Post-print of: Surface and Coatings Technology volume 206, issues 8–9, 15 January 2012,
pages 2525–2534
Analysis of multifunctional titanium oxycarbide films as a function of oxygen addition
J.M. Chappé (a), A.C. Fernandes (a), C. Moura (a), E. Alves (b), N.P. Barradas (b), N. Martin (c),
J.P. Espinós (d), F. Vaz (a)
a Centro de Física, Universidade do Minho, 4710-057 Braga, Portugal
b Departamento de Física, Instituto Tecnológico Nuclear, E.N. 10, 2686-953 Sacavém, Portugal
c Institut FEMTO-ST, UMR 6174, CNRS UFC ENSMM UTBM, 32, Avenue de l'Observatoire,
25044 BESANCON Cedex, France
d Instituto de Ciencia de Materiales de Sevilla (CSIC-Univ. Sevilla). Avda. Américo Vespucio 46,
41092 Sevilla, Spain
Abstract
Reactive magnetron sputtering was used to deposit titanium oxycarbide thin films. The overall
set of results showed that the oxygen flow rate, and thus the composition of the atmosphere
in the deposition chamber, controls the composition of the titanium oxycarbide thin films
obtained by reactive sputtering. Rutherford Backscattering Spectroscopy analysis revealed the
existence of three major types of films, indexed to their particular composition ratios. A
detailed study by X-ray photoelectron spectroscopy was carried out in order to characterize
the evolution of the Ti—C, C—O and C—C bonds induced by the increase of the oxygen partial
pressure, which was found to be closely related with the different zones of composition that
were suggested. Micro-Raman spectroscopy and X-ray diffraction measurements allowed
describing the complex nature of the film structure, namely in what concerns different phases
and their evolution, texture phenomena and grain size evolution as a function of the particular
composition and film types (different zones). Electrical conductivity revealed a transition from
a metallic to a semi-conducting behavior as a function of the oxygen concentration in the films,
in good agreement with the different zones that were suggested. Similarly, optical properties
supported this gradual change and for oxygen contents higher than 67 at.%, the films exhibited
typical reflectance of insulator materials (interferences) in the UV, visible and near IR regions.
Keywords
X-ray photoelectron spectroscopy (XPS); Carbides; Electrical conductivity; Optical properties
1. Introduction
1
The continuous progress in surface and thin film technology is largely connected to the
adjustment of structural and chemical properties of several types of films to actual
applications. Besides the chemical species that will be present in the thin film, and the
particular set of experimental parameters that are to be chosen, the relative amounts of each
element that compose a film have to be considered as essential factors for determining the
microstructural properties and chemical binding conditions of the growing film [1] and [2], and
thus governing the overall material (thin film) behavior.
In recent years there has been an increasing interest in elucidating the mechanisms of the
tailoring of a thin film system, using a combination of different chemical elements. The use of
transition-metal nitrides in engines, machines, tools and other wear-resistant components
steadily increases and has achieved by far the highest level of commercial success. In this
respect, TiN and (Ti,M)N coatings with M = B, C, Al, Cr, etc. are being widely used, mainly
because of their unique combination of properties, namely high mechanical resistance, among
others [3], [4], [5], [6], [7] and [8]. Other more recent examples include the so-called
multifunctional thin film systems, where the oxynitrides, MeOxNy, and oxycarbides, MeCxOy,
(where Me stands for transition metal) have been attracting particular interest [9], [10], [11],
[12], [13], [14], [15] and [16]. The multifunctionality of such systems arises from the fact that
the presence of oxygen allows the tailoring of film properties between those of metallic
nitrides/carbides, MeN/MeC, and those of the correspondent insulating oxides, MeOx. The
possibility to tune the oxide/nitride or oxide/carbide ratio allows one to change the material
bonding characteristics and consequently the structural arrangement between oxide and
nitride/carbide and hence the overall electric, optical, mechanical and tribological properties
of the materials. By this mixing approach (achieved simply by adding increasing amounts of
oxygen to well-known nitrides/carbides), a relatively simple method can be used to obtain
several different material responses. Furthermore, little is known about the right structure of
metallic carboxides, especially on TiCxOy thin films and their resulting properties [17] and [18].
Such films are identified as interesting compounds exhibiting vacancies on both metal and
non-metal sublattices over a wide range of compositions [19]. These structural defects may
lead to unique combination of properties that may become attractive for various applications
[20] and [21].
