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Influence of Ta on the oxidation resistance of WB 2−z coatings

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Ternary W 1−x Ta x B 2−z is a promising protective coating material possessing enhanced ductile character and phase stability compared to closely related binaries. Here, the oxidation resistance of W 1−x Ta x B 2−z thin films was experimentally investigated at temperatures up to 700 °C. Ta alloying in sputter deposited WB 2−z coatings led to decelerated oxide scale growth and a changed growth mode from paralinear to a more linear (but retarded) behavior with increasing Ta content. The corresponding rate constants decrease from k* p = 6.3 ⋅ 10 −4 µm 2 /s for WB 2−z , to k* p = 1.1 ⋅ 10 −4 µm 2 /s for W 0.66 Ta 0.34 B 2−z as well as k l = 2.6 ⋅ 10 −5 µm/s for TaB 2−z , underlined by decreasing scale thicknesses ranging from 1170 nm (WB 2−z), over 610 nm (W 0.66 Ta 0.34 B 2−z) to 320 nm (TaB 2−z) after 10 min at 700 °C. Dense and adherent scales exhibit an increased tantalum content (columnar oxides), which suppresses the volatile character of tungsten-rich as well as boron oxides, hence being a key-factor for enhanced oxidation resistance. Thus, adding Ta (in the range of x = 0.2-0.3) to α-structured WB 2−z does not only positively influence the ductile character and thermal stability but also drastically increases the oxidation resistance.
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Journal of Alloys and Compounds
journal homepage: www.elsevier.com/locate/jalcom
Influence of Ta on the oxidation resistance of WB
2−z
coatings
C. Fuger
a,
, B. Schwartz
a
, T. Wojcik
a,b
, V. Moraes
b
, M. Weiss
c
, A. Limbeck
c
, C.A. Macauley
d,e
,
O. Hunold
f
, P. Polcik
g
, D. Primetzhofer
h
, P. Felfer
d
, P.H. Mayrhofer
b
, H. Riedl
a,b
a
Christian Doppler Laboratory for Surface Engineering of high-performance Components, TU Wien, Austria
b
Institute of Materials Science and Technology, TU Wien, A-1060 Wien, Austria
c
Institute of Chemical Technologies and Analytics, TU Wien, Vienna, Austria
d
Department of Materials Science, Friedrich-Alexander-Universität Erlangen-Nürnberg, Germany
e
Interdisciplinary Center for Nanostructured Films (IZNF), 91058 Erlangen, Germany
f
Oerlikon Balzers, Oerlikon Surface Solutions AG, 9496 Balzers, Liechtenstein
g
Plansee Composite Materials GmbH, D-86983 Lechbruck am See, Germany
h
Department of Physics and Astronomy, Uppsala University, SE-75120 Uppsala, Sweden
article info
Article history:
Received 18 September 2020
Received in revised form 22 November 2020
Accepted 24 November 2020
Available online xxxx
Keywords:
Oxidation resistance
Transition metal diborides
Scale growth
Diborides
abstract
Ternary W
1−x
Ta
x
B
2−z
is a promising protective coating material possessing enhanced ductile character and
phase stability compared to closely related binaries. Here, the oxidation resistance of W
1−x
Ta
x
B
2−z
thin films
was experimentally investigated at temperatures up to 700 °C. Ta alloying in sputter deposited WB
2−z
coatings led to decelerated oxide scale growth and a changed growth mode from paralinear to a more linear
(but retarded) behavior with increasing Ta content. The corresponding rate constants decrease from
k*
p
= 6.3 10
−4
µm
2
/s for WB
2−z
, to
k*
p
= 1.1 10
−4
µm
2
/s for W
0.66
Ta
0.34
B
2−z
as well as k
l
= 2.6 10
−5
µm/s for
TaB
2−z
, underlined by decreasing scale thicknesses ranging from 1170 nm (WB
2−z
), over 610 nm
(W
0.66
Ta
0.34
B
2−z
) to 320 nm (TaB
2−z
) after 10 min at 700 °C. Dense and adherent scales exhibit an increased
tantalum content (columnar oxides), which suppresses the volatile character of tungsten-rich as well as
boron oxides, hence being a key-factor for enhanced oxidation resistance. Thus, adding Ta (in the range of
x = 0.2–0.3) to α-structured WB
2−z
does not only positively influence the ductile character and thermal
stability but also drastically increases the oxidation resistance.
© 2020 The Author(s). Published by Elsevier B.V.
CC_BY_4.0
1. Introduction
Various industrial applications, e.g. in aerospace industry, avia-
tion, or energy production, demand complex applied materials re-
quirements to increase lifetime and thus, enhance environmental
sustainability. Since bulk materials occasionally cannot fulfill all
these demands, surface engineering using physical vapor deposition
(PVD) deposited thin films is a well-established method enhancing
specific properties. Based on the most common applications, tran-
sition metal ceramics, particularly transition metal nitrides such as
TiN, CrN, or Ti
1−x
Al
x
N have been explored in depth over decades
[1–3]. However, the material class of borides is a valuable alternative
in new applications. Especially, transition metal borides (TMBs) ex-
hibit a tremendous potential to be applied in various fields ranging
from wear and corrosion resistant coatings, to superconductive thin
films, or as superhard and extremely stable protective layers in
diverse engineering applications [4–6]. One very interesting re-
presentative of the TMBs is tungsten diboride, which exhibits spe-
cific mechanical properties with respect to the generally limited
fracture tolerance of such ceramic compounds [7–9]. Previous stu-
dies on diborides revealed that transition metal diborides tend to
crystalize in two different hexagonal structures, α- (AlB
2
, space group
191, P6/mmm) and ω-prototype (W
2
B
5−x
, space group 194, P6
3
/mmc),
respectively [10,11]. By synthesizing WB
2−z
coatings using PVD
techniques, the films crystallize in the metastable α-phase rather
than in the thermodynamically preferred ω-structure. As structural
defects play a crucial role in the stabilization of α-structured WB
2−z
,
the impact on the mechanical properties, especially fracture re-
sistance, is still not fully described. However, in theoretical as well as
experimental investigations (by free-standing cantilever tests), α-
WB
2−z
exhibits a highly ductile behavior with K
IC
values of around
3.7 ± 0.3 MPa√m. In comparison, ZrB
2−z
and Zr
1−x
Ta
x
B
2−z
films obtain
values around from 3.5 to 5.5 MPa√m during cube corner indenta-
tion, which typically overestimates fracture characteristics (com-
pared to cantilever bending) by at least 2.0 MPa√m [12–14].
However, one weak point of α-structured WB
2−z
is the thermal
https://doi.org/10.1016/j.jallcom.2020.158121
0925-8388/© 2020 The Author(s). Published by Elsevier B.V.
CC_BY_4.0
Corresponding author.
E-mail address: christoph.fuger@tuwien.ac.at (C. Fuger).
Journal of Alloys and Compounds xxx (xxxx) xxx
Please cite this article as: C. Fuger, B. Schwartz, T. Wojcik et al., Influence of Ta on the oxidation resistance of WB
2−z
coatings, Journal of
Alloys and Compounds, https://doi.org/10.1016/j.jallcom.2020.158121i
stability, as the decomposition to its stable ω-modification sets in
relatively early – at around 800 °C. Furthermore, in addition to the
limited phase stability in inert atmospheres, the formation of volatile
tungsten oxides occurs at even lower temperatures and thus strongly
limits the wide usage of α-structured WB
2−z
[15–17]. With respect to
phase transition between α and ω, the addition of Ta has been proven
to shift the starting point for decomposition to increased tempera-
tures. The resulting ternary W
1−x
Ta
x
B
2−z
with x ~ 0.25 leads to an
enhancement of the decomposition temperature (in inert atmo-
sphere) by 600 °C to a maximum of T
dec
= 1400 °C [18]. Small addi-
tions of Ta also slightly enhance the fracture toughness to
3.8 ± 0.5 MPa√m at x ~ 0.15, while obtaining hardness values above
40 GPa [12,18]. All these points suggest W
1−x
Ta
x
B
2−z
as a promising
protective coating system. However, a clear study of the influence of
Ta on the oxidation resistance is still missing to complete the profile
of properties.
Therefore, based on the above-mentioned results, the focus of
this study is to investigate in detail the influence of Ta on the oxi-
dation resistance of sputter deposited W
1−x
Ta
x
B
2−z
thin films.
