26] W.G. Fahrenholtz, G.E. Hilmas, Ultra-high temperature ceramics: materials for ex-
treme environments, Scr. Mater. 129 (2017) 94–99.
[27] W.F. Piedrahita, W. Aperador, J.C. Caicedo, P. Prieto, Evolution of physical prop-
erties in hafnium carbonitride thin films, J. Alloys Compd. 690 (2017) 485–496.
[28] W. Li, D. Ming-hui, Z. Hong-sen, Z. Bin, Study on HfC N coatings deposited on
biomedical AISI 316L by radio-frequency magnetron sputtering, J. Alloys Compd.
730 (2018) 219–227, https://doi.org/10.1016/j.jallcom.2017.09.289.
[29] P. Ström, P. Petersson, M. Rubel, G. Possnert, A combined segmented anode gas
ionization chamber and time-of-flight detector for heavy ion elastic recoil detection
analysis, Rev. Sci. Instrum. 87 (2016) 103303.
[30] W.C. Oliver, G.M. Pharr, An improved technique for determining hardness and
elastic modulus using load and displacement sensing indentation experiments, J.
Mater. Res. 7 (1992) 1564–1583.
[31] R. Saha, W.D. Nix, Effects of the substrate on the determination of thin film me-
chanical properties by nanoindentation, Acta Mater. 50 (2002) 23–38.
[32] G.C.A.M. Janssen, M.M. Abdalla, F. van Keulen, B.R. Pujada, B. van Venrooy,
Celebrating the 100th anniversary of the Stoney equation for film stress: develop-
ments from polycrystalline steel strips to single crystal silicon wafers, Thin Solid
Films 517 (2009) 1858–1867.
[33] K. Matoy, H. Schönherr, T. Detzel, T. Schöberl, R. Pippan, C. Motz, G. Dehm, A
comparative micro-cantilever study of the mechanical behavior of silicon based
passivation films, Thin Solid Films 518 (2009) 247–256.
[34] S. Brinckmann, C. Kirchlechner, G. Dehm, Stress intensity factor dependence on
anisotropy and geometry during micro-fracture experiments, Scr. Mater. 127 (2017)
76–78.
[35] S. Brinckmann, K. Matoy, C. Kirchlechner, G. Dehm, On the influence of micro-
cantilever pre-crack geometries on the apparent fracture toughness of brittle ma-
terials, Acta Mater. 136 (2017) 281–287.
[36] G. Kresse, J. Furthmüller, Efficient iterative schemes for ab initio total-energy cal-
culations using a plane-wave basis set, Phys. Rev. B Condens. Matter 54 (1996)
11169–11186.
[37] G. Kresse, D. Joubert, From ultrasoft pseudopotentials to the projector augmented-
wave method, Phys. Rev. B Condens. Matter 59 (1999) 1758–1775.
[38] W. Kohn, L.J. Sham, Self-consistent equations including exchange and correlation
effects, Phys. Rev. 140 (1965) A1133–A1138.
[39] J.P. Perdew, K. Burke, M. Ernzerhof, Generalized gradient approximation made
simple, Phys. Rev. Lett. 77 (1996) 3865–3868.
[40] S. Wei, L.G. Ferreira, J.E. Bernard, A. Zunger, Electronic properties of random al-
loys: special quasirandom structures, Phys. Rev. B Condens. Matter 42 (1990)
9622–9649.
[41] R. Yu, J. Zhu, H.Q. Ye, Calculations of single-crystal elastic constants made simple,
Comput. Phys. Commun. 181 (2010) 671–675.
[42] R. Hill, The elastic behaviour of a crystalline aggregate, Proc. Phys. Soc. A. 65
(1952) 349.
[43] J. Neidhardt, S. Mráz, J.M. Schneider, E. Strub, W. Bohne, B. Liedke, W. Möller,
C. Mitterer, Experiment and simulation of the compositional evolution of Ti–B thin
films deposited by sputtering of a compound target, J. Appl. Phys. 104 (2008)
063304.