Anyway, and in spite of the apparent simplicity, the understanding of this structural
arrangement and its relationships to its formation mechanism is still very incomplete, and the
available studies are mostly related to the properties of the material itself, and less concerned
to the particular architectures of the material (micro/nano)structures. In the case of
nitrides/carbides and oxides mixing, the understanding of the structural and bonding
arrangements are of major importance, not only in the material science fundamental
knowledge point of view, but above all in the sense that a detailed understanding of the
structural and bonding evolution is of fundamental importance for the preparation of new
materials and novel devices that exhibit improved physical properties.
It will be on these bonding arrangements that the present paper will be focused on, firstly
showing the multi-phase system of titanium oxycarbide, TiCxOy, thin films, prepared by
reactive magnetron sputtering, and then establishing the correlations between the films'
2
deposition conditions and the effective bonding evolutions and the correspondent variation of
their structural, electrical and optical properties.
1.1. Experimental details
The TiCxOy films were deposited onto single crystal silicon (100) substrates, by reactive DC
magnetron sputtering, in a laboratory-size deposition system. It consists of two vertically
opposed rectangular magnetrons (unbalanced of type 2), in a closed field configuration. Films
were prepared in rotation mode (using a constant speed of seven rotations per minute). Only
one target was used during all depositions with the substrate holder positioned at 70 mm from
the target. A DC power supply was connected to the Ti target (200 × 100 mm2 — 99.6 at.%
purity), containing 12 cylindrical carbon pellets (10 mm diameter) placed in the preferential
eroded zone. The used current density was 100 A.m− 2. A gas atmosphere composed of argon
(working gas) and oxygen (the reactive gas) was used. The Ar flow was kept constant at 60
sccm and the oxygen gas flow rate varied from 0.5 to 10 sccm (corresponding to a partial
pressure variation from 7.8 × 10− 3 to 8.6 × 10− 2 Pa). The residual pressure in the chamber
was always below 2 × 10− 3 Pa before the introduction of the working and the reactive gases,
while the working pressure varied only slightly from approximately 4 × 10− 1 to 5 × 10− 1 Pa
during the depositions. The depositions were carried out with the substrates in the grounded
condition. The deposition temperature was close to 200 °C (a heating resistance positioned at
80 mm from the substrate holder and a thermocouple was placed close to the surface of the
substrate holder on plasma side to measure the temperature immediately after stopping the
discharge).
The atomic composition of the as-deposited samples was measured by Rutherford
Backscattering Spectroscopy (RBS) using (1.4, 1.75) MeV and 2 MeV for the proton and 4He
beams, respectively. The scattering angles were 140° (standard detector, IBM geometry) and
180° (annular detector), tilt angles 0° and 30°. Composition profiles for the as-deposited
samples were generated using the software code NDF [22]. For the 14N, 16O and 28Si data,
the cross-sections given by Ramos et al. were used [23]. The analyzed area was about 0.5 × 0.5
mm2. The uncertainty in the C and N concentrations is around 5 at.%, and around 2 at.% for
the Ti. In addition, for some samples, Particle Induced X-ray Emission (PIXE) measurements
were performed to check for impurities. Ball cratering tests were used to measure the
thickness of the samples (used for deposition rate determination). The characterization of the
bond-types was carried out in a XPS spectrometer from VG (ESCALAB 210). An
unmonochromatized Mg Kα (1253.6 eV) source was used for the measurements. Both XPS
core (O1s, Ti2p and C1s) and valence band spectra were recorded after successive sputteringclean processes with Ar+ ions (3.5 kV, 30° incidence), until a steady state in the surface
composition was reached (around 100 min sputtering). Spectra were taken in the ΔE constant
mode with a pass energy of 50 eV. They were energy calibrated by the position of the Ti2p3/2
in TiO2 (initial state, before sputtering) at 458.5 eV, and the Fermi level at 0 eV for the rest of
situations. The Shirley algorithm was applied for background subtraction [24], and all the plots
have been normalized to the intensity of the O1s signal.