2. Experimental
All WB
2−z
, W
1−x
Ta
x
B
2−z
, and TaB
2−z
coatings were deposited by DC
magnetron sputtering using an in-house developed deposition
system with two confocal arranged cathodes, for details see [19]. The
ultra-high vacuum coating facility was equipped with two 6-inch
TaB
2
and W
2
B
5
– powder-metallurgically produced targets obtaining
a purity of at least 99.6% (Plansee Composite Materials GmbH). Prior
to the depositions, the silicon (20 × 7 × 0.33 mm
3
, 100-oriented) and
single crystalline sapphire (10 × 10 × 0.5 mm
3
,
012-oriented) sub-
strates were ultrasonically pre-cleaned in acetone and ethanol for
5 min each. Subsequently, they were mounted on a rotating substrate
holder, and heated up in the chamber (base pressure below 4 10
−4
Pa)
to the deposition temperature, T
dep
, of 500 °C (corresponding to
300 ± 15 °C on the substrate surface) for at least 20 min. Further-
more, the substrates were etched in argon atmosphere (p
etch
= 6 Pa)
for 10 min applying a substrate potential of −500 V. The deposition
process itself was carried out in pure argon atmosphere at a de-
position pressure of 0.4 Pa. Both targets were powered by a Solvix
HIP
3
generator used in DC mode – controlling the applied target
current, I
target
, of maximum 4.2 A in total for both cathodes. To vary
the chemical composition of the deposited films (dividing the full
compositional range in steps of about x ~ 0.20) the applied current
was changed successively for both target materials. The deposition
time was varied between 60 and 85 min due to a decreased sputter
rate for TaB
2
with respect to W
2
B
5
. To gain homogeneous composi-
tions over all substrates, the substrate holder was rotated with a
frequency, f
rot
, of 0.25 Hz. The film growth was also supported by
applying a bias potential of −50 V.
All oxidation tests were carried out in a standard chamber fur-
nace in ambient air. To gain relatively flat and defined oxide scales,
all oxidation tests have been performed on sapphire substrates. The
kinetic behavior of the oxide scale formation was analyzed through
varying the exposure time (t
ox
= 1, 10, 100, and 1000 min) and tem-
perature (T
ox
= 500, 600, and 700 °C), respectively. The temperatures
and time periods were selected based on oxidation pretests. In detail,
all compositions have been fully oxidized after 1000 min at 700 °C
and therefore the oxidation temperature was not further increased.
After quantifying the oxide scale thicknesses by means of scanning
electron microscopy (SEM) investigations (FEI Quanta 200 FEG-SEM)
within fracture cross-sections, the oxidation rate constants were
calculated by linear regression of squared oxide thickness values as a
function of oxidation time.
As a proper chemical quantification of light elements in combi-
nation with heavy ones e.g. W
1−x
Ta
x
B
2−z
, presents a certain challenge
[18], different chemical analysis methods were utilized to obtain best
results. The elemental composition of all as deposited thin films on Si
substrates was analyzed by liquid inductively coupled plasma optical
emission spectroscopy (ICP-OES). Liquid ICP-OES measurements
were carried out on an iCAP 6500 RAD (Thermo Fisher Scientific,
USA), with an ASX-520 autosampler (CETAC Technologies, USA) using
a HF resistant sample introduction kit, consisting of a Miramist
nebulizer (Burger Research, USA), a PTFE spray chamber and a
ceramic injector tube. All W
1−x
Ta
x
B
2−z
coatings were acid digested
with the method presented and validated in [20]. Samples were
broken into pieces of about 3 × 3 mm
2
and the thin film with the
substrate was dissolved in a mixture of 1 mL HNO
3
and 0.25 mL HF.
After a reaction time of 15 min at a temperature of 60 °C, the thin
films including the substrates were completely dissolved. Derived
sample digests were diluted to a final volume of 20 mL with a mix-
ture of 3% HNO
3
and 0.3% HF. Quantification was done via external
calibration using matrix adjusted standards - for further details see
also [20–22].
For a 3D elemental distribution, a selected coating in the as de-
posited state was analyzed by atom probe tomography (APT). The
APT analysis was carried out on a CAMECA LEAP 4000X HR in pulsed
laser mode. This instrument is equipped with a 355 nm UV laser with
a spot size of ~2 µm and a reflectron lens resulting in a detection
efficiency of ~37%. The experiments were done with a laser pulse
energy of 50 pJ at a target evaporation rate of 1%. The as-deposited
samples were extracted using a keyhole technique [23], which places
the analysis axis in the growth direction of the film. Moreover, the
composition of selected as deposited samples as well as composi-
tional depth profiles of the oxide layers were obtained by time of
flight elastic recoil detection analysis (TOF-ERDA) using a 36 MeV I
8+
ion beam and detecting recoils in a detection angle of 45° with re-
spect to the primary beam. Details on the employed detection
system can be found in Ström et al. [24]. The expected systematic
uncertainties for light elements such as B are found on a level of ±
1 at% for relative measurements free from standards, mainly due to
uncertainties in the specific energy loss of the recoiling particles. For
absolute measurements (free from standards), systematic un-
certainties for light elements (e.g. B or C) are expected to be on a
level of 5–10% of the detected concentration. A detailed description
of sources and consequences of systematic uncertainties are dis-
cussed in more detail by Arvizu et al. [25] and Zhang et al. [26].
Depth profiles were established using the CONTES software
package [27].
The coating morphology before and after oxidation – was
analyzed in cross sectional view by a FEI Quanta 250 FEG-SEM,
equipped with a field emission gun (acceleration voltage used,
10 kV). In addition, the chemical composition of the formed scales
and unaffected coatings was analyzed by energy dispersive X-Ray
spectroscopy (EDAX EDS detector, 15 kV acceleration voltage).
Selected samples were also surveyed in cross sectional view by
transmission electron microscopy (TEM FEI TECNAI, G20, accelera-
tion voltage of 200 kV). Detailed structural information of the formed
oxide scales was gained by selected area electron diffraction (SAED)
analysis. The sample preparation was done by focused ion beam (FIB,
Quanta 200 3D Dual Beam), applying standard lift out techni-
ques [28].
3. Results and discussion
3.1. Chemical composition and phase constitution as deposited
After sputter depositing the films, the chemical compositions of
all coatings have been analyzed by ICP-OES. Fig. 1 shows the boron
content (in at%) of all deposited coatings as a function of the Ta
fraction at the metal sublattice, x. The blue, half-filled triangles are
indicating the chemical compositions of the films synthesized within
this study, while the green open squares are highlighting the
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
2
elemental fraction of W
1−x
Ta
x
B
2−z
coatings (obtained by TOF-ERDA)
prepared in a previous study [18]. For a further comparison between
different analysis techniques as well as quality and reproducibility of
the depositions, the average W, Ta (in terms of metal sublattice oc-
cupation, x) and B concentration of a surveyed atom probe tip is also
given in Fig. 1, see red open square (selected coating originates from
study [18] analyzed volume is depicted in Fig. 2). Both studies (blue
and green data set refer to this study and Ref. [18]) exhibit the similar
tendency of decreasing boron content with increasing amount of
tantalum, highlighting the strong affinity to form either sub-stoi-
chiometric structures with boron and/or Schottky defects as well as
multi-phased coatings. Nevertheless, the decreasing B contents with
increasing Ta are a strong indication, that the phase formation of α-
structured W
1−x
Ta
x
B
2−z
films at high Ta fractions is limited and di-
rected towards dual phase structures (α-W
1−x
Ta
x
B
2−z
+ α-TaB
2−z
or an
orthorhombic o-TaB phase) – as also suggested in [12]. By con-
sidering all chemical and structural data (X-Ray diffraction [XRD]
pattern for the full compositional range are presented in the Ap-
pendix), we can assume that there is a more or less ideal solid so-
lution of α-W
1−x
Ta
x
B
2−z
up to about x = 0.20 (on the metal sublattice)
and dual phased morphologies containing α-W
1−x
Ta
x
B
2−z
as well as
α-TaB
2−z
rich domains at Ta concentrations x > 0.40. Furthermore, at
Ta x > 0.85 an additional o-TaB phase is occurring. The gray shaded
areas between x = 0.20–0.40 as well as 0.75–0.85 exhibit the transi-
tion zones of upcoming TaB
2
/TaB rich domains and the change be-
tween α-dominated structures and the additional appearance of a o-
TaB phase, respectively.