[44] S. Binder, W. Lengauer, P. Ettmayer, J. Bauer, J. Debuigne, M. Bohn, Phase equi-
libria in the systems Ti-C-N, Zr-C-N and Hf-C-N, J. Alloys Compd. 217 (1995)
128–136.
[45] U. Jansson, E. Lewin, Sputter deposition of transition-metal carbide films — a cri-
tical review from a chemical perspective, Thin Solid Films 536 (2013) 1–24.
[46] R. Abermann, Measurements of the intrinsic stress in thin metal films, Vacuum 41
(1990) 1279–1282.
[47] International Center of Diffraction Data, Powder Diffraction File 04-002-0653,
(2019).
[48] International Center of Diffraction Data, Powder Diffraction File 04-004-8343,
(2019).
[49] S.F. Pugh, XCII. Relations between the elastic moduli and the plastic properties of
polycrystalline pure metals, The London, Edinburgh, and Dublin Philosophical
Magazine and Journal of Science 45 (1954) 823–843.
[50] D.G. Pettifor, Theoretical predictions of structure and related properties of inter-
metallics, Mater. Sci. Technol. 8 (1992) 345–349.
[51] H. Niu, X.-Q. Chen, P. Liu, W. Xing, X. Cheng, D. Li, Y. Li, Extra-electron induced
covalent strengthening and generalization of intrinsic ductile-to-brittle criterion,
Sci. Rep. 2 (2012) 718.
[52] C. Hu, X. Zhang, Z. Gu, H. Huang, S. Zhang, X. Fan, W. Zhang, Q. Wei, W. Zheng,
Negative effect of vacancies on cubic symmetry, hardness and conductivity in
hafnium nitride films, Scr. Mater. 108 (2015) 141–146.
[53] M.M.S. Villamayor, J. Keraudy, T. Shimizu, R.P.B. Viloan, R. Boyd, D. Lundin,
J.E. Greene, I. Petrov, U. Helmersson, Low temperature ( T s / T m \textless 0.1)
epitaxial growth of HfN/MgO(001) via reactive HiPIMS with metal-ion synchro-
nized substrate bias, J. Vac. Sci. Technol. A 36 (2018) 061511.
T. Glechner, et al. Surface & Coatings Technology 399 (2020) 126212
8
... Several strategies [20] have been established to enhance the fracture toughness of TMC based thin films, i.e. nitrogen alloying to form TM-carbonitrides [21]. This approach was experimentally found to be highly effective for the Ta-C-N [22,23] system and further underlined in a study on Hf-C-N, where the theoretically proposed [21] dependency between ductility and valence electron concentration (VEC) was confirmed [24]. According to this latter work, the fracture toughness of WC (VEC = 10) should be even higher compared to e.g. ...
... The determined fracture toughness values K IC are 3.3 ± 0.33 MPa·m 1/2 for WC 0.67 , 3.2 ± 0.24 MPa·m 1/2 for WC 0.78 , and 1.9 ± 0.24 MPa·m 1/2 for WC 3.38 (4 beams tested for each coating). The high K IC value observed for the crystalline WC 0.67 is a further confirmation of the direct correlation between the VEC and the fracture toughness of TM-carbides and nitrides, as already proposed theoretically and experimentally [21,23,24,35]. In comparison, the K IC values of group IV and V TM-carbides such as Hf-C and Ta-C are below 2 MPa·m 1/2 [23,24]. ...
... The high K IC value observed for the crystalline WC 0.67 is a further confirmation of the direct correlation between the VEC and the fracture toughness of TM-carbides and nitrides, as already proposed theoretically and experimentally [21,23,24,35]. In comparison, the K IC values of group IV and V TM-carbides such as Hf-C and Ta-C are below 2 MPa·m 1/2 [23,24]. Furthermore, it is astonishing that WC 0.67 and WC 0.78 exhibit very similar K IC values despite having a drastically different morphology, which in turn suggest that the film growth characteristics only have a minor influence on the overall fracture resistance of these films. ...