The structure and phase distribution of the coatings were accessed by X-ray diffraction (XRD),
using a conventional Philips PW 1710 diffractometer, operating with Cu Kα radiation, in a
3
Bragg–Brentano configuration. The XRD patterns were deconvoluted and fitted with a Voigt
function to determine the structural characteristics of the films, such as the peak position (2θ)
and the full width at half maximum (FWHM), which were then used to access the lattice
parameters and grain sizes. Raman measurements were performed at room temperature with
a Jobin-Yvon (T64000) triple monochromator, equipped with a liquid nitrogen cooled charge
couple device (CCD) detector, with a resolution better than 1 cm− 1. The excitation line, 488
nm, of an argon ion laser was focused onto the sample using a × 100 MS Plan objective of an
Olympus Microscope BHSM, in a backscattering geometry. The spectra were obtained with a
measured power of about 0.5 mW on the sample, to avoid heating. Electrical conductivity was
measured on glass substrates at room temperature and as a function of temperature. The van
der Pauw configuration was involved. Optical reflectance spectra were recorded in the range
of wavelengths 250–2500 nm at 10° of incidence using a 900 Lambda Perkin–Ellmer
spectrophotometer equipped with the URA (Universal Reflectance Accessory) device.
2. Results and discussion
2.1. Chemical composition
Fig. 1 shows the RBS composition results of the TiCxOy samples (with thickness ranging from
1.3 to 3.2 μm), plotted as a function of the oxygen partial pressure. Fig. 1a) shows the variation
of the atomic concentration of the film's elements, while Fig. 1b) represents the
correspondent atomic ratios, rO = CO/CTi, rC = CC/CTi and rOC = (CO + CC)/CTi.
From Fig. 1a) the first conclusion to be drawn is that there are major differences in the
evolution of the film's elemental concentrations, namely that of O in comparison to those of C
and Ti. In fact, the increase in the oxygen partial pressure up to a value of PO2 = 3.2 × 10− 2 Pa,
leads to the synthesis of films with an expected continuous increase of the O content (from
about 10 at.% up to approximately 55 at.%), while Ti revealed a clear tendency to decrease
significantly, from approximately 72 at.% down to 38 at.%. This is a very common behavior that
is observed when preparing thin films with increasing reactive gas flow partial pressures
(nitrides, oxides, among others), resulting from the well-known target oxidation/poisoning
effect. There is an increase of the oxide contamination layer at the target surface, which has a
considerably lower sputtering yield compared to that of a metallic Ti. The process works in the
compound sputtering mode leading to a significant reduction of the sputtering yield. In
addition, increasing the oxygen gas supply into the process, increases the oxygen partial
pressure as well. More oxygen species react with sputtered titanium atoms and with particles
depositing on the growing film. As a result, the oxygen concentration in the films rises and
affects the structural properties and the physical behaviors of deposited TiCO thin films.
Regarding the O and Ti content evolution, the composition results revealed some distinct
zones within the used oxygen partial pressures. First of all, the films prepared with oxygen
partial pressures up to 1.5 × 10− 2 Pa show a sudden and reverse change of O and Ti
concentrations: O contents rising from 10 to close to 40 at.%, whereas that of Ti dropping from
72 to 50 at.%. Films prepared with oxygen partial pressures between 1.5 × 10− 2 and 3.2 × 10−
2 Pa exhibit a more reduced variation of chemical composition. Oxygen content changes from
40 up to 55 at.% and Ti reduces down to 38 at.%. Regarding the films prepared with oxygen
partial pressures PO2 ≥ 3.5 × 10− 2 Pa, the composition analysis revealed a quite different
4
variation tendency. All elements seem to reach a compositional steady state. In this high
oxygen partial pressure range, the films have all a very low amount of C (less than 3 at.%),
while O and Ti are about 67 and 30 at.%, respectively.
In order to have more accurate differences in composition within the different films, as well as
to check for some possible correlations between the composition and the compounds that
may be formed, Fig. 1b) shows the different atomic ratios, rO = CO/CTi, rC = CC/CTi and rOC =
(CO + CC)/CTi. Following the composition variations mentioned above, Fig. 1b) allows dividing
the prepared films into different compositional zones. A first zone can be indexed to the films
prepared with oxygen partial pressures up to 1.5 × 10− 2 Pa. The composition analysis revealed
the highest increase of O and highest decrease of Ti. It obviously corresponds to a rapid
increase of the atomic ratio rO, which grows from values close to 0 up to approximately 0.8 in
this first zone, providing some evidence for a possible increase of the O bonds with Ti.