To reveal a more detailed picture about the 3D atomic composi-
tion in the so-called α-TaB
2
transition zone (x ≥ 0.26, Ta submetal
occupation), a detailed ATP analysis of the W
0.74
Ta
0.26
B
1.89
(given
composition obtained by TOF-ERDA) coating was conducted (see
Fig. 2a). The inhomogeneous distribution of tantalum in the growth
direction, is shown clearly in the one-dimensional concentration
profile (see Fig. 2b). Both, Ta and W show an inversely oscillating
concentration going from ~5 at% to ~15 at% for Ta (x = 0.12–0.37 Ta
occupation on the metal sublattice) and from ~35 at% to ~25 at% for
W (x = 0.88–0.63 W sublattice occupation). This behavior could be
due to, (i) a limited solubility of Ta in the α-W
1−x
Ta
x
B
2−z
structure,
indicating the occurrence of a Ta rich phase and/or (ii) a synthesis
related effect through the experimental setup of co-sputtering a
WB
2−z
and TaB
2−z
targets, respectively. By taking into account the
Fig. 1. Chemical composition obtained by ICP-OES of all coatings deposited within this study (blue data points). The green data points indicate the chemical composition
evaluated by TOF-ERDA of W
1−x
Ta
x
B
2−z
thin films taken from [18] – except TaB
2−z
which was analyzed within this study. During APT the overall elemental concentration of the
inspected tip was evaluated and is indicated in red. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 2. Elemental distribution of Ta obtained by atom probe tomography of
W
0.74
Ta
0.26
B
2−z
coating (a), showing a horizontally layered structure also confirmed by
the one-dimensional concentration profile, depicted in (b). The layered structure is
characterized by an alternating W/Ta content, highlighted by the image detail (c).
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
3
rotational frequency of the substrate holder (0.25 Hz) and the de-
position rate of both binary systems (~70 nm/min for WB
2−z
com-
pared to ~ 50 nm/min for TaB
2−z
) result in an overall deposition rate
of ~ 5 nm/s, which fits to the periodicity in Fig. 2b. Assuming that the
substrate holder is not fully covered from both, the WB
2−z
and TaB
2−z
sputter plumes, there are areas with increased WB
2−z
/TaB
2−z
con-
centration, dependent on the progress of rotation. Therefore, the
layered structure, depicted in Fig. 2, is suggested to be an unintended
variation caused by the co-sputtering process [29]. Nevertheless, a
limited solubility of Ta in the α-W
1−x
Ta
x
B
2−z
structure cannot be
ruled out, as the W-depleted sections reveal x ~ 0.37 Ta on the metal
sublattice (overall ~15 at% Ta). A more detailed view of the alter-
nating W/Ta rich sections is also given in the cutout, depicted in
Fig. 2c. Due to reasons of clarity, W atoms are not visualized in Fig. 2.
These detailed analyses are also in good agreement with the dif-
fraction results, which suggests that samples in the α-TaB
2
transition
zone (x ≥ 0.26) form imperfect as-deposited solid solutions when
deposited using co-sputtering. Nevertheless, the occurrence of
multi-phased coatings containing α-W
1−x
Ta
x
B
2−z
, α-TaB
2−z
, as well as
o-TaB is mostly related to a lack of boron and the high tendency to
form stoichiometric o-TaB. (please place Figure 2 here)
Nevertheless, a well-known deviation regarding the boron con-
tent of APT, TOF-ERDA and ICP-OES evaluated samples is also visible
within this study – especially, depicted by the green (TOF-ERDA) and
red square (APT), indicating a thin film from the same deposition
run. To obtain comparable values, impurities like H, C, O, N and Ga
have been subtracted from the total concentration of the APT tip, as
all these elements together have been less than 2 at%. Here, we need
to mention that the analyzed volume during APT is way smaller
compared to the other analysis techniques. Nonetheless, the data
sets from all utilized analysis techniques clearly reveal that the de-
tection of an absolute boron concentration is difficult and only
comparable when originating from the same technique/set-up as
highlighted by [30,31]. However, the amount of W and Ta atoms
shows almost perfect coincidence between APT and TOF-ERDA. To
simplify notations, we therefore used TMB
2−z
within the manuscript
(always normalized to the metal sublattice).
3.2. Oxide scale formation
As a constant film thickness is crucial for a clear comparison
during oxidation tests, the deposition times were adapted (from 60
to 85 min with increasing Ta) to grow films between 3.5 and 4.0 µm
for all compositions. Cross sectional micrographs obtained with SEM,
revealed very dense and smooth morphologies for all coatings in the
as deposited state (on Si substrate) – see Fig. 3a-i for WB
2−z
, 3b-i for
W
0.42
Ta
0.58
B
2−z
, and 3c-I for TaB
2−z
, respectively. In contrast to our
previous study, showing a more pronounced columnar growth
morphology, the coatings exhibited a rather fine-grained to amor-
phous, nearly featureless structure. This is mainly attributed to the
decreased deposition temperature used, T
dep
= 500 °C compared to
T
dep
= 700 °C in [12].
In addition to the as deposited state, the corresponding cross-
sectional micrographs after 100 min at 600 °C (3a-ii, 3b-ii, and 3c-ii)
as well as 100 min at 700 °C (3a-iii, 3b-iii, and 3c-iii) oxidation are
also presented. The coatings show stable oxide growth after 100 min
at 600 °C and a decreased scale thickness with increasing Ta content
(see Fig. 3-ii). However, after 100 min at T
ox
= 700 °C all coatings
reveal extensive oxide scales suggesting imminent break-through
oxidation, which is also underlined by the presence of lateral cracks
(see Fig. 3c-iii). Beyond t
ox
= 100 min at T
ox
= 700 °C all coatings start
to spall off, and therefore no 1000-min oxidation tests at that tem-
perature have been conducted. However, all scale thicknesses as well
as scale morphologies were visually and analytically inspected by
cross sectional SEM investigations.
The oxide scale thicknesses in relation to the oxidation times are
summarized exemplarily for all compositions oxidized at T
ox
= 600 °C
see Fig. 4a. In general, the results clearly reveal decreasing oxide
thickness with increasing Ta content – i.e. by comparing WB
2−z
to
TaB
2−z
. The scale thickness significantly decreases from 8.0 µm for
WB
2−z
compared to 1.6 µm for TaB
2−z
after t
ox
= 1000 min. Further-
more, the growth mode changes from a more paralinear behavior for
W-rich coatings to a more linear mode with increasing Ta content.
This is in good agreement with the pure metals, as described by
Kubaschewski and Hopkins [32]. Nevertheless, the more linear
growth mode for the Ta-rich coatings seems to be retarded – which
could to be related to the morphology and density of the formed
oxide – compared to the paralinear behavior for WB
2−z
rich coatings.
Unfortunately, there is no datapoint for the W
0.19
Ta
0.81
B
2−z
coating
(purple diamond) after 1000 min at T
ox
= 600 °C of oxidation because
of spallation (see purple diamonds in Fig. 4a).
3.3. Oxidation kinetics
To gain further information about the oxidation kinetics, the rate
constants
k*
p
(for paralinear scale growth dominated by bulk and
grain boundary diffusion – also called “Regime 2) and k
l
(for linear
rate law dominated by surface reaction of oxygen – so called “Regime
1) were calculated according to the procedure described in [33]. The
calculations yield to an oxidation rate constant of
k*
p
= 6.3 10
−4
µm
2
/s
(oxidation is following a paralinear growth mode) for WB
2−z
(full
squares) which is decreasing to k
l
= 2.6 10
−5
µm/s (linear oxidation
growth mode, but one order of magnitude lower) for TaB
2−z
(full
stars). Coatings with a Ta content of x = 0.34–0.58 show a similar
oxidation behavior, resulting in
k*
p
values of around 1.1 10
−4
µm
2
/s.
Based on this dataset, the oxidation kinetics of W
0.19
Ta
0.81
B
2−z
is
suggested to be similar. The subsections (b–d) of Fig. 4 depict the
oxidation behavior of pure WB
2−z
, W
0.42
Ta
0.58
B
2−z
and pure TaB
2−z
,
respectively, at T
ox
of 500, 600, and 700 °C. The oxide scales increase
drastically for T
ox
= 700 °C already after t
ox
= 100 min, leading to
lateral cracks (see Fig. 3c-iii) and moreover, to non-adherent scales.