Article
Full-text available
During the growth of WC based thin films, carbon can be introduced by either a non-reactive or reactive deposition route. In this study, we compare the influence of the carbon origin on the coating properties, sputtering three different target materials – a ceramic WC, a ceramic WC including a conventional cobalt binder, and a metallic tungsten (W) target – in reactive (acetylene, C2H2) as well as non-reactive (pure Ar) atmospheres. The morphology changes, independently to the target type and atmosphere used, from crystalline (hex-W2C rich to pure fcc-WCx) to a nanocomposite (fcc-WCx nanometre sized grains embedded in an amorphous matrix) structure, up to amorphous coatings, only dominated by the prevalent C/W ratio. The cobalt binder however leads to a preferred amorphization of the coatings. The highest hardness is obtained for predominantly fcc structured WC0.67 (WC ceramic target), H = 40±1.7 GPa, exhibiting also an excellent intrinsic fracture toughness of KIC = 3.3±0.33 MPa·m1/2 obtained by micro-mechanical testing. Furthermore, the bonding nature of carbon is distinctly affected by the reactive carbon source, leading to more pronounced π-bonded carbon peak with increasing C2H2/Ar flow rates.
... The obtained values are summarized in Table 1. Based on DFT calculations presented in Ref. [16,32], the gained energy through the conversion into HfO 2 is highest for metallic Hf followed by HfC, HfB 2 , and HfN. This trend is also reflected in the activation energies obtained for Hf, HfN 1.5 , and HfB 2.3 , where HfN 1.5 obtains the high-est value and Hf the lowest -see Table 1. ...
Article
Full-text available
The influence of the non-metal species on the oxidation resistance of transition metal ceramic based thin films is still unclear. For this purpose, we thoroughly investigated the oxide scale formation of a metal (Hf), carbide (HfC0.96), nitride (HfN1.5), and boride (HfB2.3) coating grown by physical vapor deposition. The non-metal species decisively affect the onset temperature of oxidation, ranging between 550 °C for HfC0.96 to 840 °C for HfN1.5. HfB2.3 and HfN1.5 obtain the slowest oxide scale kinetic following a parabolic law with kp values of 4.97∙10⁻¹⁰ and 5.66∙10⁻¹¹ kg² m⁻⁴ s⁻¹ at 840 °C, respectively. A characteristic feature for the oxide scale on Hf coatings, is a columnar morphology and a substantial oxygen inward diffusion. HfC0.96 reveals an ineffective oxycarbide based scale, whereas HfN1.5 features a scale with globular HfO2 grains. HfB2.3 exhibits a layered scale with a porous boron rich region on top, followed by a highly dense and crystalline HfO2 beneath. Furthermore, HfB2.3 presents a hardness of 47.7 ± 2.7 GPa next to an exceptional low inward diffusion of oxygen during oxidation. This study showcases the strong influence of the non-metallic bonding partner despite the same metallic basis, as well as the huge potential for HfB2 based coatings also for oxidative environments.
... One of the approaches that allow enhancing properties of such materials is the formation of carbonitrides by replacing carbon atoms with nitrogen on the nonmetallic sublattice. Due to this substitution, the concentration of valence electrons changes, allowing hardness and fracture toughness to increase [26]. Also, it was predicted that HfC 0.5 N 0.38 may have an even higher melting point than HfC [27]. ...
Article
The oxidation resistance of the hafnium carbonitride (HfCN), HfCN+10 wt% SiC and HfCN+20 wt% SiC ceramics were investigated using thermogravimetric analysis (TGA) and static oxidation experiments at 1200 °C. TGA data showed that in the 1400 °C temperature range weight gain of the HfCN-SiC ceramics almost halted, while the HfCN sample continued to oxidize. Experiments at 1200 °C and further microstructure investigations at the ceramic/oxide interfaces showed formation of a dense HfSiO4 barrier layer with low oxygen diffusivity allowing increased oxidation resistance compared to the pure HfCN. A possible formation mechanism for the protective layer and its influence on oxidation was suggested. Also, it was shown that addition of the SiC powder to the HfCN allows decreasing consolidation temperature by 200 °C without mechanical properties loss. Hf(C,N) with 20 wt% SiC showed a high fracture toughness (5.1 MPa∙m1/2), high hardness (20.2 GPa), and lower density compared to the HfCN ceramics.
