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Crystallization characteristics and chemical bonding properties of nickel carbide thin film nanocomposites

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The crystal structure and chemical bonding of magnetron-sputtering deposited nickel carbide Ni 1−x C x (0.05 x 0.62) thin films have been investigated by high-resolution x-ray diffraction, transmission electron microscopy, x-ray photoelectron spectroscopy, Raman spectroscopy, and soft x-ray absorption spectroscopy. By using x-ray as well as electron diffraction, we found carbon-containing hcp-Ni (hcp-NiC y phase), instead of the expected rhombohedral-Ni 3 C. At low carbon content (4.9 at%), the thin film consists of hcp-NiC y nanocrystallites mixed with a smaller amount of fcc-NiC x . The average grain size is about 10–20 nm. With the increase of carbon content to 16.3 at%, the film contains single-phase hcp-NiC y nanocrystallites with expanded lattice parameters. With a further increase of carbon content to 38 at%, and 62 at%, the films transform to x-ray amorphous materials with hcp-NiC y and fcc-NiC x nanodomain structures in an amorphous carbon-rich matrix. Raman spectra of carbon indicate dominant sp 2 hybridization, consistent with photoelectron spectra that show a decreasing amount of C–Ni phase with increasing carbon content. The Ni 3d–C 2p hybridization in the hexagonal structure gives rise to the salient double-peak structure in Ni 2p soft x-ray absorption spectra at 16.3 at% that changes with carbon content. We also show that the resistivity is not only governed by the amount of carbon, but increases by more than a factor of two when the samples transform from crystalline to amorphous.
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Crystallization characteristics and chemical bonding properties of nickel carbide thin film
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2014 J. Phys.: Condens. Matter 26 415501
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Journal of Physics: Condensed Matter
J. Phys.: Condens. Matter 26 (2014) 415501 (11pp) doi:10.1088/0953-8984/26/41/415501
Crystallization characteristics and chemical
bonding properties of nickel carbide thin
film nanocomposites
Andrej Furlan1,3, Jun Lu1, Lars Hultman1, Ulf Jansson2and
Martin Magnuson1
1Thin Film Physics Division, Department of Physics, IFM, Link¨
oping University, SE-58183 Link¨
2Department of Chemistry, Ångstr¨
om Laboratory, Uppsala University, Box 538, SE-751 21 Uppsala,
Received 22 June 2014, revised 29 July 2014
Accepted for publication 9 August 2014
Published 19 September 2014
The crystal structure and chemical bonding of magnetron-sputtering deposited nickel carbide
Ni1xCx(0.05 x0.62) thin films have been investigated by high-resolution x-ray
diffraction, transmission electron microscopy, x-ray photoelectron spectroscopy, Raman
spectroscopy, and soft x-ray absorption spectroscopy. By using x-ray as well as electron
diffraction, we found carbon-containing hcp-Ni (hcp-NiCyphase), instead of the expected
rhombohedral-Ni3C. At low carbon content (4.9 at%), the thin film consists of hcp-NiCy
nanocrystallites mixed with a smaller amount of fcc-NiCx. The average grain size is about
10–20 nm. With the increase of carbon content to 16.3 at%, the film contains single-phase
hcp-NiCynanocrystallites with expanded lattice parameters. With a further increase of carbon
content to 38 at%, and 62 at%, the films transform to x-ray amorphous materials with
hcp-NiCyand fcc-NiCxnanodomain structures in an amorphous carbon-rich matrix. Raman
spectra of carbon indicate dominant sp2hybridization, consistent with photoelectron spectra
that show a decreasing amount of C–Ni phase with increasing carbon content. The Ni 3d–C 2p
hybridization in the hexagonal structure gives rise to the salient double-peak structure in Ni 2p
soft x-ray absorption spectra at 16.3 at% that changes with carbon content. We also show that
the resistivity is not only governed by the amount of carbon, but increases by more than a
factor of two when the samples transform from crystalline to amorphous.
Keywords: amorphous nanocomposites, thin film coatings, transition metal carbides,
magnetron sputtering
(Some figures may appear in colour only in the online journal)
1. Introduction
Transition metal carbides are useful in various applications
ranging from wear and oxidation resistant protective coatings
to low friction solid lubricants [1,2]. This flexibility is due
to the nanocomposite nanocrystalline/amorphous-C structure
that governs the coating’s properties depending on the amount
of amorphous matrix, and crystallite size of the carbide [3,4].
3Present address: AMB Industry, Kvarnv¨
agen 26, 361 93 Broakulla, Sweden.
Early transition metals such as Ti, Zr, and V form strong
covalent metal carbon bonds often in cubic crystals in contrast
to late transition metals such as Fe and Ni that form less
strong Me–C bonds [5]. The late transition metals usually
form completely amorphous or mainly amorphous materials
with complex nanocrystallites above a threshold value around
20 at%. An important exception is the Ni–C system, where
metastable rhombohedral-Ni3C nanocrystallites are easily
formed in a large composition range [6,7], and it is more
difficult to form completely amorphous films [8]. Using RF
0953-8984/14/415501+11$33.00 1 © 2014 IOP Publishing Ltd Printed in the UK
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
sputtering [6], partly amorphous Ni1xCxfilms with Ni3C
crystallites embedded in an amorphous Ni1xCxphase have
previously been obtained for x=0.35 [6]. Amorphous Ni–
C films have also been obtained for x>0.5 using reactive
co-sputtering with CH4as a carbon source [7].
Metallic Ni usually has a cubic fcc structure with space
group Fm–3m. In addition, the hcp-Ni phase with space group
P63/mmc has been reported [9,10]. Pure hcp-Ni metal is very
unstable, and most previous investigations lack a material com-
position analysis [6,7]. One study showed that nano crystallites
of hcp-Ni metal could be synthesized, and likely stabilized by
carbon, but easily transformed into fcc-Ni metal when the size
extends more than 5 nm [11]. Because both hcp-Ni with space
group P63/mmc and Ni3C with space group R-3cH contain
carbon and have very similar crystal structures, most previ-
ous work did not experimentally show how to discriminate
hcp-Ni from Ni3C[12,13]. Only recently, Schaefer et al [14]
were able to distinguish between hcp-Ni and Ni3C structures
by means of low-angle x-ray diffraction. Schaefer pointed out
that all previously reported hcp-Ni contains carbon, and should
be described as rhombohedral Ni3C[13,14] in a superstruc-
ture with interstitially ordered carbon. The superstructure can
be approximated by a hexagonal subcell that is nearly identical
in size to that of hcp-Ni with lattice constants a=2.682 Å,
and c=4.306 Å [15].