The abovementioned gradients in chemical composition due to
Fig. 3. Cross sectional micrographs of (a) WB
2−z
, (b) W
0.42
Ta
0.58
B
2−z
and (c) TaB
2−z
.
Sections (a) – (c) show selected coatings in (i) as deposited state, (ii) after 100 min of
oxidation at 600 °C and (iii) after 100 min of oxidation at 700 °C. The dashed lines are
highlighting the interface between substrate and thin film as well as between thin
film and oxide scale (from bottom to top).
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
4
co-sputtering may lead to this behavior, but also stress relief due to
oxide scale formation in general can be a possible scenario. However,
increased Ta contents clearly retard the oxide scale formation.
3.4. Oxide scale morphology
For a more detailed view on the oxide scale morphology and
structure, TEM investigations of WB
2−z
, W
0.66
Ta
0.34
B
2−z
, and TaB
2−z
after oxidizing at 700 °C for 10 min have been conducted. The results
for α-WB
2−z
, α-W
0.66
Ta
0.34
B
2−z
, as well as multi-phased α-TaB
2−z
+ o-
TaB are presented in Figs. 5–7, respectively. Fig. 5a depicts a cross
sectional bright-field (BF) image of the unaffected coating and scale
on top (from left to right hand side). Here, the highly fine-grained
growth morphology of the unaffected WB
2−z
coating is conceivable
as no features are visible. In contrast, the thermally grown scale
exhibits a granular-like, dense structure that is coarser towards the
outermost regions. The Pt layer on top was deposited during FIB
preparation of the TEM lamella. High resolution images – see Fig. 5b
and c show a fine-grained and dense oxide morphology near the
coating-oxide interface (see Fig. 5b), whereas at the outermost
region relatively large (about ~ 100 nm) crystallites are formed (see
Fig. 5c). This morphological change of the oxide with increasing
thickness is possibly related to the changing growth mode of the
forming WO
3
crystals, whereby the outermost regions undergo the
longest oxidation time and consequently ongoing diffusion (having
also the strongest oxygen supply). However, this area is also slightly
porous at some positions. SAED also emphasizes the formation of a
WO
3
crystals, see indexed pattern in Fig. 5d.
Fig. 6a shows a Scanning TEM – high angle annular dark-field
(STEM-HAADF) image of the oxidized W
0.66
Ta
0.34
B
2−z
thin film. The
formed oxide scale exhibits fine crystallites embedded in an amor-
phous matrix, also showing lateral cracks near the coating-oxide in-
terface (see BF image in Fig. 6b). Here, the crack formation is related to
the growth of WO
3−x
and Ta
2
O
5−x
, which possess different thermal
expansion coefficients (WO
3
~ 12 10
−6
C
−1
, Ta
2
O
5
~ 4 10
−6
C
−1
[34,35])
also with respect to the unaffected W
0.66
Ta
0.34
B
2−z
coating material. The
difference in CTE needs also to be seen with respect to the unintended
variation of the chemical composition originating from the co-sput-
tering process – see APT investigations in Fig. 2 – and hence suggesting
for a partly layered oxide scale and related crack formation. Compared
Fig. 4. Oxide thickness as a function of oxidation time for (a) 600 °C, of all coatings deposited. The oxidation rate constants k
p
(parabolic or paralinear growth mode) and k
l
(linear
growth mode) are depicted in section (a) indicated by individual symbols representing different coating compositions (see legend). Furthermore, the oxidation behavior of (b)
WB
2−z
, (c) W
0.42
Ta
0.58
B
2−z
and (d) TaB
2−z
for 500, 600, and 700 °C oxidation temperature is represented. (For interpretation of the references to colour in this figure legend, the
reader is referred to the web version of this article.)
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
5
to the oxide grown on pure WB
2−z
, the crystallite size of the oxide that
originates from W
0.66
Ta
0.34
B
2−z
stays rather constant over the entire
scale (BF micro-graph, Fig. 6c). The amorphous oxide scale is also
highlighted by the SAED pattern, depicted in Fig. 6d. However, WO
3
based crystallites can be indicated considering the very broad ring-
shaped patterns as depicted in Fig. 6d. (Please place Figure 6 here)
Fig. 7a displays a BF cross section of the oxidized TaB
2−z
coating,
pointing out the areas for recorded SAED (Fig. 7b) and higher
resolution BF image (Fig. 7c) of the formed oxide scale. Here, the very
dense, columnar growth morphology with grain sizes up to 100 nm
is very remarkable. Unfortunately, it was not possible to allocate the
SAED pattern, which can be seen in Fig. 7b, to one distinct crystal
structure. As the oxide layer on top of the TaB
2
film is composed of
various grains revealing different crystal orientations and potentially
different crystal structures, a very diffuse diffraction pattern appears.
Nevertheless, Ta
2
O
5
(in both versions orthorhombic [SG 25, Pmm2]
as well as trigonal [SG 146, R3]) and TaBO
4
(orthorhombic [SG 141,
I4
1
/amd]) crystal structures revealed the best and most reliable
matches in terms of the depicted SAED pattern. (Please place Figure
7 here)
3.5. Oxide scale constitution
To get a deeper insight into the oxide scale constitution, EDS
investigations of oxidized WB
2−z
, W
0.66
Ta
0.34
B
2−z
, and TaB
2−z
coatings
were performed during TEM – again after oxidizing at 700 °C for
10 min. In Fig. 8 the chemical composition is plotted as a function of
the distance, which is highlighted in the STEM cross sections next to
the EDS plots. A clear distinction between the oxide scale and the
unaffected diboride coatings is possible for all three compositions.
The EDS line scan in Fig. 8a reveals a W to O ratio of around 1–3,
which indicates the occurrence of a WO
3
based scale – confirming
the SAED results in Fig. 5d. However, also small amounts of boron
can be detected, but being not dominant for the scale constitution
and related to residuals during a volatile boron-oxide formation.
Increasing the Ta content within the film also lead to a change in the
scale constitution, as suggested by the line scan presented in Fig. 8b.
The chemical composition depicted in Fig. 8b suggests for a highly
oxygen rich scale, but still containing a certain amount of B. This is in
good relation with the structural analysis of the scale presented in
Fig. 6d indicating an amorphous state of the formed oxide scale, as
boron oxide rich scales tend to be amorphous. Nevertheless, the
metal (W,Ta) to O ratio is in the range of 1:4 emphasizing a change in
the scale constitution. Additionally, the chemical composition of the
Fig. 5. Detailed TEM investigations of an oxidized WB
2−z
coating after oxidation at
700 °C for a duration of 10 min. Section (a) shows a bright-field (BF) TEM cross section
of the oxidized coating, pointing out the areas for the HR images displayed in (b) and
(c) and the recorded SAED shown in (d). The indicated ring patterns correspond to a
WO
3
crystal.
Fig. 6. Detailed TEM investigations of an oxidized W
0.66
Ta
0.34
B
2−z
coating after oxi-
dation at 700 °C for a duration of 10 min. Section (a) shows a STEM-HAADF cross
section of the oxidized coating, pointing out the area for the BF image displayed in (b).
Another BF cross section is depicted in (c) indicating the recorded SAED shown in (d).
Fig. 7. TEM micrographs of an oxidized TaB
2−z
coating after oxidation at 700 °C for a
duration of 10 min. Section (a) shows a BF TEM cross section of the oxidized coating,
pointing out the areas for the recorded SAED displayed in (b) and the detailed BF
image of the oxide scale shown in (c). An even more detailed view (high magnified-
TEM) on the oxide is depicted in (d).
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
6
oxidized TaB
2−z
coating (see Fig. 8c) underlines this trend, that the
presence of Ta lead to a decelerated (volatile) boron oxide formation.
The chemical analysis suggests for mixed oxide structures such as
Ta
2
O
5
and TaBO
4
– being in correspondence with the structural
analysis in Fig. 7b.