This initial increase of the oxygen amount in the coatings is due to the well-known high affinity
of metal atoms, such as those of titanium, towards oxygen rather than to carbon [25]. Taking
into account the available thermochemistry data for titanium oxides and carbides, a
progressive oxidation of the films would be expected as long as sufficient oxygen is available,
since the formation of a titanium–oxygen bond is energetically more favorable (ΔHf°TiO2: −
672.4 kJ.mol− 1[26], ΔHf°TiO: − 520 kJ.mol− 1) than that of a titanium–carbon bond (ΔHf°TiC: −
183 kJ.mol− 1[26]). Anyway, as both elemental and atomic ratio evolutions reveal in this first
zone (which will be noted hereafter as zone I), (Fig. 1a) and b)), the O and C concentrations
remain lower than or approximately equal to the concentration of Ti, and thus the kinetics of
the reaction may play a dominant role, inducing that both elements most likely form mixed
bonds with Ti. Anyway, since no ion bombardment was used in the preparation of the films
(except some possible negative oxygen ions, which are neglected at first), and thus the atom
mobility might be too reduced to insure complete phase segregations (carbide and oxide
ones), it may be admitted that some C can also form some bonds with Ti. Therefore it is
possible that phase mixtures may occur, including elemental ones (such as C within grain
boundaries — nanocomposite-type structures, etc.). In fact, one must also consider that during
the sputtering process, some clusters or groups of C atoms may arrive on the growing film,
providing the conditions to form graphite-like species.
For oxygen partial pressures between 1.9 × 10− 2 and 3.2 × 10− 2 Pa, a second zone is defined,
called zone II in Fig. 1b). This figure shows that, although rather smoothly, the rO ratio keeps
growing, varying between roughly 1 and 1.5. It provides the conditions to form overstoichiometric-type compounds (rOC varies between 1.3 and 1.6), or even some increasing
mixture of different phases. Anyway, the values of oxygen are still relatively below those
required to form stoichiometric TiO2 (rO = 2), and thus some possible formation of mixed
carboxide phases, Ti(C,O), may be induced, but now with clearly more oxygen than carbon. The
fact that the ratio rC is relatively low in this second zone (below 0.26), may suggest that O is
playing the fundamental role in the formation of any crystalline (or non-crystalline) films, but
with amounts that may give rise to some difficulties to form any of the common crystalline
structures (both monoxides and dioxides). It is also important to note that this indexation into
two distinct zones is basically done to separate the films in terms of their stoichiometry, since
5
in fact both Fig. 1a) and b) reveal some kind of continuous behavior, which seems particularly
evident in the evolution of rO and rOC in Fig. 1b).
Finally, the samples prepared with oxygen partial pressures above 3.5 × 10− 2 Pa, zone III in
Fig. 1b), revealed atomic ratios rO and rOC around 2, meaning that the films are now in
conditions close to those of TiO2-type compounds. Moreover, the atomic ratio rC is rather low
(below 0.1), which associated with the TiO2-type stoichiometric condition of the films and the
thermodynamics data, may induce some tendency to form TiO2-type films, with C occupying
some secondary phases in the film structure (isolated at the grain boundaries or even mixed in
the oxide phases). An important note about this third zone is that the actual content of oxygen
seems to be even higher than that of TiO2 (rO is round 2.2), which may result from the
uncertainty in the composition measurements, as well as some defects in the oxide lattices
that may be formed.
2.2. Bonding states: X-ray photoelectron spectroscopy (XPS)
In order to follow the influence of the chemical composition variations, Fig. 1a) and b), on the
type of bonds that are formed in the films, and specially to check if there are differences in the
bonding states within each compositional zone mentioned in the previous section, a detailed
study of the chemical changes in the films was carried out, using X-ray photoelectron
spectroscopy (XPS).