In this work, we investigate the nanocomposite-to-
amorphous structure, and the nature of chemical bonding be-
tween Ni and C for a range of C concentrations (0.05 x
0.62) in magnetron sputtered Ni1xCxfilms. As a non-
equilibrium process, magnetron sputtering may increase the
solubility of C into Ni, and the carbide phase of the film
structure may be influenced by the total C content. By
employing a combination of x-ray diffraction, high-resolution
transmission electron microscopy (HR-TEM), x-ray photo-
electron spectroscopy (XPS), Raman, and soft x-ray absorp-
tion spectroscopy (XAS), we analyze the Ni carbide, and Ni
metal-like nanocrystalline to amorphous contributions to the
structure, and the dependence on the carbon content. In par-
ticular, we identify the crystallization of the Ni–C system into
hcp-Ni and fcc-Ni by combining x-ray diffraction (XRD) with
HR-TEM for low carbon contents. We show that the samples
do not form a superstructure of Ni3C with ordered carbon as
previously thought. XPS characterization gives a quantitative
analysis of the compositions of different Ni–C phases with par-
ticular emphasis on the variation of the carbon content in the
carbide phase. The electrical resistivity of the Ni1xCxfilms
is correlated to the amount of C–Ni and C–C bonds, the degree
of crystallization, and the total carbon content.
2. Experimental details
2.1. Synthesis and deposition
All of the investigated films were deposited by dual dc mag-
netron sputtering in ultra high vacuum (UHV) on single-crystal
Si(0 0 1) (10 ×10 mm) substrates. Prior to deposition, the sub-
strates were cleaned in ultrasound baths of acetone and iso-
propyl alcohol. During deposition, the substrates biased to
50 V, and preheated to 250 C from the back side by a resis-
tive heater built-into the substrate holder. This made it possible
to synthesize the films with a high degree of purity, and with
precisely tuned composition. The Ni1xCxthin films were
deposited in a UHV chamber with a base pressure of 109Pa
from a double current regulated 2-inch magnetron setting in an
Ar discharge generated at 3.0 mTorr, and with a gas flow rate
of 30 sccm. The magnetrons were directed towards a rotating
substrate holder at a distance of 15 cm. As separates puttering
sources, graphite, and a non-elemental Ni+C target were used
(99.999% pure C, and 99.95% pure Ni). To enable the mag-
netic field from the magnetron to reach the plasma through the
ferromagnetic Ni target, a segmented design was used in which
a circular center part of the target was removed, and placed on
a graphite plate. In this way, simultaneous sputtering of Ni
and C from the same target was accomplished [4]. The tuning
of the film composition was achieved by keeping the graphite
target at a constant current of 300 mA, and tuning the current
on the Ni target. The resulting thicknesses of the as-deposited
coatings were 740 nm (x=0.05), 635 nm (x=0.16), 309 nm
(x=0.38), 250 nm (x=0.62) and 200 nm (x=1.0: a–C) as
determined by XRR.
2.2. Characterization
The structural properties of the thin films were determined
by high-resolution XRD analysis. In order to avoid
diffraction signal from the Si substrate, grazing incidence
(GI) XRD measurements were carried out on a PANanalytical
EMPYREAN using a Cu Kαradiation source, and a parallel
beam geometry with a 2incidence angle to avoid substrate
peaks and minimize the influence of texture. Each XRD scan
was performed with 0.1resolution, 0.05step length with a
total of 1800 points for 6 h.
HR-TEM, and selected area electron diffraction (SAED)
were performed by using a Tecnai G220 U-Twin 200kV
FEGTEM microscope. Cross-section samples were mechan-
ically polished, and ion milled to electron transparency by a
Gatan Precision Ion Polishing System (PIPS).
The chemical compositions of the films were determined
by x-ray photoelectron spectroscopy (XPS) using a Physical
Systems Quantum 2000 spectrometer with monochromatic Al
Kαradiation. Depth profiles of the films were acquired by
rastered Ar+-ion sputter etching over an area of 2 mm2×2mm
with ions being accelerated by the potential difference of 4 kV.
The high-resolution scans of the selected peaks were acquired
after 6, 30, and 45 min of Ar+-ion sputter etching with ions
being accelerated by the potential difference of 4500 kV and
200 kV, respectively. The XPS analysis area was set to a diam-
eter of 200 μm and the step size to 0.05 eV with a base pressure
of 109Pa during all measurements. The peak fitting was made
by Voigt shape functions to account for the energy resolution of
the instrument and chemical disorder (Gaussian part) and the
lifetime width of the photoionization process (Lorenzian part).
Raman scattering spectroscopy was used in order to
correlate the nanostructuring, and sp2/sp3ratio of the films to
the carbon concentration. The Raman spectra were acquired
at room temperature in the range 800–1900 cm1in the back
scatteringconfiguration using UV 325 nm laser excitation.
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
X-ray absorption spectroscopy (XAS) measurements were
performed in total fluorescence yield (TFY) mode at the undu-
lator beamline I511-3 at MAX II (MAX-IV Laboratory, Lund,
Sweden), comprising a 49-pole undulator, and a modified SX-
700 plane grating monochromator [16]. The measurements
were made at a base pressure lower than 6.7·107Pa. The
XAS spectra were measured at 5grazing incidence angle from
the surface plane and a detection angle of 30from the inci-
dent photon direction. All samples were measured in the same
geometry with energy resolutions of 0.2 eV and 0.1 eV at the
Ni 2p and C 1sabsorption edges, respectively. The XAS spec-
tra were normalized to the step before and after the absorp-
tion edges and corrected for background and self-absorption
effects [17] with the program XANDA [18] in figures 6and 7.
Cross-sectional scanning electron microscopy (SEM)
images were obtained in a LEO 1550 microscope using
accelerating voltages of 15 kV in in-lens imaging mode. The
obtained images were used for thickness measurements, and
structural analysis of the coatings.
Sheet resistance measurements were made with a four-
point probe ‘4-dimensions’ 280C. For each sample, four
readings were made with a different 4-sensor orientation
around the center of the sample. As a final value of the
electrical resistivity, a mean value over four measurements was
made. Each set of measurements on a sample showed similar
values indicating negligible influence from surface oxide.
3. Results
3.1. X-ray diffraction (XRD) and high-resolution transmission
electron microscopy (HR-TEM)
Figure 1(a) shows x-ray diffractograms (XRD) performed to
characterize the microstructure of the nickel carbide Ni1xCx
films for x=0.05,0.16,0.38,and 0.62. The XRD data were
refined by the Rietveld method using the MAUD program [20].
Five peaks in the top diffractogram (x=0.05) are indexed
as fcc-NiCxstructure (space group Fm-3m) with a lattice
parameter of a=3.610(1)Å. This lattice parameter is larger
than that of fcc-Ni metal (3.524 Å) [19], but smaller than
that of fcc-NiC (4.077 Å) [20]. Interpolation of these values
yields an estimated phase composition of fcc-NiCx, where
x=0.23–0.30. However, the lattice parameter also depends
on the nano structured grain size of fcc-NiCxthat has larger
cell parameters than in the case of bulk materials. Therefore,
x0.30 for fcc-NiCxin the 0.05 sample represents an upper
limit of the composition.