... This restricts their application and often leads to premature failure in use. Recent studies [33,34] in the field of non-reactively sputtered transition metal carbonitrides highlighted an enhancement of fracture characteristics through a non-reactively deposition route. ...
Article
Full-text available
In the field of hard protective coatings, nano-crystalline Ti-B-N films are of great importance due to the adjustable microstructure and mechanical properties through their B content. Here, we systematically study this influence of B on Ti-B-N during reactive as well as non-reactive DC magnetron sputtering. The different deposition routes allow for an additional, very effective key parameter to modify bond characteristics and microstructure. Plasma analysis by mass spectroscopy reveals that for comparable amounts of Ti⁺, Ti²⁺, Ar⁺, and Ar²⁺ ions, the count of N⁺ ions is about 2 orders of magnitude lower during non-reactive sputtering. But for the latter, the N⁺/N2⁺ ratio is close to 1, whereas during reactive sputtering this ratio is only 0.1. This may explain why during reactive deposition of Ti-B-N, the B-N bonds dominate (as suggested by X-ray photoelectron spectroscopy), whereas the B-B and Ti-B bonds dominate for non-reactively prepared Ti-B-N. Chemically, reactively versus non-reactively sputtered Ti-B-N coatings follow the TiN-BN versus TiN-TiB2 tie line, respectively. Detailed X-ray diffraction and transmission electron microscopy studies reveal, that up to 10 at% B can be dissolved in the fcc-TiN lattice when prepared by non-reactive sputtering, whereas already for a B content of 4 at% a BN-rich boundary phase forms when reactively sputtered. Thus, we could not only observe a higher hardness (35 GPa instead of 25 GPa) as well as a higher indentation modulus (480 GPa instead of 260 GPa), but also a higher fracture energy (0.016 instead of 0.009 J/m during cube-corner indentations) for Ti-B-N coatings with 10 at% B, when prepared non-reactively.
Article
The exceptional mechanical properties of transition metal carbide coatings are known to be governed by the carbon content and its morphological distribution. Here, we verify the influence of the target peak power density on the chemical composition, microstructure, and mechanical properties of NbCx coatings grown by non-reactive high-power impulse magnetron sputtering (HiPIMS). By tuning the pulse parameters, the power density can be increased from 0.11 to 1.48 kW/cm2 leading to a decrease in the C/Nb ratio from 1.52 to 0.99 within the films – proven by combined elastic backscattering and time-of-flight elastic recoil detection analysis. This decrease in the C/Nb ratio is accompanied by microstructural changes from nanocomposite morphologies with an average grain size of 6.6 ± 2.5 nm at 0.13 kW/cm2 into more columnar structures with an average column width of 65.2 ± 18.7 nm at 1.48 kW/cm2. Independent from the C/Nb ratio, all films exhibit a single face-centered cubic structure. The mechanical properties correlate with the enhanced growth behavior dominated by ions at higher peak power densities and the varied C/Nb ratios. A maximum in hardness and fracture toughness of H = 38.7 ± 3.6 GPa and KIc = 2.78 ± 0.13 MPa∙m1/2 (at 3.2 GPa residual compressive stress), is obtained for the nearly stoichiometric NbC coating exhibiting C/Nb ratio of 1.06.