As it is very difficult to gain reliable values for boron (and partly
oxygen) contents with EDS, the chemical composition of the formed
oxide scales, after 100 h at 600 °C, was also investigated by TOF-
ERDA. Due to resolution limits regarding the distinction of heavy W
and Ta atoms, the elemental amount of tungsten and tantalum is
summarized in a single dataset (blue lines). Fig. 9a reveals a boron
depleted oxide layer (at least ~200 nm from the top of the oxide) on
top of the WB
2
coating, indicating a clear volatile character of boron
oxides during oxidation tests. However, the gradually increasing B
contents for the Ta rich coatings as W
0.19
Ta
0.81
B
2−z
or pure TaB
2−z
suggest a change from the volatile behavior to a more partly-viscous
(glassy) character of the boron oxide through Ta. A similar tendency
was also observed in the EDS line scans conducted during TEM
analysis – see Fig. 8. TOF-ERDA provides information about the
number of recoiled atoms per area, which were hit by impinging
heavy ions, represented by the depth value plotted on the abscissa in
Fig. 9. Therefore, the theoretical layer thickness can be calculated by
means of the density and the molar weight of the containing ele-
ments. For the measured oxides on TaB
2−z
coatings, the interface
between oxide layer and coating was suggested to be located at
around 500 × 10
15
atoms/cm
2
, where no more oxygen could be de-
tected (see Fig. 9c). The atomic density of the oxide (assuming a
mean value of 8.4 g cm
−3
considering orthorhombic and trigonal
Ta
2
O
5
as well as orthorhombic TaBO
4
) leads to a calculated layer
thickness (d
ox_calc
) of ~60 nm. Of course, the calculations do not
consider boron oxide-based structures which are present to small
quantities (as seen in the EDS linescans and ERDA profiles). As SEM
investigations revealed an oxide thickness of ~100 nm after oxidation
at 600 °C and 100 min, a slightly under-dense (~60% density) oxide
layer or a gradually variety/mixed oxide (leading to deviating den-
sities) on TaB
2−z
films is presumed. In Fig. 9b the calculated oxide
thickness for W
0.19
Ta
0.81
B
2−z
was ~220 nm leading to a density of
~40% - relative to a measured oxide thickness of around 600 nm.
Here we applied a weighted atomic density for the oxide, in corre-
spondence to the line scan presented in Fig. 8b, of about 8.2 g cm
−3
.
Unfortunately, the penetration depth of ERDA was too low for de-
tecting the oxide/coating interface of WB
2−z
(see Fig. 9a). Hence, no
thickness and oxide density could be calculated for pure WB
2−z
.
Nevertheless, an even lower oxide density compared to Ta-alloyed
coatings is suggested by analyzing the TEM images represented in
the foregoing Figures. These results lead to the assumption that the
addition of Ta not only decreases the oxide scale growth kinetics but
also increases the scale density.
4. Conclusion
To investigate the oxidation resistance of W
1−x
Ta
x
B
2−z
coatings,
compositions in the full range from WB
2−z
, W
0.85
Ta
0.15
B
2−z
,
W
0.66
Ta
0.34
B
2−z
, W
0.42
Ta
0.58
B
2−z
, W
0.19
Ta
0.81
B
2−z
, to TaB
2−z
were de-
posited by DC magnetron co-sputtering utilizing W
2
B
5
and TaB
2
targets, respectively. ICP-OES as well as TOF-ERDA exhibit similar
tendency of decreasing B contents with increasing amounts of tan-
talum, indicating that the phase formation of α-structured
W
1−x
Ta
x
B
2−z
films at high Ta fractions is limited and directed towards
dual phase structures (α-TaB
2−z
and o-TaB). As highlighted by atom
probe tomography, a layered structure composed of Ta-rich and W-
rich zones is occurring at the nanoscale (in the range of 5–10 nm) –
but it is assumed to be an artifact of co-sputtering.
Fig. 8. EDS line scan of WB
2−z
(a), W
0.66
Ta
0.34
B
2−z
(b) and TaB
2−z
(c) thin films after
oxidation with 700 °C for 10 min. The corresponding STEM micrographs on the right-
hand side, indicate location and distance (white arrow – the tail refers to zero dis-
tance) of the performed EDS measurements. The uncertainty of quantifying light
elements (like B and O) in EDS measurements can be in the range of 5–10%.
Fig. 9. Chemical composition of the oxide layers of WB
2−z
(a), W
0.19
Ta
0.81
B
2−z
(b), and
TaB
2−z
(c) thin films after oxidation at 600 °C for 100 min, investigated by TOF-ERDA
depth profiling. The gray area indicates the calculated oxide thickness by assuming a
WO
3
structured oxide for WB
2
, and mixed oxide phases for W
0.19
Ta
0.81
B
2−z
as well as
TaB
2−z
.
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
7
After oxidation tests in ambient air at 500, 600, 700 °C for 1, 10,
100 and 1000 min, respectively, the formed oxide layers were ana-
lyzed by SEM, TEM, and TOF-ERDA. The investigations showed a
decrease in oxide layer thickness with increasing Ta content for all
temperature settings. At T
ox
= 600 °C and t
ox
= 100 min the oxide
scale thickness decreases steadily from 2.2 µm to 0.1 µm from WB
2−z
to TaB
2−z
. A detailed evaluation of the oxidation kinetics revealed a
paralinear scale growth for W-rich coatings becoming more linear
but retarded with increasing Ta. The corresponding rate constants
decelerate from
k*
p
= 6.3 10
−4
µm
2
/s for.
WB
2−z
to k
l
= 2.6 10
−5
µm/s (linear oxidation growth mode) for
TaB
2−z
. After 100 min at T
ox
= 700 °C all coatings exhibit extensive
scaling suggesting imminent break-through oxidation, also underlined
by the presence of lateral cracks. SAED investigations exhibited a WO
3
based oxide layer for WB
2−z
and mixed oxide scales (e.g. Ta
2
O
5
as well
as TaBO
4
) for TaB
2−z
films, also underlined by TEM-EDS line scans.
Furthermore, TOF-ERDA measurements showed an increased amount
of boron within the oxide scales for coatings with rising Ta content,
suggesting for a decelerated formation of a volatile boron oxides.
In summary, the addition of Ta to WB
2−z
based coatings retards
the oxide scale kinetics through the formation of denser, less volatile,
and adherent scales, thus being a key factor for the enhanced oxi-
dation resistance. An optimum composition of W
1−x
Ta
x
B
2−z
based
coatings would be in the range of x = 0.2–0.3, combining enhanced
fracture toughness (≥3.0 MPa√m) as well as super hardness (≥40 GPa)
next to a decent oxidation resistance.
CRediT authorship contribution statement
C. Fuger: Conceptualization, Software, Investigation, Writing - ori-
ginal draft. B. Schwartz: Investigation. T. Wojcik: Investigation. V.
Moraes: Investigation. M. Weiss: Investigation, Writing - review &
editing. A. Limbeck: Investigation. C.A. Macauley: Investigation,
Writing - review & editing. O. Hunold: Writing - review & editing.
P. Polcik: Writing - review & editing. D. Primetzhofer: Investigation,
Writing - review & editing. P. Felfer: Investigation, Writing - review &
editing. P.H. Mayrhofer: Writing - review & editing. 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 in-
fluence the work reported in this paper.
Acknowledgements
The authors greatly acknowledge the financial support of Plansee
Composite Materials GmbH, Oerlikon Balzers Surface Solutions AG and
the Christian Doppler Gesellschaft within the framework of the Christian
Doppler Laboratory for Surface Engineering of high-performance
Components. SEM and TEM investigations were carried out using facil-
ities of the USTEM center of TU Wien, Austria. CM and PF acknowledge
financial support by the Bavarian Ministry of Economic Affairs and Media,
Energy and Technology for the joint projects in the framework of the
Helmholtz Institute Erlangen-Nürnberg for Renewable Energy (IEK-11) of
Forschungszentrum Jülich. CAM and PF would also like to acknowledge
funding by the Deutsche Forschungsgemeinschaft (DFG) via the Cluster of
Excellence ‘Engineering of Advanced Materials’ (project EXC 315). Support
by VR-RFI (contracts #821-012-5144 and #2017-00646_9) and the
Swedish Foundation for Strategic Research (SSF, contract RIF14-0053)
supporting accelerator operation is gratefully acknowledged.
Appendix
See Appendix Fig. A1.
Fig. A 1. Structural evolution of W
1−x
Ta
x
B
2−z
thin films obtained by XRD analysis, realized by a Philips XPERT diffractometer in Bragg-Brentano configuration equipped with a Cu-
Kα (λ = 1.54 Å) radiation source. The green, red and yellow dashed vertical lines are indicating the α-WB
2
, α-TaB
2
and o-TaB phase, respectively [12].
