![Raman spectra for the carbon peak of the Ni1−xCx films 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].](https://www.researchgate.net/profile/Martin-Magnuson/publication/265846112/figure/fig5/AS:11431281208416132@1701393346368/Raman-spectra-for-the-carbon-peak-of-the-Ni1-xCx-films-for-x016-038-062-and_Q320.jpg)
Content uploaded by Martin Magnuson
Author content
All content in this area was uploaded by Martin Magnuson on Oct 21, 2020
Content may be subject to copyright.
This content has been downloaded from IOPscience. Please scroll down to see the full text.
Download details:
IP Address: 130.236.82.7
This content was downloaded on 20/09/2014 at 11:34
Please note that terms and conditions apply.
Crystallization characteristics and chemical bonding properties of nickel carbide thin film
nanocomposites
View the table of contents for this issue, or go to the journal homepage for more
2014 J. Phys.: Condens. Matter 26 415501
(http://iopscience.iop.org/0953-8984/26/41/415501)
Home Search Collections Journals About Contact us My IOPscience
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¨
oping,
Sweden
2Department of Chemistry, Ångstr¨
om Laboratory, Uppsala University, Box 538, SE-751 21 Uppsala,
Sweden
E-mail: Martin.Magnuson@ifm.liu.se
Received 22 June 2014, revised 29 July 2014
Accepted for publication 9 August 2014
Published 19 September 2014
Abstract
The crystal structure and chemical bonding of magnetron-sputtering deposited nickel carbide
Ni1−xCx(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 Ni1−xCxfilms with Ni3C
crystallites embedded in an amorphous Ni1−xCxphase 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 Ni1−xCxfilms. 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 Ni1−xCxfilms
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 Ni1−xCxthin films were
deposited in a UHV chamber with a base pressure of 10−9Pa
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 2◦incidence angle to avoid substrate
peaks and minimize the influence of texture. Each XRD scan
was performed with 0.1◦resolution, 0.05◦step 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
2
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 10−9Pa 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 cm−1in the back
scatteringconfiguration using UV 325 nm laser excitation.
2
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 6.7·10−7Pa. The
XAS spectra were measured at 5◦grazing incidence angle from
the surface plane and a detection angle of 30◦from 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 Ni1−xCx
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 [21–24]. 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 Ni1−xCxfilms 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).
3
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 Ni1−xCxfilms 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=82◦is 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 Ni1−xCx
films with x=0.05, 0.16, 0.38, and 0.62. As observed, at least
4
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 3. C1s XPS spectra of the Ni1−xCxfilms 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 Ni1−xCx
films, it is impossible to deconvolute the feature at 283.9–
284.5 eV into separate C–C (sp2) and Ni–C∗peaks. However,
a comparison with Ti1−xCx,Cr
1−xCxand Fe1−xCxfilms show
that the Me–C∗contribution 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
nanocrystallites.
Figure 4shows the relative amount of the carbide and
a–C phases as a function of total carbon content (assuming
that the Ni–C∗contribution 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 Ni1−xCxfilms
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 1410cm−1
and 1570 cm−1, respectively. Both peaks are slightly shifted
together with respect to their positions for the pure a–C of
1384 cm−1and 1585 cm−1, 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.
5
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 Ni1−xCxfilms
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 2p3/2,1/2→3ddipole transitions of
the Ni1−xCxfilms 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
6
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 6. Ni 2p TFY-XAS spectra of Ni1−xCxfor 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 Ni1−xCxfilms 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 [33–35]. 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 Ni1−xCxfilms 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 π∗states, and the higher states (3–6)
above 290eV are associated with unoccupied σ∗states. The
empty π∗orbitals (1) consist of the sum of two contributions
in Ni1−xCx: (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 σ∗states. Additional σ∗states (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 π∗peak 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 π∗peaks occur below and σ∗peaks above
290 eV as indicated in figure 7. This analysis method gives an
estimation of the relative amount of π∗(sp2,sp
1hybridization
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 Ni1−xCx
films as a function of carbon content x. Compared to the
electrical resistivity of 6.93 μcm [38] for Ni metal, the
small introduction of C of ∼3 at% increases the resistivity
to ∼30 μcm. 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 μcm, and continues
approximately linearly to ∼150 μcm 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 μcm. Above 20.7 at% C, the resistivity
increases approximately linearly up to a maximum value of
∼775 μcm 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
7
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
Figure 7. C1sTFY-XAS spectra of Ni1−xCxfor different C contents compared to amorphous carbon (a–C, x=1.0).
Figure 8. (a) Resistivity of the Ni1−xCxfilms 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
8
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 Ni1−xCxas compared
to other late transition-metal carbides such as Cr1−xCx[4], and
Fe1−xCx[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 Ni1−xCxfilms 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 Ni1−xCx
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-Ni3C1−x. 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 Ni1−xCx
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 Ni1−xCxsamples 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
structure.
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
9
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 Ni1−xCxfilms 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 Ni1−xCxfilm 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
amorphous.
Acknowledgments
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.