As an introductory note and beyond the powerful nature of this technique in checking the
chemical environments present in materials (especially in the first superficial layers), it is worth
to notice that one of most common sources of errors in XPS interpretation comes from the socalled surface contamination. This contamination normally comprises the first few dozens of
angstroms of the surface and is at the origin of several deficient conclusions about material
behavior. Enyashi and Ivanovskii [27] studied the surface chemical environments on titanium
oxide and titanium carbide thin films and suggested that oxygen adsorbs dissociatively at 300
K, and it is immediately incorporated into the first few atomic layers of the solid. Furthermore,
Leng et al. [28] showed that the appearance of this oxide-like layer (either stoichiometric or
even slightly sub-stoichiometric) is also commonly observed in several other film systems, such
as those of nitrides, carbides, etc., just to mention a few. In the case of titanium compounds,
this contamination layer is normally made of titanium oxide [29], which may lead to a
misunderstanding in the identification of chemical states, since the same species may also be
present in the “bulk”. In order to overcome this problem, a sputter-cleaning procedure was
adopted in this work. The ion bombardment of the ternary system Ti–C–O causes the
preferential sputtering of oxygen and carbon, inducing a metal rich surface, promoting the
appearance of sub-stoichiometric phases where titanium is in a oxidation state lower than + 4
[30]. This behavior is often claimed to result in the formation of several different titanium suboxide phases [28] and [30]. In order to minimize this problem, an energy sputter cleaning
process was performed with Ar+ ions (3.5 kV, 30° incidence), during relatively large periods of
time (up to 105 min), until a compositional stationary regime was achieved. Fig. 2 shows the
result of such sputter cleaning process, carried out in one of the samples studied in this work
(sample prepared with 6 sccm, PO2 = 3.7 × 10− 2 Pa). This figure shows that after sputtering
6
for 105 min, a steady state was reached, where composition of the surface layer does not
change any more by the sputter cleaning.
2.2.1. Ti2p, C1s and O1s core level spectra
In terms of the detailed analysis of the TiCxOy samples prepared within the framework of this
study, a set of four samples was used for this chemical bonding analysis, selected according to
their particular composition and located in each of the three distinct compositional zones.
From zone I, sample TiC0.24O0.14 (rOC = 0.38) was selected, while from zone II, samples
TiC0.26O1.04 (rOC = 1.30) and TiC0.19O1.44 (rOC = 1.62) were selected according to their
particular positioning at the boundaries between the two other zones. Finally, sample
TiC0.09O2.27 (rOC = 2.36) was selected from zone III.
Fig. 3 shows the Ti2p spectrum for the four selected samples prepared within each of the three
distinct zones, as well as a Ti2p spectrum of a “pure” TiO2 thin film, prepared with similar
conditions as those used for the preparation of the TiCxOy samples (same sputtering
parameters, not no C pellets were included in the Ti target), and subjected also to the same
sputtering treatment (3.5 kV, 30°), for reference. As expected, and in agreement with several
works in the literature [28] and [30], the result of the sputter-cleaning is the appearance of a
mixture of oxidation states for Ti. The vertical lines show the binding energies for the two
extreme species: Ti+ 4 (in TiO2, peaks at about 458.5 (2p3/2) and 464.2 eV (2p1/2)) and Ti0 (in
metallic Ti, peaks at 453.8 eV (2p3/2) and ~ 460 eV (2p1/2)), as well as, most likely, from two
others species, Ti+ 3 and Ti+ 2. The binding energy for these two species is between that of Ti+
4 and Ti0, but the exact values depend on the chemical nature of the atoms (ions) bonded to
Ti+ 2 and Ti+ 3 (in this case, C, or O) and their coordination number. Binding energy values
found in the literature for Ti+ 3 or Ti+ 2 must be taken as a first approximation value, since a
precise knowledge of the composition and structure of the surface layer of materials used as
references is needed.
The appearance of these new species can have two reasons: the film compositions (which are
graded from the atmosphere corroded surface to the virgin bulk), and the sputtering process
itself, as O and C atoms initially bonded to Ti are preferentially removed, as it happened for the
pure TiO2. As a first approximation these results are consistent with some kind of transition
behaviors in what concerns samples from zones I and II, and with a TiO2-like behavior of
samples from zone III, as suggested by the composition results. Anyway, it is fair to claim that
the results obtained in the samples from zone II are not significantly different from those
obtained in the sample from zone I, where only a significant reduction in the Ti0 peak intensity
is clearly observed, while their position remains almost the same.