C. Fuger, B. Schwartz, T. Wojcik et al. Journal of Alloys and Compounds xxx (xxxx) xxx
8
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... Therefore, the introduction of tantalum or zirconium into the WB2 crystal lattice can give the opportunity to create a new group of hard and refractory materials. Hitherto, the W Ta B2 system has possessed greater interest of researchers [3,11,12]. The theoretical and experimental studies have been mainly concerned on the obtaining of this material in the form of coatings. ...
... Experimental studies have shown that the obtained layers are characterized by great hardness of around 45 GPa and fracture toughness KIC values 3.0 MPa √ m what make this compound much better than common TiN , Ti Si N and (T i, Al)N [11]. Also, the good thermal stability and high oxidation resistance at temperatures up to 700 • C [12] makes them a promising candidate for industrial applications. Additional function of theoretical calculations is possibility of explanation of material properties origin. ...
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Potentially superhard $W_{1-x}Zr_xB_2$ polymorph hP6-P6$_3$/mmc-$WB_2$ with zirconium doping in the range of x=0.0-0.25 was thoroughly analyzed within the framework of first-principles density functional theory from the structural and mechanical point of view. The obtained results were subsequently compared with properties of material deposited by magnetron sputtering method. All predicted structures are mechanically and thermodynamically stable. Due to theoretical calculations zirconium doping reduces hardness and fracture toughness $K_{IC}$ of $WB_2$. Deposited films are characterized by greater hardness $H_v$ but lower fracture toughness $K_{IC}$. The results of experiments show that not only solid solution hardening is responsible for strengthening of predicted new material but also change of microstructure, Hall-Petch effect and boron vacancies.
... Such environments are relevant for cutting tools or the aerospace industry [6]. However, in practice, their application is limited through their poor oxidation resistance [7,8]. Pure borides tend to form a B 2 O 3 layer upon oxidation, which evaporates at elevated temperatures, a process further enhanced by the presence of water vapor by forming the more volatile boric acid, thus limiting the oxidation resistance [9][10][11]. ...
Article
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The increased demand for sustainability requires, among others, the development of new materials with enhanced corrosion resistance. Transition metal diborides are exceptional candidates, as they exhibit fascinating mechanical and thermal properties. However, at elevated temperatures and oxidizing atmospheres, their use is limited due to the fact of their inadequate oxidation resistance. Recently, it was found that chromium diboride doped with silicon can overcome this limitation. Further improvement of this protective coating requires detailed knowledge regarding the composition of the forming oxide layer and the change in the composition of the remaining thin film. In this work, an analytical method for the quantitative measurement of depth profiles without using matrix-matched reference materials was developed. Using this approach, based on the recently introduced online-LASIL technique, it was possible to achieve a depth resolution of 240 nm. A further decrease in the ablation rate is possible but demands a more sensitive detection of silicon. Two chromium diboride samples with different Si contents suffering an oxidation treatment were used to demonstrate the capabilities of this technique. The concentration profiles resembled the pathway of the formed oxidation layers as monitored with transmission electron microscopy. The stoichiometry of the oxidation layers differed strongly between the samples, suggesting different processes were taking place. The validity of the LASIL results was cross-checked with several other analytical techniques.
... Contrary to bulk diborides, the B 2 O 3 phase evaporates rapidly at temperatures >~400 • C from overstoichiometric TMB 2 thin films. This leads to the formation of a highly porous oxide layer with no oxidation protection [20,32,33]. As also proved for TiN-based coatings, alloying diboride thin films with Al significantly improves their oxidation resistance properties [20]. ...
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We study microstructure, mechanical, and corrosion properties of Zr1-xCrxBy coatings deposited by hybrid high-power impulse/DC magnetron co-sputtering (CrB2-HiPIMS/ZrB2-DCMS). Cr/(Zr+Cr) ratio, x, increases from 0.13 to 0.9, while B/(Zr+Cr) ratio, y, decreases from 2.92 to 1.81. As reference, ZrB2.18 and CrB1.81 layers are grown at 4000 W DCMS. ZrB2.18 and CrB1.81 columns are continual from near substrate toward the surface with open column boundaries. We find that the critical growth parameter to achieve dense films is the ratio of Cr+-dominated ion flux and the Zr+B neutral flux from the ZrB2 target. Thus, the alloys are categorized in two groups: films with x < 0.32 (low Cr+/(Zr+B) ratios) that have continuous columnar growth, rough surfaces, and open column boundaries, and films with x ≥ 0.32 (high Cr+/(Zr+B) ratios) that Cr+-dominated ion fluxes are sufficient to interrupt continuous columns, resulting in smooth surface and dense fine-grain microstructure. The pulsed metal-ion irradiation is more effective in film densification than continuous Ar+ bombardment. Dense Zr0.46Cr0.54B2.40 and Zr0.10Cr0.90B1.81 alloys are hard (> 30 GPa) and almost stress-free with relative nanoindentation toughness of 1.3 MPa√m and 2.3 MPa√m, respectively, and remarkedly low corrosion rates (∼1.0 × 10-6 mA/cm2 for Zr0.46Cr0.54B2.40 and ∼2.1 × 10-6 mA/cm2 for Zr0.10Cr0.90B1.81).
... Moreover, WB 2±z is known for its enhanced material properties like hardness, Young's modulus as well as fracture toughness, and therefore often consulted as a base system forming ternary TM I TM II B 2±z [10,11,20,21]. In previous studies, we successfully showed the positive effect of Ta alloying on the mechanical properties (H, K IC ), thermal stability, and oxidation resistance of WB 2±z thin films [10,21,22]. ...
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Transition metal diboride-based thin films are promising candidates to replace state-of-the-art protective and functional coating materials due to their unique properties. Here, we focus on hexagonal WB2−z , showing that the AlB2 structure is stabilized by B vacancies exhibiting its energetic minima at sub-stoichiometric WB1.5. Nanoindentation reveals super-hardness of 0001 oriented α-WB2−z coatings , linearly decreasing by more than 15 GPa with predominant 1011 orientation. This anisotropy is attributed to differences in the generalized stacking fault energy of basal and pyramidal slip systems, highlighting the feasibility of tuning mechanical properties by crystallographic orientation relations.
... The layered oxide structure with B 2 O 3 is well reported in various publications, such as for the oxidation of bulk ZrB 2 or HfB 2 [8,39]. On the other hand studies on the oxidation behavior of TM-boride coatings [18,40], done in a conventional furnace in ambient air, do not report any B 2 O 3 at the surface. Therefore, we performed comparative oxidation treatments also in a conventional furnace in ambient air confirming this observation. ...
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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.
... The inclusion of Al was accompanied by a slight reduction in hardness from 42 GPa in the TiB 2.4 thin film, to 39 GPa for the Ti 0.68 Al 0.32 B 1.35 thin film. Reduction in oxidation rate has also been observed in WB x through alloying with Ta [21], with many recent studies expanding the understanding of the oxidation kinetics in hard coatings [22][23][24]. ...
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Transition-metal diboride thin films, which have high melting points, excellent hardness, and good chemical and thermal conductivity, severely suffer from rapid oxidation in air. Here, we explore the influence of varying B content and resulting nanostructure change on the oxidation properties of TiBx thin films, with x = 1.43, 2.20, and 2.70. Results show that all as-deposited layers have columnar structure. The column boundaries of as-deposited TiB2.20 and TiB2.70 films grown by direct current magnetron sputtering (DCMS) are B-rich, while the as-deposited TiB1.43 films grown by high-power impulse magnetron sputtering (HiPIMS) show no apparent grain boundary phases and contain Ti-rich planar defects. The oxidation rate of TiB1.43 air-annealed at 400 oC up to 48 h is significantly lower than that of TiB2.20 and TiB2.70 films. The oxidation rate of TiB1.43, TiB2.20, and TiB2.70 films was measured at 2.9±1.5, 7.1±1.0, and 20.0±5.0 nm/h, respectively, with no spallation of even as thick oxide scales as 0.5 μm in any of the films. The improved oxidation resistance can be explained by the absence of B-rich tissue phase at the column boundaries of understoichiometric TiBx films, a phase that interlaces the nanocolumnar TiB2 structures in the corresponding overstoichiometric films. An easy oxidation pathway is thus eliminated.
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Within physical vapor deposited Hf-Si-B2±z thin films, selective diffusion-driven oxidation of Si is identified to cause outstanding oxidation resistance at temperatures up to 1500 °C. After 60 h at 1200 °C, the initially 2.47 µm thin Hf0.20Si0.23B0.57 thin film exhibits a dense oxide scale of only 1.56 µm. The thermally induced decomposition of metastable Hf-Si-B2±z leads not only to the formation of Si precipitates within the remaining thin film (related to a non-homogenous Si distribution after the deposition) but also to pure Si layers on top and bottom of the Hf-Si-B2±z coatings next to the excellent adherend SiO2 based scales.