The other marked reflections in the top diffractogram can
be indexed by either hcp-Ni or rhombohedral Ni3C structures.
From a structure point of view, rhombohedral Ni3C has the
same Ni position as hcp-Ni but the ordered interstitial C atoms
create additional reflections 01-12 and 1-10-4, which are ab-
sent for the hcp-Ni structure. Thus, the rhombohedral structure
can be excluded based on the absence of 01-12 and 1-10-4 re-
flections [2124]. Furthermore, compared to the pure hcp-Ni
phase, the slight peak shift to low angle indicates an expansion
of the lattice, due to the carbon occupation. With increas-
ing carbon content, the cell parameters are further increased
Figure 1. (a) X-ray diffractograms of Ni1xCxfilms with C content
ranging from 4.9 to 61.8 at%. (b) Enlargement of the XRD data of
the 16.3% C sample at low angles in comparison to Rietveld
refinement of hcp-Ni and rhomb. Ni3C0.5with a similar composition
as the 16.3% sample.
(see the second diffractogram of the 16.3 at% C sample in
figure 1(a)). Thus, the phase with hcp-Ni structure should be
termed as hcp-NiCy(y < 1) instead of hcp-Ni.
The diffractograms are further analyzed with Rietveld re-
finement. Both hcp-NiCyand rhombohedral Ni3C structures
were used in the refinement. For comparison, the same carbon
content was used for the two structure models, i.e. hcp-NiC0.67
and rhombohedral Ni3C0.5. The refined results in figure 1(b)
shows that the 01-12 and 1-10-4 peaks are present in rhom-
bohedral Ni3C0.5but absent in hcp-NiC0.67. The simulated
diffractogram of the hcp-NiC0.67 agrees well with our ex-
perimental data (see figure 1(b)). Moreover, the refinement
gives rise to the accurate lattice parameters of hcp-NiCyas
a=2.611(2)Å and c=4.328(7)Å for the 4.9 at% C sam-
ple and a=2.653(7)Å, c=4.337(2)Å for the 16.3 at%
C sample, respectively. Figure 1(a) also shows that with in-
crease of carbon content from 4.9 to 16.3 at%, the fcc-NiCx
disappears and single phase of hcp-NiCyforms. Applying
Scherrer’s equation for the 16.3 at% C sample yields a grain
size of 23 nm for all diffraction peaks in agreement with the
average grain size in the TEM images. A further increase of
carbon content to 37.9 at% and above leads to amorphous-like
structures, as shown in figure 1(a).
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 2. HR-TEM micrographs of low (left), and high (right) resolution with corresponding SAED patterns for the Ni1xCxfilms with
increasing xvalues, top to bottom: (a,b) 0.05; (c,d) 0.16; (e,f) 0.38; and (g,h) 0.62. The diffraction rings for the fcc-NiCx(2 0 0) and
hcp-NiCy(0 11), (0 0 2), (01 0) reflections are indicated in the SAED.
The structural evolution of the films with composition is
also observed by HRTEM and SAED as shown in figure 2. The
sharp SAED in figure 2(c) clearly has an absence of 012 and
104 reflections, which confirms the hcp-NiCystructure rather
than rhombohedral Ni3C structure. No significant texture is
observed in the SAED. The films with low carbon contents
(x=0.05, and 0.16) are polycrystalline, and the sharp dots
of reflections in the corresponding SAED pattern show that
the film with 16.3 at% C consists of hcp-NiCy, while the
film with 4.9 at% C contains two phases: hcp-NiCy, and
cubic fcc-NiCx(the 200 reflection is consistent with XRD).
In contrast, the films with higher carbon contents (38, and 62
at% C) consist of Ni-rich nanocrystalline domains surrounded
by amorphous carbon-rich matrix domains. The average grain
size of the nanocrystalline domains is approximately 3–5 nm
(x=0.38 and 0.62). Although it is difficult to identify the
exact phases for these x-ray amorphous samples with high
carbon content, the XRD peak intensity distribution profiles
indicate that the strongest broad structure at 2θ=42.6
includes both hcp-NiCy010, 002 and 011 reflections, and a
fcc-NiCx111 reflection, respectively. The broad structure at
2θ=82is formed by fcc-NiCx200, and hcp-NiCy110 and
103 reflections. Thus, as observed by the HR-TEM in figure 2,
the samples with high carbon contents of 38 and 62 at% likely
consist of both cubic fcc-NiCx, and hcp-NiCynanocrystallites
both with a small grain size, approximately 3–5 nm.
3.2. X-ray photoemission spectroscopy (XPS)
Figure 3shows C 1s core-level XPS spectra of the four Ni1xCx
films with x=0.05, 0.16, 0.38, and 0.62. As observed, at least
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 3. C1s XPS spectra of the Ni1xCxfilms with carbon
content ranging from 5 to 62 at%. The deconvoluted peaks at
283.3 eV and 285.3eV, indicated by the dashed vertical lines
corresponds to carbidic NiC carbon in Ni–C bonds, and free
carbon in C–C bonds, respectively. A third structure is identified
at 283.9 eV, and can be associated with charge-transfer C–Ni
bonds or C–C in sp2hybridized bonds.
three peaks are required to deconvolute the spectra. A peak
at 283.3 eV can be assigned to Ni–C bonds, while a second
peak at 285.3 eV can be assigned to sp3hybridized carbon
(C–C-sp3)[25,26]. Between these two peaks, a third feature is
clearly present. The intensity of the middle peak increases with
carbon content and is also shifted from about 283.9 eV in the
most Ni-rich film to about 284.5 eV in the most C-rich film.
Most likely, several types of carbon are contributing to this
feature. Firstly, sp2-hybridized carbon is known to exhibit a
peak at about 0.9 eV lower binding energy than sp3-hybridized
carbon, i.e. at about 284.0 eV [25]. It is well known that binary
sputter-deposited metal carbide films are often nanocomposites
with a carbide phase in an amorphous carbon (a–C) matrix.
Previous C 1sXPS studies on the a–C phase have shown a
mixture of sp2—and sp3-hybridized carbon. Consequently,
it is most likely that a part of the intensity of the feature at
283.9–284.5 eV originates from free carbon in an a–C matrix.