Article
In this work chemical vapor deposited (CVD) coatings of (Tix,W1-x)Ny from TiCl4, WF6, NH3 and Ar were investigated. This coating material has previously been deposited using other vacuum techniques but no publication has so far demonstrated CVD of (Tix,W1-x)Ny. The studied (Tix,W1-x)Ny coatings had a metallic molar ratio (Ti:W) close to 2:1 and 1:1, and were slightly over-stoichiometric with regard to N. The coatings appeared homogeneous and crystallised in a rock salt structure on an α-Al2O3 substrate. The cell parameter varied between 4.16 and 4.23 Å as a function of the deposition conditions, ranging from a pure TiNx to a pure WNx coating. The texture in the normal direction was 〈100〉 for the TiNx and (Tix,W1-x)Ny coatings and 〈111〉 for WNx. Electron backscattered diffraction (EBSD) results showed that a strong correlation to the substrate existed but random in-plane orientation was also present. The microstructure showed columnar grains with well defined facets growing. Adding a mixture of TiCl4 and WF6 to produce (Tix,W1-x)Ny did increase the grain size significantly when compared to the case when only one metal precursor was present. The down-stream thickness profile, using only WF6 and NH3, displayed mass transport control behaviour, with the coating thickness converging to zero within the deposition zone. Using only TiCl4 on the other hand showed a uniform deposition profile, the signs of a surface kinetics controlled process.
Article
Full-text available
Motivated by density functional theory (DFT)-derived ductility indicators for face centred cubic (fcc, rocksalt) structured TiN/MoN0.5 superlattices and Ti0.5Mo0.5N0.75 solid solutions, TiN/MoNy superlattice (SL) thin films with bilayer periods Λ of 2.4, 3.9, 6.6, 9.9, and 23.0 nm and corresponding solid solutions were developed by DC reactive magnetron sputtering. These SLs allow for improved hardness H and critical fracture toughness KIC, with both peaking at the same bilayer period Λ of 9.9 nm (where the MoN0.5 layers crystallize with the ordered β-Mo2N phase); H = 34.8±1.6 GPa and KIC = 4.1±0.2 MPa√m. The correspondingly prepared fcc-Ti0.5Mo0.5N0.77 solid solution has H = 31.4±1.5 GPa and KIC = 3.3±0.2 MPa√m. Thus, especially the fracture toughness shows a significant superlattice effect. This is suggested by DFT – by the increase of the Cauchy pressure from -19 to +20 GPa for the 001-direction (while that in the 100-direction remained high, above 83 GPa) upon increasing Λ from 3 to 4 nm. Together, experimental and computational investigations prove the importance of optimized bilayer periods for highest strength and fracture toughness, as well as optimized N-content for the solid solutions.
Article
Full-text available
With a considerable amount of commonly used material systems consisting of individual, rather confined layers, the question for mechanical behaviour of their individual interfaces arises. Especially, when considering varying interfacial structures as a result of the processing environment. Furthermore, the interaction between pronounced plasticity and fracture processes can lead to challenges with regards to separation between sole interface- or bulk properties. The present work investigates the interfacial fracture characteristic of a WTi-Cu sytem commonly found in the microelectronics industry as a heterogeneous model material with pronounced plasticity in the Cu phase. To study this behaviour on a rather limited scale (<6 µm), microcantilever experiments were conducted and evaluated using a continuous J-Δa curve evaluation scheme with classical elastic plastic considerations in mind. A change in interface chemistry, resulting from air exposure between processing steps was probed and found to show distinct crack propagation along the interface opposed to crack tip blunting as found in the vacuum processed sample. Complementary density functional theory calculations also showed a strong reduction of interface cohesion upon oxygen accumulation and a model framework based on classical dislocation plasticity considerations revealed the transition from plasticity to fracture processes to be a result of shielding and following change in mode mixity.
Article
In this study, we report high entropy carbides synthesis with reactive bipolar high‐power impulse magnetron sputtering (HiPIMS). Uncontrolled microstructure and stoichiometry development with reactive gas flow rate are major limitations of conventional DC and RF magnetron sputtering of multicomponent carbides. With HiPIMS these chemically disordered crystals structurally and compositionally transform from a carbon deficient metallic (C/M<1), to a stoichiometric ceramic zone (C/M∼1), and to a nanocomposite embodiment (C/M>1), as a function of the carbon content during HiPIMS deposition. X‐ray diffraction, X‐ray photo electron spectroscopy, Raman spectroscopy, scanning electron microscopy and nanoindentation hardness measurements are combined to demonstrate the three regions of synthesis domain. HiPIMS provides access to metallic, ceramic and composite carbides with great control over the microstructure and stoichiometry which is elusive in case of conventional DC and RF magnetron sputtering. Notably, the stoichiometric ceramic zone maintains a constant carbon to metal ratio (C/M∼1) over an extended amount of methane flow before transitioning to a nanocomposite microstructure(C/M>1). The transition zone breadth depends on materials affinity for carbon that correlates with valence electron concentration (VEC). As such, synthesis conditions for new high entropy carbides can be understood and predicted based on VEC.