Article
Magnetron sputtered WB2 coatings doped with 8, 11 and 16 at.% zirconium were analysed using energy dispersive spectroscopy, X-ray diffraction and nanoindentation under the load of 4, 7 and 10 mN. It has been observed that these coatings crystallize in the α-AlB2 and ω-W2B5 prototype structure. Phenomenon responsible for this is an increase of the zirconium content which causes an increase in the ω-W2B5 phase. All the deposited coatings have a hardness of about 45 GPa while Young's modulus drops down from 497 to 480 GPa with increasing zirconium content. Coatings without doping and doped with 16 at.% zirconium were annealed at 650 °C and subjected to cyclic thermal loads using a maximum temperature 600 °C and cooling with a compressed air. It has been observed that addition of zirconium improved coatings phase stability.
Article
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The concept of Si alloyed transition metal (TM) diborides – well explored for bulk ceramics – is studied for five different physical vapor deposited TM-Si-B2±z (TM = Ti, Cr, Hf, Ta, W) coatings, focusing on the oxidation behavior up to 1200 °C. In their as deposited state, all coatings exhibit single phased AlB2 prototype structures, whereby the addition of Si results in dense, refined morphologies with no additional phases visible in the X-ray diffractograms. With already low amounts of Si, the slope of the mass increase during dynamic oxidation flattens, especially for Ti-Si-B2±z, Cr-Si-B2±z, and Hf-Si-B2±z. Above distinct Si contents, the formation of a steady state region exhibiting no further mass increase is promoted (starting at around 1000 to 1100 °C). Best results are obtained for Hf0.21Si0.18B0.61 and Cr0.26Si0.16B0.58 (both around 2.4 μm thick in the as deposited sate), revealing drastically retarded oxidation kinetics forming 400 nm thin oxide scales after 3 h at 1200 °C in ambient air (significantly lower compared to bulk ceramics). This highly protective oxidation mechanism is attributed to the formation of an amorphous Si rich oxide scale. The Si content needed to form these oxide scales largely differs between the TM-Si-B2±z coatings investigated, also diversifying the prevalent oxidation mechanism, especially for Cr-Si-B2±z.
Article
As an efficient tool for theoretical investigation, the first-principles calculation based on Density Functional Theory (DFT) has been applied to various fields of material development. In the present work, the first-principles calculation has been employed to predict the mechanic characteristics and thermal performances of three novel WB4 tetraborides. The single crystal constant Cij calculated by the stress-strain method shows that hP10-WB4 and hR15-WB4 tetraborides have stable crystal structures and elastic anisotropies. However, the calculated phonon dispersions reveal that hP20-WB4 is dynamically unstable. According to the calculated elastic modulus and anisotropy indexes A, the sequence of elastic anisotropy is hP10-WB4 < hR15-WB4. Besides, the Debye temperature of WB4 has such a sequence: hP10(737.042K) > hR15(640.517K), and the sound velocities have great anisotropy. Moreover, the calculated kmin shows the strong anisotropy and their values show the order: hP10 > hR15.
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We compare the chemical composition of TiAlN thin films determined by ion beam analysis and laser-assisted atom probe tomography (APT). The laser pulse energy during APT was increased subsequently from 10 to 20, 30, 40, 50, 100 and 200 pJ within a single measurement, covering the range that is typically employed for the analysis of transition metal nitrides. The laser pulse energy-dependent Ti, Al and N concentrations were compared to ion beam analysis data, combining Rutherford backscattering spectrometry (RBS) and elastic recoil detection analysis (ERDA) with the total measurement uncertainty of 2.5% relative deviation. It can be learned that the absolute N concentration from APT is underestimated by at least 5.5 at.% (up to 8.2 at.%) and the absolute Al concentration from APT is overestimated by at least 4.5 at.% (up to 6.2 at.%), while absolute Ti concentration values are for both techniques in good agreement with maximum deviations < 2 at.%. Hence, the here presented comparative analysis clearly shows that absolute Al and N concentration values obtained by ion beam analysis deviate significantly to the APT data for the laser pulse energy range from 10 to 200 pJ. Possible causes for the compositional discrepancy between Rutherford backscattering spectrometry/elastic recoil detection analysis and APT, such as molecular ions, multiple detection events and preferential evaporation/retention of species with different evaporation fields, are discussed. The presented data emphasize that laser-assisted APT is a precise tool to quantify the chemical composition of TiAlN thin films, that lacks accuracy.
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Based on density functional theory, we recently suggested that metastable α-WB2 is a promising candidate combining very high hardness with high toughness. These calculations further suggested that the addition of Tantalum supports the crystallization of α-structured W1-xTaxB2-z, with only minor reduction in toughness. Thus, various Ta containing WB2-based coatings have been synthesized using physical vapor deposition. With increasing Ta content, the hardness increases from ∼41 GPa (WB2) to ∼45 GPa (W0.74Ta0.26B2). In situ micromechanical cantilever bending tests exhibit fracture toughness KIC values of 3.7 to 3.0 MPa√m for increasing Ta content (single-phased up to 26 at.% Ta).
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The demand to discover new materials is scientifically as well as industrially a continuously present topic, covering all different fields of application. The recent scientific work on thin film materials has shown, that especially for nitride-based protective coatings, computationally-driven understanding and modelling serves as a reliable trend-giver and can be used for target-oriented experiments. In this study, semi-automated density functional theory (DFT) calculations were used, to sweep across transition metal diborides in order to characterize their structure, phase stability and mechanical properties. We show that early transition metal diborides (TiB2, VB2, etc.) tend to be chemically more stable in the AlB2 structure type, whereas late transition metal diborides (WB2, ReB2, etc.) are preferably stabilized in the W2B5−x structure type. Closely related, we could prove that point defects such as vacancies significantly influence the phase stability and even can reverse the preference for the AlB2 or W2B5−x structure. Furthermore, investigations on the brittle-ductile behavior of the various diborides reveal, that the metastable structures are more ductile than their stable counterparts (WB2, TcB2, etc.). To design thin film materials, e.g. ternary or layered systems, this study is important for application oriented coating development to focus experimental studies on the most perspective systems.
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Reactive high-power impulse magnetron sputtering (R-HiPIMS) is seen as a key-technology for the deposition of future hard and multifunctional coatings. Increased ionisation rates allow for additional possibilities in tuning specific coating characteristics based on growth mechanisms varied by surface-diffusion. Especially within the well-established Ti-Al-N system, the Al solubility limit (xmax) of metastable face-centred-cubic (fcc) Ti1−xAlxN is an everlasting scientific topic. Here, we investigate in detail the dependence of xmax on various deposition parameters (i.e. pulse frequency and duration, N2-to-Ar flow rate ratio, and substrate bias potential) during R-HiPIMS of Ti-Al-N coatings using Ti0.6Al0.4, Ti0.5Al0.5 and Ti0.4Al0.6 composite targets. The systematic studies showed that the highest solubility limit of xmax ∼0.55 could be obtained for duty cycles around 3.75% (peak power densities of ∼1.0 kW/cm2) and a N2-to-Ar flow rate ratio of 0.3. A further decisive fact for the deposition of high Al containing fcc-structured Ti1−xAlxN coatings is surface diffusion controlled by bias potentials (DC as well as modulated pulses) ensuring sufficient intermixing of the arriving film species. Despite the presence of very small amounts of wurtzite-typed phases, excellent hardness values of ∼36 GPa for Ti0.40Al0.60N – which further increased to ∼ 40 GPa upon annealing for 1 h at 700 °C – could be achieved for a DC bias potential of -50 V, irrespective of all variations conducted. Based on our results we can further conclude, that the ratio and energy of Tin+- and Aln+-ions, simultaneously arriving at the substrate surface, are decisive for stabilising the highly preferred cubic modification. A distinct promotion of specific discharge regimes – selected by synchronised bias pulses – can thus positively influence the cubic phase formation through altered gas-to-metal ion ratios arriving at the film surface.