Secondly, studies on sputter-deposited Me–C films have also
shown an additional Me–C feature at a slightly higher binding
energy [3] that can be due to sputter damage of the metal
carbide grains. Thirdly, a contribution originates from surface
Me atoms in the carbide grains. This is caused by charge-
transfer effects where charge is transferred from the metal
surface atoms to the more electronegative carbon atoms in the
a–C matrix [7]. In nanocomposites with very small grains or
domains, the relative amount of surface atoms is large and
will show up as a high-energy shoulder on the main C 1s
Me–C peak (denoted Me–C)[9]. However, for the Ni1xCx
films, it is impossible to deconvolute the feature at 283.9–
284.5 eV into separate C–C (sp2) and Ni–Cpeaks. However,
a comparison with Ti1xCx,Cr
1xCxand Fe1xCxfilms show
that the Me–Ccontribution is small in XPS compared to
the C–C (sp2) peak [27,28]. For this reason, we assign the
entire peak at 283.9–284.5 eV to C in a–C, although it will
give a slight overestimation of the amount of the C–C (sp2)
phase compared to the NiC carbide phases. The XPS data
supports the TEM and XRD studies and confirms that the
films consist of at least two phases: fcc-NiCxand hcp-NiCy
carbide nanocrystallites dispersed in an amorphous carbon (a–
C) phase. The intense C–Ni peak for x=0.05 and 0.16 is an
indication of the localized character of the Ni–C bonds in the
Figure 4shows the relative amount of the carbide and
a–C phases as a function of total carbon content (assuming
that the Ni–Ccontribution to 283.9–284.5 eV peak can be
neglected). As can be seen, the relative amount of the a–C
phase increases non-linearly with the total carbon content.
The total composition analysis in table 1is valid under the
assumption that the total photoemission cross-section in all of
the samples is constant for carbon. The combined carbon in the
fcc-NiCxand hcp-NiCycarbide phases can now be estimated
using the data in figure 4, as presented in table 1. The analysis
show that the carbon content of the carbide phase strongly
increase with the total carbon content from 15.7 at% (0.16
at% total), 36 at% (0.38 at% total) to 60 at% (0.62 at% total).
However, the estimated carbon content in the carbide phase
represents a lower limit since the contribution of Ni–C* has
been neglected in the analysis of the C 1s spectra. The variation
of the carbon content in the NiC carbide phase is consistent
with the small dispersion of the Ni 2p3/2XPS peak position
from 852.7 eV (x=0.05), 852.9 eV (x=0.16), 853.0 eV
(x=0.38), to 853.1 eV (x=0.62).
3.3. Raman spectroscopy
Figure 5shows carbon Raman spectra of the Ni1xCxfilms
with xvalues of 0.16, 0.38, and 0.62, in comparison to pure
amorphous carbon (a–C). As the Raman scattering cross sec-
tion from C–Ni is low, the Raman spectra are dominated by the
segregated part of the carbon in the compounds. The two band
components in the spectra, the disordered (D), and graphite (G)
peaks, were deconvoluted by Voigt shape functions. In pure
graphite, the vibrational mode that gives rise to the G-band
is known to be due to the relative motion of sp2hybridized C
atoms while the D band is due to the breathing vibrational mode
of the six-membered rings [29]. The energies of the D and the
G bands have an almost constant position around 1410cm1
and 1570 cm1, respectively. Both peaks are slightly shifted
together with respect to their positions for the pure a–C of
1384 cm1and 1585 cm1, respectively. The ID/IGheight ra-
tios of the films are 1.53, 1.59, 1.36 while for the a–C film, the
ratio is lower (0.87) due to a more graphitic character of the
carbon bonds [30]. These ratios approximately correspond to
sp2fractions of 0.68, 0.70, and 0.65 while it is lower for a–C
(0.59). The predominant sp2hybridization is consistent with
the observations in the XPS spectra of the same samples.
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 4. Relative amount of C–Ni, and C–C bonds determined as the proportions of the areas fitted on the XPS C 1s peak. The two curved
lines are guidelines for the eye.
Tab le 1. Composition of the Ni–C films for x=0.05,0.16,0.38,and 0.62. The amount of carbon in the carbide phase and the sp2fractions
were determined by integrating the areas under the corresponding peak structures in C 1s XPS spectra. The sp2fractions in Raman were
estimated from [28].
Total composition Ni0.95C0.05 Ni0.84C0.16 Ni0.62 C0.38 Ni0.38C0.62 a–C
at% C in NiCyphase 5.0 15.7 36 60 100
XPS sp2fraction 0.64 0.80 0.69 0.70 —
Raman sp2fraction 0.77 0.79 0.89 0.71
C1sXAS π/[π+σ] 0.42 0.53 0.60 0.56 0.72
Figure 5. Raman spectra for the carbon peak of the Ni1xCxfilms
for x=0.16,0.38,0.62, and amorphous carbon (a–C). The two
vertical dashed lines indicate the disorder (D), and graphite (G)
peaks of the fitted peak components [29,30].
3.4. Ni 2p x-ray absorption spectroscopy
Figure 6shows Ni 2p XAS spectra of the 3d, and 4s conduction
bands following the Ni 2p3/2,1/23ddipole transitions of
the Ni1xCxfilms with different carbon content in comparison
to Ni metal. The Ni 2p XAS spectra mainly represent the
nickel contribution in the fcc-NiCxand hcp-NiCycarbide
phases. The main peak structures are associated with the Ni
2p3/2and the 2p1/2core-shell spin–orbit splitting of 17.3 eV.
A comparison of the spectra shows four interesting effects: (i)
the intensity of the main 2p3/2peak decreases with carbon
content. The Ni 2p XAS intensity is proportional to the
unoccupied 3d states, and the intensity trend indicates that
the Ni 3d electron density decreases around the absorbing
Ni atoms for higher carbon concentration. The intensity of
the normally sharp Ni 2p3/2XAS peak in pure crystalline Ni
phase, is largely suppressed by the broadening and distribution
of different types of chemical bonds. (ii) For comparison,
the XAS spectrum of fcc Ni metal (x=0.0) has narrower,
and more intense 2p3/2, and 2p1/2absorption peaks, whereas
the XAS spectra of the carbon-containing films are broader
and shifted by 0.6 eV towards higher photon energy. This
energy shift is an indication of higher ionicity of Ni as a
result of charge-transfer from Ni to C. (iii) For x=0.16, a
pronounced double-peak structure with 1.4 eV splitting from
the main peak at the 2p3/2peak is observed. The double-peak
feature is similar to the t2g-egcrystal field splitting observed
in TiC nanocomposites [16]. It is a signature of a change
in orbital occupation to the hcp crystal structure while the
sharp single-peak feature of Ni metal is a signature of fcc
structure (cubic). For x=0.05, the double-peak structure
has essentially vanished in comparison to at x=0.16 due
to the superposition of the strong fcc contribution. (iv) The
6 eV feature [19,31] above the main 2p3/2peak is prominent
in Ni metal x=0.0 that is associated with electron correlation
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 6. Ni 2p TFY-XAS spectra of Ni1xCxfor different C contents in comparison to bulk Ni (x=0.0).
effects and narrow-band phenomena [32]. The intensity of the
6 eV feature is very low in the Ni1xCxfilms in comparison to
Ni metal even at x=0.05 due to more delocalized bands.