Regarding the relative positions, the peaks around 454.9 and ~ 461 eV can be attributed to TiC
(Ti2p3/2 and 2p1/2, respectively), and/or TiO, which are claimed to appear for very similar
binding energies [31], [32], [33], [34] and [35]. Although not clearly visible due to the mixture
of oxidation states mentioned above, the peaks at about 458.5 (2p3/2) and 464.2 eV (2p1/2)
must be attributed to Ti—O bonds in TiO2. Nevertheless, and following the compositional
changes observed in Fig. 1a) and b), the relative abundances of the Ti+ n species are
dependent on the analyzed sample, or more precisely, on the samples zone: the lower the rOC
atomic ratio (lower oxygen amounts), the higher is the concentration of Ti0 species at the
7
steady state. In fact, a closer look to Fig. 3 reveals that the main peaks that might be attributed
to both elemental Ti and/or TiC and TiO, are clearly visible in the sample from zone I, but
decrease significantly in intensity when moving to samples from zone II and especially from
that of zone III. It is also important to note the behavior of the two samples from zone II,
where it seems that the traces for these Ti—Ti bonds are now significantly reduced in this
zone. This is consistent with increasing amounts of oxygen (especially within zone II, where rO
varies from about 1 to approximately 1.4) and its high reactivity towards titanium, while
keeping relatively low amounts of carbon (rC varies from about 0.25 to approximately 0.2). It
would be expected that a change from a Ti—C bonding dominant zone (especially within zone
I) to a more Ti—O bonding-type would occur in this second zone.
Concerning the O1s peak, Fig. 4, the spectra of the four TiCxOy samples (including that of
“pure” TiO2) seem quite identical, whatever the deposition conditions or the composition.
Again, the main difficulty arises from the fact that the binding energies of O—Ti bonds from
both TiO2 and TiO-type compounds are quite similar (varying between 529.4 and 530.1 eV),
where actually some authors claim the values around 530.1 eV for both types [34], [36], [37],
[38] and [39]. Thus the fact that all O1s peaks are so similar for all samples suggests that, in the
ion altered surface layer of the films, there is mainly a unique chemical state for O, whatever
the deposition condition of the samples. However, the absence of TiO2-type bonds in the Ti2p
spectrum for the sample in zone I (Fig. 3), indicates that they must really be monoxide Ti—Otype bonds, which then evolve to TiO2-type ones in zone II (and especially within zone III). In
fact, the O1s binding energy is the same as the one for oxide species in TiO2 and for TiOx (x <
2) obtained by Ar+ sputtering of TiO2[40]. The peak is slightly asymmetrical in the high binding
energy side, but this asymmetry is also found in sputtered TiO2. However, if the samples
contain minor amounts of O bond to C (O—C) (see fitting treatment of C1s signals), they
should also appear at this position.
Finally, regarding the last present element, C, a very large signal due to hydrocarbon like or
graphite-like species is found initially. As the intensity of this signal decreases very fast by
sputtering, it can be concluded that it is cause by contamination. After sputtering for 85 min
(at the steady state), two clear signals are found at the vicinity of the binding energy values of
281.3 (± 0.1) and 283.5 (± 0.1) eV, which can be ascribed to carbide (carbon bonded to metallic
Ti [35] and [36]) and graphite (or carbon bond to carbon [41] and [42]), respectively.
Accordingly, there are some points that deserve to be noted. First, C—Ti bonds decrease
significantly from zone I to zone II, as expected. Second, there are still some traces of C—Ti
bonds, in spite of the very low content of carbon (between 2 and 3 at.%) present in the films
from zone III. And third, and somehow surprisingly, there is always evidence of the existence
of C—C bonds, even for films from zone III (Fig. 5). In the case of zone I, this might be
interesting due to the under-stoichiometric nature of the films. Considering these last results,
one might assume that not only a change/mixture of Ti—C and Ti—O bonds may exist, but also
some “free” C, which would induce the possibility that some kind of nanocomposite material is
forming, in which a Ti(C,O) phase would co-exist with another one of carbon. A final note
concerns the “pure” TiO2 sample, which reveals almost no traces of C—Ti bonds, which
confirms the “cleaning” of the surface contamination layer, and thus reinforcing all
observations that were made previously about the composition of the “bulk” layer. This
effective sputter-cleaning process also enables us to make some extra quantitative analysis, as
8
demonstrated by Fig. 6, related to the elemental concentration obtained by the XPS peak
fitting.
From the results plotted in Fig. 6, the first important note is that, although the two techniques
contain information from different regions of the sample (XPS gives information of only the
very first surface layers, while RBS gives detailed information on the composition of the bulk of
the films), the obtained results are relatively close in spite of some discrepancies for films
produced with oxygen partial pressures of the second zone. Anyway, it seems clear that the
evolution tendencies are quite similar, which can give complementary information about the
evolution of the film composition. First of all, and as a result of the fitting results, Fig. 7 shows
the evolution of the types of carbon bonds that are present in the coatings.