Article
Building nanocrystalline/amorphous biphase nanostructure has recently emerged as a new strengthening-toughening route that can combine the strengthening benefits of nanocrystallinity and amorphization. Therefore, it is required exploring on the strain rate-dependent deformation behavior with above different nanostructure in certain system. Herein, three typical nanostructures: polycrystalline, uniform nanocrystalline/amorphous core-shell nanostructure and amorphous matrix with nanoclusters, have been achieved in the sputtered Ti2AlNb films by individually regulating the sputtering bias voltage (Ubias). Note that a high Ubias can promote formation of Nb−rich amorphous tissues via Al preferential resputtering and Nb segregation. Moreover, microstructural evolution and strain rate-dependent hardness and deformation behavior were further explored. Firstly, at a low Ubias (−40V), Ti2AlNb film exhibited polycrystalline character; it yielded a relatively low hardness (~10.0 GPa) but a high strain sensitivity coefficient of 0.1505. Subsequently, the novel nanocrystalline/amorphous core-shell nanostructure consisting of Ti2AlNb nanocrystalline cores uniformly encapsulated by thin amorphous shells was achieved at a higher Ubias (−120V); this nanostructure provided a remarkable hardness enhancement (~15.2 GPa) without sacrificing its ductility and intermediate strain sensitivity coefficient of 0.1382. Finally, a transition to amorphous matrix with nanoclusters occurred with further increasing Ubias to −200V, thus yielding slight decrease in hardness (~12.5 GPa) and a minimum strain sensitivity coefficient of 0.0915 when shear-band deformation was activated.
Article
Full-text available
Recently, amorphous Hf-B-Si-C-N coatings found to demonstrate superior high-temperature oxidation resistance. The microstructure evolution of two coatings, Hf7B23Si22C6N40 and Hf6B21Si19C4N47, annealed to 1500 °C in air is investigated to understand their high oxidation resistance. The annealed coatings develop a two-layered structure comprising of the original as-deposited film followed by an oxidized layer. In both films, the oxidized layer possesses the same microstructure with HfO2 nanoparticles dispersed in an amorphous SiOx-based matrix. The bottom layer in the Hf6B21Si19C4N47 coating remains amorphous after annealing while Hf7B23Si22C6N40 recrystallized partially showing a nanocrystalline structure of HfB2 and HfN nanoparticles separated by h-Si3N4 and h-BN boundaries. The HfB2 and HfN nanostructures form a sandwich structure with a HfB2 strip being atomically coherent to HfN skins via (111)-Hf monolayers. In spite of the different bottom layer structure, the oxidized/bottom layer interface of both films was found to exhibit a similar microstructure with a fine distribution of HfO2 nanoparticles surrounded by SiO2 quartz boundaries. The high-temperature oxidation resistance of both films is attributed to the particular evolving microstructure consisting of HfO2 nanoparticles within a dense SiOx-based matrix and quartz SiO2 in front of the oxidized/bottom layer interface acting as a barrier for oxygen and thermal diffusion.