A comparison of the spectral shapes at different carbon
contents shows that the 2p3/2/2p1/2branching ratio estimated
by the peak height is largest (3.0) for the lowest carbon content
(x=0.05) and is similar as for Ni metal. With increasing
carbon content, the branching ratio decreases to 2.3 (x=0.16),
1.8 (x=0.38) and 1.3 (x=0.62). Integration of peak areas by
Gaussian functions give the same trend as a comparison of the
peak heights but yields a higher branching ratio for Ni metal
than for the other samples. A lower 2p3/2/2p1/2branching
ratio is an indication of higher ionicity (lower conductivity) for
the highest carbon content [3335]. However, the 2p3/2/2p1/2
branching ratio is a result of the ionicity for Ni, mainly in the
fcc-NiCxand hcp-NiCy, carbide components, and not for the
entire film. For the higher carbon contents, when the main
part of the film consists of C-rich matrix areas, this phase
determines the resistivity.
3.5. C 1sx-ray absorption spectroscopy
Figure 7shows C 1sXAS spectra of the Ni1xCxfilms probing
the unoccupied C 2p conduction bands as a superposition of
the NiC carbide phases, and the changes in the C matrix phase
with composition. The first peak structure (1) at 285 eV is
associated with empty πstates, and the higher states (3–6)
above 290eV are associated with unoccupied σstates. The
empty πorbitals (1) consist of the sum of two contributions
in Ni1xCx: (i) sp2(C=C), and sp1hybridized C states in
the amorphous carbon phase, and (ii) C 2p–Ni 3d hybridized
states in the fcc-NiCxand hcp-NiCycarbide phases. The
peak at 288.5 eV is also due to C 2p-Ni 3d hybridization
with a superimposed contribution from the carbon phase
[16]. The energy region above 290eV is known to originate
from sp3hybridized (C–C) σ* resonances, where the peak
(4) at 291.5 eV forms a shape resonance with multielectron
excitations towards higher energies [16]. The most intense
structure shows highest intensity for x=0.05 that originates
from sp3hybridized σstates. Additional σstates (5), (6)
at 295 eV and 298 eV are also associated with sp3bonding.
Contrary to the case of TiC [16], there is no pre-peak below
the πpeak at 283.3 eV for NiC. The integrated π/[π+σ]
intensity ratio was calculated by fitting a step-edge background
with a Gaussian function to each peak following the procedure
in [36,37] in order to not overestimate the σcontribution.
We assumed that πpeaks occur below and σpeaks above
290 eV as indicated in figure 7. This analysis method gives an
estimation of the relative amount of π(sp2,sp
content in the samples as shown in table 1. The fraction of
sp2is smallest for x=0.05 and highest for a–C, following a
similar trend as the XPS and the Raman results.
3.6. Resistivity measurements
Figure 8(a) shows the electrical resistivity in the Ni1xCx
films as a function of carbon content x. Compared to the
electrical resistivity of 6.93 μcm [38] for Ni metal, the
small introduction of C of 3 at% increases the resistivity
to 30 μcm. However, it is well known that metallic thin
films display higher electrical resistivity compared to the bulk
metals [38,39]. When increasing the C content another 2
at% C, the resistivity doubles to 62 μcm, and continues
approximately linearly to 150 μcm at 16.3 at% C. Above
this carbon content, the samples transform from polycrystalline
to amorphous and the resistivity increases more by a factor
of two to 400 μcm. Above 20.7 at% C, the resistivity
increases approximately linearly up to a maximum value of
775 μcm for the C content of 61.8 at%. The increase of
electrical resistivity with increasing C content in the films is
correlated to the increased amount of C–C bonding (figure 4),
and the interstitial incorporation of inter-bonded C atoms
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 7. C1sTFY-XAS spectra of Ni1xCxfor different C contents compared to amorphous carbon (a–C, x=1.0).
Figure 8. (a) Resistivity of the Ni1xCxfilms depending on the C content as determined from sheet resistance measurements by a four-point
probe and film thickness determined by SEM. (b) The resistivity is plotted as a function of the relative amount of C bound in the C–Ni
bonds. The dashed least-square fitted curves are guides for the eye.
between the Ni lattice sites forming the fcc-NiCxand hcp-
NiCycarbide phases. The a–C phase is known to be a very
poor conductor compared toNi metal. As the XPS analysis
shows contributions from C–C bonds for all the investigated
films, an increase of resistivity is not only governed by the
increase of the amorphous C phase but is also influenced by
the crystallinity. Figure 8(b) shows the exponential decrease of
electrical resistivity with the increasing proportion of the NiC
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
carbide phases. The symmetry of the C–Ni component of the
C1s XPS peak also suggests a low electrical conductivity of
the carbide. However, as shown by the XPS analysis, there is
only a small amount of C incorporated into Ni. Therefore, the
NiC phases could be regarded as a solid solution [39,40] with
a low amount of C solved into Ni rather than as an ordinary
carbide phase. Thus, the C–Ni component of the structure is
mostly metallic, leaving the a–C matrix phase to determine the
general trend in the electrical resistivity of the films.
4. Discussion
The most significant difference between Ni1xCxas compared
to other late transition-metal carbides such as Cr1xCx[4], and
Fe1xCx[41], is the precipitation of Ni-based phases in the
form of nanocrystals, while Cr and Fe form mainly completely
amorphous films for a large range of compositions. In our
combined analysis of XRD and HR-TEM, we find that at low
carbon content (x=0.05), the Ni1xCxfilms consist of hcp-
NiCynanocrystals with a smaller contribution (25–30%) of
fcc-NiCynanocrystals, where y0.30. For low carbon
contents (x=0.05 and 0.16), the estimated average grain
size is relatively large, 10–20nm.
Previous investigations of hcp-Ni nanocrystals are lacking
a material composition analysis [6,7], but pure hcp-Ni metal
is known to be unstable or metastable. To our knowledge,
only one previous experiment showed that crystallites of hcp-
Ni metal could be synthesized, and it was easily transformed
into fcc-Ni metal when its size was larger than 5 nm [11]. Most
of the reported hcp-Ni metals are likely stabilized by carbon.