Article
Using a combination of density functional theory calculations and nanomechanical testing of sputter-deposited, 110-oriented Ta0.47C0.34N0.19 thin films, we show that non-metal alloying – substituting C with N atoms – in TaC results in a super-hard material with enhanced ductility. Based on the calculated elastic constants, with Pugh and Pettifor criteria for ductile character, we predict that stoichiometric and sub-stoichiometric Ta-C-N alloys are more ductile than Ta-C compounds. From nanoindentation of the as-deposited coating, we measure hardness of 43 ± 1.4 GPa. In situ scanning electron microscopy (SEM) based micro-compression of cylindrical pillars, prepared via focused ion beam milling of the coating, revealed that Ta-C-N alloys are ductile and undergo plastic deformation with a yield strength of 17 ± 1.4 GPa. The post-compression SEM images of the pillars show {111} as the active slip system operating during compression. Additional in situ SEM based cantilever tests suggest that the Ta-C-N films exhibit superior fracture toughness compared to Ta-C coatings. Our results provide a new perspective on the role of alloying on the mechanical behavior of ultra-high temperature compounds such as transition-metal carbides.
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The present study investigates the performance of four selected PVD coating materials in application near hot-forming procedure of Al-Si coated 22MnB5 steel sheets. These are: a Ti0.57Al0.43N/Mo-Si-B multilayer, a TiN base layer having a Mo-Si-B top layer, a DLC, and an Al-Cr-O coating. We determined the amount of adherent material (after the hot-forming procedure) by 3D topographical analysis and characterised the build-up volume as well as the coating/build-up interface constitutions with respect to the diffusivity of the participating elements. The smallest build-up volume could be identified for the Ti0.57Al0.43N/Mo-Si-B multilayer as well as the DLC coating, whereas the highest adhesive wear peak-value during the whole testing sequence was detected for the Al-Cr-O coating. Based on detailed TEM analysis of the diffusion zones (between build-ups and coating materials), we can conclude that the smallest adhesive wear is obtained, if the interface formation (between build-up and coatings) is narrow and Fe-free or at least Fe-lean.
Article
Refractory transition-metal diborides exhibit inherent hardness. However, this is not always sufficient to prevent failure in applications involving high mechanical and thermal stress, since hardness is typically accompanied by brittleness leading to crack formation and propagation. Toughness, the combination of hardness and ductility, is required to avoid brittle fracture. Here, the authors demonstrate a strategy for simultaneously enhancing both hardness and ductility of ZrB2-rich thin films grown in pure Ar on Al2O3(0001) and Si(001) substrates at 475 °C. ZrB2.4 layers are deposited by dc magnetron sputtering (DCMS) from a ZrB2 target, while Zr1−xTaxBy alloy films are grown, thus varying the B/metal ratio as a function of x, by adding pulsed high-power impulse magnetron sputtering (HiPIMS) from a Ta target to deposit Zr1−xTaxBy alloy films using hybrid Ta-HiPIMS/ZrB2-DCMS sputtering with a substrate bias synchronized to the metal-rich portion of each HiPIMS pulse. The average power 𝒫Ta (and pulse frequency) applied to the HiPIMS Ta target is varied from 0 to 1800 W (0 to 300 Hz) in increments of 600 W (100 Hz). The resulting boron-to-metal ratio, y = B/(Zr+Ta), in as-deposited Zr1−xTaxBy films decreases from 2.4 to 1.5 as 𝒫Ta is increased from 0 to 1800 W, while x increases from 0 to 0.3. A combination of x-ray diffraction (XRD), glancing-angle XRD, transmission electron microscopy (TEM), analytical Z-contrast scanning TEM, electron energy-loss spectroscopy, energy-dispersive x-ray spectroscopy, x-ray photoelectron spectroscopy, and atom-probe tomography reveals that all films have the hexagonal AlB2 crystal structure with a columnar nanostructure, in which the column boundaries of layers with 0 ≤ x < 0.2 are B-rich, whereas those with x ≥ 0.2 are Ta-rich. The nanostructural transition, combined with changes in average column widths, results in an ∼20% increase in hardness, from 35 to 42 GPa, with a simultaneous increase of ∼30% in nanoindentation toughness, from 4.0 to 5.2 MPa√m.
Article
WB 2 /CrN multilayer films with thick modulation periods over 50 nm (Λ = 1400, 315, 235, 150, 55 nm) were synthesized by direct-current magnetron sputtering, and the influence of modulation period on microstructure and mechanical properties for the multilayer films was systematically studied. In WB 2 /CrN multilayer films, CrN sublayers present the columnar microstructure. As Λ decreases, the structure of WB 2 sublayers evolves from (110) orientation to (001) orientation to amorphous structure, and critical crystalline thickness for WB 2 sublayers is over 150 nm here. A transition layer, which shows the columnar crystal with size of 10–11 nm high and 2.5–3.5 nm wide caused by the effect of the crystalline interface of the CrN sublayers, is detected in WB 2 sublayers. Additionally, a-BN, WB 2 , WB 2 (N), CrN, Cr 2 N and Cr 2 O 3 phase are formed in the multilayer films. Moreover, film hardness mainly obeys the rule of mixture. The maximum hardness of 31.2 GPa is obtained at Λ = 315 nm due to crystalline WB 2 sublayers with (001) preferred orientation, and amorphous WB 2 sublayers greatly reduce the film hardness to only 22.3–24.3 GPa at Λ ≤ 235 nm. Consequently, the poor hardness leads to the higher wear rates (5.7–7.8 × 10 ⁻⁷ mm ³ /mN) of multilayer films with Λ ≤ 235 nm compared with those (2.9–3.3 × 10 ⁻⁷ mm ³ /mN) of other films. However, both the fracture toughness and adhesive strength of the films present an increasing trend with decreasing Λ resulting from the soft CrN and BN phases and a certain amount of interface. In conclusion, decreasing the critical crystal-thickness of the WB 2 sublayers, controlling the N content in WB 2 sublayers and getting sharp interfaces will play important roles in developing the higher-performance WB 2 /CrN multilayer films.
Article
Solid solutions of tungsten diboride (WB2) with increasing substitution of tungsten (W) by tantalum (Ta) and niobium (Nb) - ranging from 0 to 50 at. % on a metals basis - were synthesized through resistive arc melting. Samples were characterized using a combination of powder X-ray diffraction (PXRD) for phase identification, energy-dispersive X-ray spectroscopy and X-ray photoelectron spectroscopy for elemental composition, Vickers microindentation for hardness measurements, and thermogravimetric analysis for thermal stability. The solubility limit was found to be less than 8 at. % for Nb and less than 10 at. % for Ta, as determined by PXRD. Vickers hardness (Hv) values were measured to be 40.3 ± 1.6 and 41.0 ± 1.2 GPa at 0.49 N for 6 at. % Nb and for 8 at. % Ta substitution, respectively. In addition, the hardest solid solution (W0.92Ta0.08B2) showed oxidation resistance up to ∼570 °C, approximately 70 °C higher than that of tungsten carbide (WC). Although pure WB2 is known not to be superhard, these results demonstrate the formation of superhard solid solutions through the substitution of tungsten by small amounts of transition metals. This increase in hardness can be attributed to solid solution hardening.
Article
Nano-composites often represent a challenge for spatially resolved analysing methods. The overlap of specific signals, for example in the mass-to-charge spectrum during atom probe tomography (APT) analysis or the intersection of electron energy loss edges in electron energy loss spectroscopy (EELS), leads to highly challenging results. Therefore, a complementary use of different techniques is essential to fully characterise a given system. Within this work, we studied our Ti-Al-N/Mo-Si-B multilayer (comprising alternating 31 nm thin arc evaporated fcc-Ti0.57Al0.43N layers and amorphous 6 nm thin Mo0.58Si0.28B0.14 layers), by high-resolution transmission electron microscopy, EELS as well as APT and completed the results by comprehensive X-ray diffraction measurements. We focused on the microstructure and crystallographic evolution during thermal loading - up to 1400 °C - and hence a detailed analytical description applying a wide set of high resolution techniques. When exposed to high temperatures, the as-deposited amorphous MoSiB layers form the oxidation resistant, intermetallic phases T1-Mo5Si3 and T2-Mo5SiB2. The layered arrangement between Ti-Al-N and Mo-Si-B allows to postpone the formation of the comparative soft, hexagonal wurtzite type AlN (due to the decomposition of fcc-Ti1-xAlxN layers) up to 1200 °C, which is by ∼200–300 °C above the typical formation temperature of w-AlN in homogeneously grown single phase fcc-Ti1-xAlxN thin films.