Article
Full-text available
Low-temperature epitaxial growth of refractory transition-metal nitride thin films by means of physical vapor deposition has been a recurring theme in advanced thin-film technology for several years. In the present study, 150-nm-thick epitaxial HfN layers are grown on MgO(001) by reactive high-power impulse magnetron sputtering (HiPIMS) with no external substrate heating. Maximum film-growth temperatures Ts due to plasma heating range from 70 to 150 °C, corresponding to Ts/Tm = 0.10–0.12 (in which Tm is the HfN melting point in K). During HiPIMS, gas and sputtered metal-ion fluxes incident at the growing film surface are separated in time due to strong gas rarefaction and the transition to a metal-ion-dominated plasma. In the present experiments, a negative bias of 100 V is applied to the substrate, either continuously during the entire deposition or synchronized with the metal-rich portion of the ion flux. Two different sputtering-gas mixtures, Ar/N2 and Kr/N2, are employed in order to probe effects associated with the noble-gas mass and ionization potential. The combination of x-ray diffraction, high-resolution reciprocal-lattice maps, and high-resolution cross-sectional transmission electron microscopy analyses establishes that all HfN films have a cube-on-cube orientational relationship with the substrate, i.e., [001]HfN||[001]MgO and (100)HfN||(100)MgO. Layers grown with a continuous substrate bias, in either Ar/N2 or Kr/N2, exhibit a relatively high mosaicity and a high concentration of trapped inert gas. In distinct contrast, layers grown in Kr/N2 with the substrate bias synchronized to the metal-ion-rich portion of HiPIMS pulses have much lower mosaicity, no measurable inert-gas incorporation, and a hardness of 25.7 GPa, in good agreement with the results for epitaxial HfN(001) layers grown at Ts = 650 °C (Ts/Tm = 0.26). The room-temperature film resistivity is 70 μΩ cm, which is 3.2–10 times lower than reported values for polycrystalline-HfN layers grown at Ts = 400 °C.
Article
Over the past decades, enormous effort has been dedicated to enhancing the hardness of refractory ceramic materials. Typically, however, an increase in hardness is accompanied by an increase in brittleness, which can result in intergranular decohesion when materials are exposed to high stresses. In order to avoid brittle failure, in addition to providing high strength, films should also be ductile, i.e., tough. However, fundamental progress in obtaining hard-yet-ductile ceramics has been slow since most toughening approaches are based on empirical trial-and-error methods focusing on increasing the strength and ductility extrinsically, with a limited focus on understanding thin-film toughness as an inherent physical property of the material. Thus, electronic structure investigations focusing on the origins of ductility vs. brittleness are essential in understanding the physics behind obtaining both high strength and high plastic strain in ceramics films. Here, we review recent progress in experimental validation of density functional theory predictions on toughness enhancement in hard ceramic films, by increasing the valence electron concentration, using examples from the V1-xWxN and V1-xMoxN alloy systems.
Article
First-principles calculations are employed to determine the mechanical properties of rock-salt structure binary and ternary transition metal nitrides, carbides, and carbonitrides from groups 4 to 12, predicting a unified indicator for mechanical properties: the valence electron concentration (VEC). Pugh's and Poisson's ratios indicate an increasing ductility with increasing VEC, with a brittle-to-ductile transition at a critical VEC = 10. The calculated C44 of carbonitrides and ternary nitrides monotonically decreases from 164 ± 12 GPa at VEC = 8 to −39 ± 46 GPa at VEC = 11, indicating a transition to mechanical instability at VEC = 10.6. Similarly, the average isotropic elastic modulus decreases slightly from 420 GPa for VEC = 8 to 388 GPa for VEC = 10, but then steeply to −98 GPa for VEC = 11, while the corresponding hardness decreases from 25 to 12 to 2 GPa. The overall softening with increasing VEC is attributed to the increasing electron density in d–t2g orbitals, which overlap upon shear and cause a decrease in C44. Phonon dispersion curves, calculated at 0 K for binary nitrides and carbides, exhibit imaginary frequencies for VEC ≥10, indicating a dynamical stability-to-instability transition between VEC = 9 and 10, which is smaller than the critical VEC = 10.6 for the mechanical stability-instability transition. In addition, mechanical stability is increased by magnetic ordering but decreased when accounting for on-site Coulomb repulsion, while temperature and vacancies cause a reduction in the magnitude of C44 for both stable and unstable compounds, likely leading to an increase in the critical VEC for the stability-instability transition. The overall results indicate a narrow region between VEC = 9 and 10 where rocksalt carbonitrides are ductile but also exhibit a high hardness, mechanical and dynamical stability, and therefore are expected to exhibit the highest toughness.