The Ni atom in rhombohedral-Ni3C and hcp-Ni structures
occupy exactly the same positions and yield identical XRD
data at high angle (>38) and this is the reason why it has
not been identified before. Recently, Schaefer et al pointed
out that it is possible to experimentally distinguish hcp-Ni and
rhombohedral-Ni3C structures by using low-angle XRD. In
addition, He and Schaefer claimed that there exists no hcp-Ni
because they inferred that all the previously reported hexagonal
Ni carbides contained carbon, presented 01-12 and 1-10-4
reflections, and should therefore be described as rhombohedral
Ni3C[13,14]. However, in our sputter-deposited Ni1xCx
samples, the absence of 01-12, and 1-10-4 reflections indicates
that the superstructure of rhombohedral Ni3C is not formed,
and instead, a hcp-Ni structure occurs. It should be noted that
our hcp-Ni structure does contain C and its cell parameters
depend on the carbon content. This is consistent with previous
works, showing that hcp-Ni is stabilized by carbon. Thus, our
hcp-Ni phase should be described as hcp-NiCyinstead of hcp-
Ni or rhombohedral-Ni3C1x. Uhlig et al also found a carbon
containing Ni structure [9]. However, the absence of low-angle
reflections indicates that their films consist of hcp-Ni with a
carbon content rather than Ni3C.
The XPS and XAS measurements confirm the structure
to be carbidic and do not show spectral profiles of metallic
Ni. From a structural point of view, the difference between
hcp-Ni3C and rhombohedral Ni3C is the carbon position:
ordered interstitial C in rhombohedral Ni3C and disordered
interstitial C in hcp-Ni3C. The formation of hcp-NiCyinstead
of rhombohedral-Ni3C may be due to the non-equilibrium
sputtering process. Moreover, at low carbon content, hcp-
NiCyor rhombohedral-Ni3C is likely more stable than fcc-
NiCx. Further theoretical calculations will be performed to
verify this hypothesis, and consequently give an interpretation
as to why NiCxform crystalline phases, whereas CrCxand
FeCxform amorphous phases at low carbon content.
The difference in XPS binding energy of the C–Ni peak
in comparison to the C–C peak (2.0 eV) is due to the different
types of bonding environments. A small low-energy shift of
0.15 eV between the samples with 4.9%. This observation is
consistent with a small XPS high-energy shift of 0.25 eV at the
Ni 2p3/2edge for these crystalline samples. As the structure of
the samples change from crystalline to amorphous at 38 and 62
at Fe, one would also expect a smaller chemical shift in the case
of Ni carbides. This scenario with a smaller chemical shift for
the late transition metal carbides is consistent when comparing
to Ti–C (2.5eV) [5] but not for Fe–C (1.7 eV) [41] and Cr–C
(1.5 eV) [3], where it is smaller. In this respect, XAS gives
important complementary information to XPS about charge-
transfer effects. The general intensity trend in the 2p branching
ratio of the Ni XAS spectra is a signature of charge-transfer
from Ni to C that is largest for the sample that contains most C
(i.e. x=0.62). The variation in intensity of the unoccupied
states reflects changes in orbital occupation and bonding of
the atoms at the carbide/matrix interface between crystallites,
and amorphous domains. It can be assumed that charge-
transfer occurs within the fcc-NiCxand hcp-NiCynanocrystal
carbide phases, but more significant across the carbide/matrix
interface with the surrounding amorphous C-phase or between
nanocrystals, that depends on the nanocrystalline size.
Moreover, the 6 eV feature in the Ni XAS spectra
that signifies electron correlation effects and narrow-band
phenomena in metallic Ni [19,31] is washed out in the Ni1xCx
samples due to the Ni 3d-C 2p orbital overlap that changes the
properties of Ni already at very low carbon content. Thus,
the spectral profiles of the Ni1xCxsamples exhibit carbide
signatures and exclude metallic nickel. Furthermore, for the
carbon content of x=0.16, the ligand-field type of splitting
by 1.4 eV that occurs in Ni XAS signifies a change in the local
coordination and orbital occupation with the formation of the
single-phase hcp-NiCycarbide phase. The most stable NiC
phase is cubic [15], but from the combined shape of the Ni
2p (no crystal-field splitting) and C 1sXAS spectra (absence
of pre-peak), this is excluded and consistent with the XRD
observations. Using surface-sensitive TEY measurements,
Choo et al [42] associated the double-structure in Ni 2p XAS
of hcp Ni with surface oxidation. With bulk-sensitive TFY-
XAS, we find that this feature is due to the intrinsic hcp-Ni
Since charge-transfer effects are clearly observed in bulk-
sensitive XAS, this contribution is also expected in the more
surface sensitive XPS spectra. Although this contribution is
difficult to separate in the XPS data, part of the third peak
between the C–C, and C–Fe peaks should be associated with
charge-transfer effects at the interfaces between nanocrystals.
The size and the number of fcc-NiCxand hcp-NiCynanocrys-
tals affect the amount of interface, and charge-transfer between
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
the domains. As observed by the XPS analysis, we find that the
carbon content in the carbide phase varies significantly with
the total carbon content (table 1). However, the amount of
sp2-fraction (0.6–0.8) as observed in XPS, Raman and XAS is
higher in all samples in comparison to a–C and does not change
significantly when the samples transform from crystalline to
amorphous. On the other hand, the trend in the resistivity de-
pends on the carbon content as well as the crystallinity. To
energetically explain why particular nanocrystals form in the
Ni–C system and not in other late transition metal carbides,
further experimental and theoretical work will be carried out
including the effect of magnetic properties.
5. Conclusions
Magnetron sputtered nanocomposite Ni1xCxfilms were
investigated for a large composition range (0.05 x0.62).
We discovered a novel hcp-NiCyphase, and show how it is
different from rhombohedral Ni3C. At low carbon content (4.9
at%), the Ni1xCxfilm consists of hcp-NiCyand fcc-NiCx
nanoparticles with an average grain size of 10–20nm as ob-
served by high-resolution x-ray diffraction and transmission
electron microscopy. With increasing carbon content (16 at%),
single-phase hcp-NiCyis formed also with an average grain
size of 10–20 nm. A double structure in the x-ray absorption
spectra reveals a change in the orbital occupation and bond-
ing for hcp-NiCyin comparisonto the other samples. A fur-
ther increase of carbon content to 38 and 62 at%, transforms
the films to complex x-ray amorphous materials with a mix-
ture of randomly-oriented short-range ordered hcp-NiCyand
fcc-NiCxnanodomain structures surrounded by an increasing
amount of amorphous carbon-rich matrix. X-ray photoelec-
tron and x-ray absorption spectroscopy analyses reveal that
interbonding states between the nanocrystallites, and domain
structures represent a third type of phase that increases with
carbon content. The general trend of increased electrical resis-
tivity with increasing carbon content in the films is correlated
to the increased amount of C–C bonding observed in x-ray
photoelectron spectroscopy. As the C–Ni phase component of
the structure is metallic, the carbon matrix phase determines
the electrical resistivity of the films when the films are amor-
phous. We also find an increase of the resistivity by more than
a factor of two when the samples transform from crystalline to
We would like to thank the staff at the MAX IV Laboratory
for experimental support, and Jill Sundberg, UU, for help with
the Raman measurements. The work was supported by the
Swedish Research Council (VR) Linnaeus, and Project Grants.