References
[1] Cavaleiro A and DeHossonJTHM(ed) 2006
Nanostructured Coatings (Berlin: Springer)
[2] Cheng Y Q and Ma E 2011 Prog. Mater. Sci. 56 379–473
[3] Andersson M, H¨
ogstr¨
om J, Urbonaite S, Furlan A, Nyholm L
and Jansson U 2012 Vacuum 86 1408
[4] Magnuson M, Andersson M, Lu J, Hultman L and Jansson U
2012 J. Phys.: Condens. Matter 24 225004
[5] Jansson U and Lewin E 2013 Thin Solid Films 536 1
[6] Laidani N, Calliari L, Speranza G, Micheli V and
Galvanetto E 1998 Surf. Coat. Technol. 100–1 116
[7] Shi J, Hashiba Y and Nittono O 2001 J. Mater. Sci. 36 343
[8] Bauer-Grosse E and Aouni A 2007 J. Non-Cryst. Solids
353 3644
[9] Uhlig S, Struis R, Schmid-Engel H, Bock J, Probst A-C,
Freitag-Weber O, Zizak I, Chernikov R and Schultes G
2013 Diamond Relat. Mater. 34 25
[10] Kopplert R et al 2012 Diamond Relat. Mater. 25 50
[11] Tian W, Sun H P, Pan X Q, Yu J H, Yeadon M,
Boothroyd C B, Feng Y P, Lukaszew R A and Clarke R
2005 Appl. Phys. Lett. 86 131915
[12] Ingham B, Gaston N, Fahy K, Chin X Y, Dotzler C J, Rees E,
Haslam G, Barber Z H, Burstein G T and Ryanet M P
2012 J. Phys. Chem. C116 6159
[13] He L 2010 J. Magn. Magn. Mater. 322 1991
[14] Schaefer Z L, Weeber K M, Misra R, Schiffer P and
Schaak R E 2011 Chem. Mater. 23 2475–80
[15] Jacobson B and Westgren A 1933 Z. Phys. Chem. Abt. B
20 361
[16] Magnuson M, Lewin E, Hultman L and Jansson U 2009
Phys. Rev. B80 235108
[17] Magnuson M, Wassdahl N and Nordgren J 1997 Phys. Rev. B
56 12238
[18] Klementiev K V 2001 J. Phys. D: Appl. Phys. 34 209–17
Klementiev K V XANES dactyloscope for Windows,
freeware: www.xafs.org/software/ViperAndXanda (2014)
www.cells.es/old/Beamlines/CLAESS/software/ (2013)
[19] Magnuson M, Wassdahl N, Nilsson A, F¨
ohlisch A,
Nordgren J and Mårtensson N 1998 Phys. Rev. B58 3677
[20] Lutterotti L 2010 Nucl. Instrum. Methods Phys. Res. B
268 334
[21] Suh I K, Ohta H and Waseda Y 1988 J. Mater. Sci. 23 757
[22] Gibson J S, Uddin J, Cundari T R, BodifordNKand
Wilson A K 2010 J. Phys.: Condens. Matter 22 4455031
[23] Nagakura S 1958 J. Phys. Soc. Japan 13 1005
[24] Wijeyesekera S D and Hoffmann R 1984 Organometallics
3949
[25] Diaz J, Paolicelli G, Ferrer S and Comin F 1996 Phys. Rev. B
54 8064
[26] KovacsGYJet al 2008 Thin Solid Films 516 7942
[27] Popov C et al 2004 Diamond Relat. Mater. 13 2071
[28] WilsonJIBet al 2001 J. Electron Spectrosc. 121 183
[29] Ferrari A C and Robertson J 2000 Phys. Rev. B61 14095–107
[30] Schwan J, Ulrich S, Batori V, Ehrhardt H and SilvaSRP
1996 J. Appl. Phys. 80 440
[31] Magnuson M, Nilsson A, Weinelt M and Mårtensson N 1998
Phys. Rev. B60 2436
[32] Magnuson M, Butorin S M, Agui A and Nordgren J 2002 J.
Phys.: Condens. Matter 14 3669
[33] Magnuson M, Wilhelmsson O, Mattesini M, Li S, Ahuja R,
Eriksson O, H¨
ogberg H, Hultman L and Jansson U 2008
Phys. Rev. B78 035117
[34] Magnuson M, Wilhelmsson O, Palmquist J-P, Jansson U,
Mattesini M, Li S, Ahuja R and Eriksson O 2006 Phys.
Rev. B74 195108
[35] Laskowski R and Blaha P 2010 Phys. Rev. B82 205104
[36] Diaz J, Anders S, Zhou X, Moler E J, KellarSAand
Hussain Z 2001 Phys. Rev. B64 125204
10
J. Phys.: Condens. Matter 26 (2014) 415501 A Furlan et al
[37] Saikubo A, Kanda K, Niibe M and Matsui S 2006
New Diamond Front. Carbon Technol. 16 235
[38] Haynes W M (ed) 2011 CRC Handbook of Chemistry and
Physics 92nd edn (Boca Raton, FL: CRC Press)
www.hbcnetbase.com
[39] Portnoi V K, Leonov A V, Mudretsova S N and Fedotov S A
2010 Phys. Met. Metallogr. 109 153
[40] Ryabtsev S I, Bashev V F, Belkin A I and Ryabtsev A S 2006
Phys. Met. Metallogr. 102 305
[41] Furlan A, Lu J, Jansson U, Hultman L and Magnuson M
2014 submitted to J. Phys. Cond. Mat.
[42] Choo S, Lee K, Jo Y, Yoon S-M, Choi J-Y, Kim J-Y,
Park J-H, Lee K-J, Lee J-H and Jung M-H 2011
J. Nanosci. Nanotechnol. 11 6126
11
A preview of this full-text is provided by IOP Publishing.
Content available from Journal of Physics: Condensed Matter
This content is subject to copyright. Terms and conditions apply.