Article
Abstract Using nonreactive sputter deposition allows the preparation of single-phase fcc structured TaCy thin films over a wide compositional range with y between 0.63 and 1.04. Among this composition range, the C-deficient TaC0.78 exhibits the highest as deposited hardness of 43.4 ± 0.65 GPa combined with the highest thermal stability. Even after vacuum annealing to 2400 °C, no vacancy-ordered or faulted Ta-C based phases can be detected. The stabilization of carbon deficient fcc structured TaCy near y of about 0.75, revealed the decisive character of vacancy engineered thin films materials for ultra-high temperature applications.
Article
In the paper, it is stated that the HfCxN1-x coating was deposited by magnetron sputtering on the AISI 316L substrate materials at various substrate temperatures (Ts). The results of XRD and TEM indicated that the HfCxN1-x coatings were almost amorphous when the Ts were below 400 °C, whereas the nanostructured Hf2CN was formed at the Ts of 400 °C. The Hf2CN nanocrystalline formation resulted in a significant improvement in the mechanical properties of HfCxN1-x coatings. The electrochemical tests demonstrated that the stable HfCxN1-x coatings improved the corrosion resistance of the AISI 316L. The platelet adhesion tests and hemolysis ratio tests indicated that the HfCxN1-x coatings presented an improved hemocompatibility when the Ts varied from 25 °C to temperatures exceeding 300 °C, which might have occurred due to the surface topography, the interfacial tension and the γsd/γsp ratio.
Article
Focused ion beam machined microcantilevers are frequently used for fracture mechanics analysis of inhomogeneous solids at the micrometer scale. A finite element method study about the influence of the pre-crack geometry on the apparent fracture toughness is provided. We discuss the influence of material bridges and the effect of rounded pre-crack corners when two dimensional models are employed to evaluate the fracture toughness. We conclude with a guideline for introducing an optimized pre-crack.
Article
Transition metal carbides are known for their exceptional thermal stability and mechanical properties, notably governed by the carbon content and the prevalent vacancies on the non-metallic sublattice. However, when using reactive deposition techniques, the formation of amorphous C-containing phases is often observed. Here, we show that non-reactive magnetron sputtering of HfC0.89 or TaC0.97 targets lead to fully crystalline coatings. Their C content depends on the target-to-substrate alignment and globally increases from HfC0.66 to HfC0.76 and from TaC0.69 to TaC0.75 with increasing bias potential from floating to − 100 V, respectively, when using a substrate temperature Tsub of 500 °C. Increasing Tsub to 700 °C leads to variations from TaC0.71 to TaC0.81.
Article
This paper identifies gaps in the present state of knowledge and describes emerging research directions for ultra-high temperature ceramics. Borides, carbides, and nitrides of early transition metals such as Zr, Hf, Nb, and Ta have the highest melting points of any known compounds, making them suitable for use in extreme environments. Studies of synthesis, processing, densification, thermal properties, mechanical behavior, and oxidation of ultra-high temperature ceramics have generated a substantial base of knowledge, but left unanswered questions. Emerging research directions include testing/characterization in extreme environments, composites, computational studies, and new materials.
Article
A dedicated detector system for heavy ion elastic recoil detection analysis at the Tandem Laboratory of Uppsala University is presented. Benefits of combining a time-of-flight measurement with a segmented anode gas ionization chamber are demonstrated. The capability of ion species identification is improved with the present system, compared to that obtained when using a single solid state silicon detector for the full ion energy signal. The system enables separation of light elements, up to Neon, based on atomic number while signals from heavy elements such as molybdenum and tungsten are separated based on mass, to a sample depth on the order of 1 μm. The performance of the system is discussed and a selection of material analysis applications is given. Plasma-facing materials from fusion experiments, in particular metal mirrors, are used as a main example for the discussion. Marker experiments using nitrogen-15 or oxygen-18 are specific cases for which the described improved species separation and sensitivity are required. Resilience to radiation damage and significantly improved energy resolution for heavy elements at low energies are additional benefits of the gas ionization chamber over a solid state detector based system.