M M, U J and J L also acknowledge support from the SSF
synergy grant FUNCASE Functional Carbides and Advanced
Surface Engineering.
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Supplementary resource (1)

... The Ni-C peaks presented at high decomposition temperatures are induced by the decomposition of amorphous metastable Ni 3 C to Ni and carbon atoms, whereas the high temperature facilitates the mixing of partial carbon and Ni atoms at the boundaries of crystalline. This process eventually forms an intermediate, solid solution with a small amount of carbon solved into nickel atoms rather than the ordinary carbide phase [39]. ...
... At the same time, the metastable state Ni3C species between Ni NP and carbonaceous support decompose into metallic Ni and graphene [38], and the formed carbon layers gradually cover the surface of Ni NPs. The Ni-C peaks presented at high decomposition temperatures are induced by the decompo sition of amorphous metastable Ni3C to Ni and carbon atoms, whereas the high tempera ture facilitates the mixing of partial carbon and Ni atoms at the boundaries of crystalline This process eventually forms an intermediate, solid solution with a small amount of car bon solved into nickel atoms rather than the ordinary carbide phase [39]. To verify the assistance of carbon support, the C 1s XPS spectra are investigated. ...
... At about 400 • C, the carboxyl carbon content is negligible, and the broken coordination results in an increased content of sp3 carbon (Table S3). This conversion is also accompanied by the formation of the C species at 283.8 eV, which corresponds to the interstitial carbon in the amorphous Ni 3 C species [39]. Moreover, the location of C−O shifted to 286.6 eV, suggesting the C−O sort changes from ligand phenolic hydroxyl to oxygenic groups on carbon. ...
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... 34,35 The C 1s showed binding energy of >290 eV due to sp 3 -C C ϭ* resonance. 36 Thus the XPS results exhibit that MoO3/NiO/PdO/Pd were functionalized with organic species. It proved the extract to be a good stabilizing/capping agent. ...
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... Similar interactions are also observed in Mg-Ni-rGO system. However, an additional Ni-C interaction is observed at 283.5-284.5 eV (Zhao et al. 2011;Adamska et al. 2012; Furlan et al. 2014). The Ni-O-C and Ni-C interactions after ball milling and hydrogenation respectively could be responsible for the low onset temperature during dehydrogenation from Mg-Ni-rGO nanocomposite (Fig. 1b). ...
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Facing the environmental crisis, new technologies are needed to sustain our society. In this context, this thesis aims to describe the properties and applications of carbon-based sustainable materials. In particular, it reports the synthesis and characterization of a wide set of porous carbonaceous materials with high nitrogen content obtained from nucleobases. These materials are used as cathodes for Li-ion capacitors, and a major focus is put on the cathode preparation, highlighting the oxidation resistance of nucleobase-derived materials. Furthermore, their catalytic properties for acid/base and redox reactions are described, pointing to the role of nitrogen speciation on their surfaces. Finally, these materials are used as supports for highly dispersed nickel loading, activating the materials for carbon dioxide electroreduction.
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Direct methane PCFCs are promising electrochemical devices that address the technical and economic challenges associated with using pure hydrogen, such as the high cost of green production, transportation, and long-term storage. However, Ni, a conventional monometallic catalyst has sluggish reaction kinetics and a low tolerance for carbon cocking under CH 4 operation, limiting its wider applications. Herein, we develop a self-assembled Ni-Rh bimetallic catalyst through Ni exsolution and Rh surface decoration in a fuel electrode. The Ni-Rh bimetallic catalyst shows remarkably high catalytic activity with an exceptional performance of ~ 0.50 W/cm ² at 500 ℃. Moreover, the catalyst significantly improves the stability with a degradation rate of 0.02%/h at 500 ℃; this value is ~ 20-fold lower than that of conventional PCFC (0.4%/h). Synchrotron-based in situ X-ray photoelectron spectroscopy reveals that the Ni-Rh bimetallic catalyst initiates a self-carbon cleaning process due to its high-water dissociation reaction, allowing sustainable operation.
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The same method used to determine {sigma}{sup *} excitations in the x-ray-absorption spectra of organic molecules is applied to analyze the structure of the {sigma}{sup *} band in the x-ray absorption spectra of amorphous carbon (a-C) films. The analysis assumes that only a {sigma} bond interaction to first neighbors is relevant to explain the structure of the {sigma}{sup *} band in a-C films. This is justified by the local character of the x-ray absorption probe and the short- range order existing in these films. The identification of the different {sigma}{sup *} components is based on the dependence of the {sigma}{sup *} binding energy with bond distance. The {sigma}{sup *} band is built up by summing the components resulting from the possible different types of {sigma} bonds in a-C. This method serves to separate the {pi}{sup *} states from the {sigma}{sup *} states and to identify the kind of chemical bonds existing between carbon atoms. This analysis yields a proportion of sp{sup 3}-bonded atoms of 60% in a film with a density of about 2.9 g/cm{sup 3}, which is a value closer to what is expected from theoretical calculations. The analysis identifies a component at about 288.5 eV which is associated to strained {sigma} bonds of the type of the existing in sp{sup 3} bonded rings like cyclopropene. Raman and photoemission spectroscopies help in the interpretation of the x-ray absorption spectra and the assignment of the {sigma}{sup *} components.
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Two different types of amorphous carbon films were deposited on Si substrates, with film hardness of 22 GPa and 40 GPa, by pulsed laser evaporation of graphite targets. The x-ray photoemission spectra (XPS) of the C 1s core level in these films shown two components at 284.3±0.1 eV and 285.2±0.1 eV, which were identified with the sp2 and sp3 hybrids forms of carbon. The sp3/sp2 concentration ratio deduced from the area of the components had a value of 2/5 for the harder amorphous carbon film and 1/4 for the softer. Upon annealing the harder film at different temperatures, the sp3/sp2 ratio remained nearly constant up to about 900 K and then decreased until reaching a value of zero above 1100 K. The C 1s core level shifted 0.3±0.1 eV toward lower binding energy in the films for annealing temperatures above 900 K. This shift was correlated with an increase in the asymmetry of the C 1s XPS spectra and of the density of states at the Fermi level, as observed by ultraviolet photoemission spectroscopy. There was no detectable π plasmon in the harder films below 900 K, despite the presence of sp2 atoms and π bonds at those temperatures.
Controlling the performance of structures and components of all sizes and shapes through the use of engineered coatings has long been a key strategy in materials processing and technological design. The ever-increasing sophistication of en- neered coatings and the rapid trend toward producing increasingly smaller devices with greater demands on their fabrication, properties and performance have led to signi?cant progress in the science and technology of coatings, particularly in the last decade or two. Nanostructured coatings constitute a major area of sci- ti?c exploration and technological pursuit in this development. Withcharacteristic structural length scales on the order of a few nanometers to tens of nanometers, nanostructured coatings provide potential opportunities to enhance dramatically performance by offering, in many situations, extraordinary strength and hardness, unprecedented resistance to damage from tribological contact, and improvements in a number of functional properties. At the same time, there are critical issues and challenges in optimizing these properties with ?aw tolerance, interfacial adhesion and other nonmechanical considerations, depending on the coating systems and applications. Nanostructured coatings demand study in a highly interdisciplinary research arena which encompasses: surface and interface science study of defects modern characterization methodologies cutting-edge experimental developments to deposit,synthesize, conso- date, observe as well as chemically and mechanically probe materials at the atomic and molecular length scales state-of-the-art computational simulation techniques for developing - sightsintomaterialbehaviourattheatomicscalewhichcannotbeobtained in some cases from experiments alone The interdisclipinary nature of the subject has made it a rich playing ?eld for scienti?c innovation and technological progress.
Local structures of diamond-like carbon (DLC) films formed by various methods were studied by near-edge X-ray absorption fine-structure (NEXAFS) spectroscopy. The DLC films are characterized by the sp2/sp 3 ratio, which influences the mechanical and electronic properties. NEXAFS spectroscopy is sensitive to the sp2/sp3 ratio, because the isolated peak corresponding to the 1s→π* resonance transition can be observed. Carbon K-edge NEXAFS spectra for DLC thin films, which were synthesized by various methods, were measured using the total electron yield mode in the range of 275eV-320eV. A peak due to the coupling of carbon with oxygen was observed in the spectra of some DLC films, whereas it was not observed in the spectra of hydrogenated carbon films formed by RF sputtering. The obtained relative sp2 contents of the DLC films were distributed in the range of ≈20%. The minimum sp2/sp3 ratio was obtained from DLC films formed by vacuum arc deposition from graphite, and large sp2/sp3 ratios were obtained from DLC films formed by plasma chemical vapor deposition from hydrocarbons. The local structure of a DLC film was concluded to depend on the synthesis method, and in particular, the carbon source material.
Nanoparticles of elemental nickel underpin a large number of magnetic and catalytic applications, and the possibility of tuning these properties via the formation of different allotropes is intriguing. While bulk elemental nickel adopts a face centered cubic (fcc) structure, a growing number of reports suggest that colloidal nickel nanoparticles can crystallize in the metastable hexagonal close packed (hcp) structure. However, there is some disagreement in the literature concerning the formation of hcp-Ni, particularly with respect to the crystallographically-related Ni3C phase. Most notable is a range of lattice constants and magnetic properties that have been attributed to hcp-Ni. Here, we show that reaction time can be used to tune the carbon content of a Ni3C1-x solid solution. Importantly, colloidal nanoparticles of Ni3C1-x can help to experimentally rationalize the range of lattice constants and magnetic properties reported for hcp-Ni and Ni3C, effectively bridging these two end-member systems. All samples, including those isolated immediately upon reduction of Ni2+ to Ni0, contained some carbon, as evidenced by XRD, XPS, TGA, DSC, TEM, and SQUID magnetometry. As reaction time increases, the average carbon content increases, and this correlates with a systematic increase in unit cell volume and a systematic decrease in saturation magnetization. These results also provide a straightforward pathway for tuning the magnetic properties of isomorphous Ni nanoparticles.
The structure of metallic glasses (MGs) has been a long-standing mystery. On the one hand, MGs are amorphous materials with no long-range structural order; on the other hand, topological and chemical short-to-medium range order is expected to be pronounced in these alloys, due to their high atomic packing density and the varying chemical affinity between the constituent elements. The unique internal structure of MGs underlies their interesting properties, which render MGs potentially useful for various applications. While more and more glass-forming alloys have been developed in recent years, fundamental knowledge on the structural aspect of MGs remains seriously lacking. For example, how atoms pack on the short-to-medium range, how the structure differs in different MGs and changes with composition, temperature, and processing history, and more importantly, how the structure influences the properties of MGs, are still unresolved questions. In this paper, we review the tremendous efforts over the past 50 years devoted to unraveling the atomic-level structure of MGs and the structural origin of their unique behaviors. Emphasis will be placed on the progress made in recent years, including advances in structural characterization and analysis of prototypical MGs, general structural models and fundamental principles, and the correlations of thermodynamic, kinetic, and mechanical properties with the MG structures. Some widely observed property–property correlations in MGs are also examined from the structural perspective. The insights summarized are shown to shed light on many intriguing behaviors of the MG-forming alloys and expected to impact the development of MGs. Outstanding questions in this important research area will also be outlined.
Thin films in the Cr–C system with carbon content of 25–85 at.% have been deposited using non-reactive DC magnetron sputtering from elemental targets. Analyses with X-ray diffraction and transmission electron microscopy confirm that the films are completely amorphous. Also, annealing experiment show that the films had not crystallized at 500 °C. Furthermore, X-ray spectroscopy and Raman spectroscopy show that the films consist of two phases, an amorphous CrCx phase and an amorphous carbon (a-C) phase. The presence of two amorphous phases is also supported by the electrochemical analysis, which shows that oxidation of both chromium and carbon contributes to the total current in the passive region. The relative amounts of these amorphous phases influence the film properties. Typically, lower carbon content with less a-C phase leads to harder films with higher Young’s modulus and lower resistivity. The results also show that both films have lower currents in the passive region compared to the uncoated 316L steel substrate. Finally, our results were compared with literature data from both reactively and non-reactively sputtered chromium carbide films. The comparison reveals that non-reactive sputtering tend to favour the formation of amorphous films and also influence e.g. the sp2/sp3 ratio of the a-C phase.
Thin films based on transition-metal carbides exhibit many interesting physical and chemical properties making them attractive for a variety of applications. The most widely used method to produce metal carbide films with specific properties at reduced deposition temperatures is sputter deposition. A large number of papers in this field have been published during the last decades, showing that large variations in structure and properties can be obtained. This review will summarise the literature on sputter-deposited carbide films based on chemical aspects of the various elements in the films. By considering the chemical affinities (primarily towards carbon) and structural preferences of different elements, it is possible to understand trends in structure of binary transition-metal carbides and the ternary materials based on these carbides. These trends in chemical affinity and structure will also directly affect the growth process during sputter deposition. A fundamental chemical perspective of the transition-metal carbides and their alloying elements is essential to obtain control of the material structure (from the atomic level), and thereby its properties and performance. This review covers a wide range of materials: binary transition-metal carbides and their nanocomposites with amorphous carbon; the effect of alloying carbide-based materials with a third element (mainly elements from groups 3 through 14); as well as the amorphous binary and ternary materials from these elements deposited under specific conditions or at certain compositional ranges. Furthermore, the review will also emphasise important aspects regarding materials characterisation which may affect the interpretation of data such as beam-induced crystallisation and sputter-damage during surface analysis.