The first important remark from this figure is that there is some constancy in the amount of
C—O bonds and a significant reduction in both carbide (Ti—C) and graphite-like bonds (C or
C—C), as the oxygen partial pressure was increased. This result indicates that as the oxygen
increases in the samples, an expected decrease of Ti—C bonds is obtained, which is most
probably resulting from some C being replaced by O. Furthermore, and although always
present, the reduction of graphite-like bonds indicates that the isolated carbon in the samples
(which could be present as an amorphous tissue at the grain boundaries — a nanocompositelike behavior), is also reduced.
Fig. 8 shows the evolution of the ratio between oxygen and carbide bonds, revealing some
extra insights about this progressive change from Ti—C to Ti—O bonds, as previously
mentioned. This plot shows that the ratio O/carbide increases from approximately 3 (sample
from zone I) to about 11 (sample from zone III), revealing an important and progressive change
of the Ti—C species towards those of Ti—O. It is also worth to note that both samples from
zone II are relatively similar in terms of this species ratio, but with a value which is about twice
of that observed for the sample within zone I. Again, an increase to a value twice as high is
observed when going from then zone II samples to that of zone III.
2.2.2. Valence band
Fig. 9 shows experimental XPS valence band spectra for representative samples lying in the
different compositional zones previously identified. It clearly illustrates that the electronic
spectra of a TiCxOy film system adopt a relatively complicated structure. Four main bands can
be distinguished at least, where the highest binding energy bands are mostly composed of C2s
states, and then follow O2p and C2p bands, and the near-Fermi states are composed mainly by
Ti3d bands (relative to standard references like graphite for C2s and C2p bands, Ti for Ti3d and
TiO2 for O2p). In fact, Fig. 9 shows that there is a decrease of the Ti band at the Fermi level,
and a parallel growth of a broad band at energies between 3 and 9 eV, with the variation of
the O content in the films. Associated to the decrease of the d band at the Fermi level is the
decrease in the intensity of the C2s bands and a relatively high intensity of the O2s bands,
when going from films from zone I to zone III. This behavior is, as already mentioned,
consistent with a variation from C to O dominant films, which in its turn is associated with the
composition evolution, showing some change in the preferential bonding of Ti that is probably
evolving for Ti—C in zone I towards a most probable one with oxygen, Ti—O.
9
Also from Fig. 9, it is possible to have an idea about the energy gap between the top of the
valence band and the Fermi level, which can be claimed to be around 1.5 eV (energy distance
between the p band and Fermi level) for the films corresponding to the oxide-like zone (zone
III), reinforcing the idea of a progressive tendency to form insulating-type films. This could be
interpreted as coming from the fact that, contrarily to f.c.c. titanium (oxy)carbide (TiCxOy films
within zone I), which is basically a metallic-like material; titanium dioxide (and corresponding
films from zone III) is a semi-conductor, with a relatively large energy difference between the
2p and the 3d orbitals. This energy difference originates an intrinsic band gap of about 3.0 eV
for rutile TiO2 and 3.2 eV for anatase TiO2. Anyway, it is widely recognized that these values
can be tailored for electrical and photophysical-based applications by electronic state
associated with defects and dopants [43].
2.3. Structural features: X-ray diffraction and Raman spectroscopy
In order to follow the influence of the variations of the chemical composition in the structural
features of the films, a detailed structural characterization was carried out. The differences in
phase formation and orientation can be observed from X-ray diffraction patterns in Fig. 10 and
Fig. 11. The films prepared within the metallic mode, with values of composition ratios CO/CTi
between 0.14 and 1.44, and (CO + CC)/CTi between about 0.4 and 1.6, corresponding to partial
pressures of oxygen between 0 and 3.2 × 10− 2 Pa, exhibit patterns suggesting that the films
crystallize in a B1-NaCl crystal structure. One can note the presence of traces of h.c.p. Ti peaks
in the samples prepared without oxygen flow and with a partial pressure of oxygen of 7.8 ×
10− 3 Pa (not shown), together with two other diffraction peaks located at 2θ = 37.1° and
62.5°, that seem to correspond to a face centered cubic structure (f.c.c.). The f.c.c. structure
observed can be of TiC (ICDD card no. 32-1383) or TiO (ICDD card no. 77-2170). The fact that
these two samples are the most sub-stoichiometric ones (CO/CTi of 0.14 and 0.17 and (CO +
CC)/CTi of 0.42 and 0.49, respectively) could explain the tendency to have the mixture of both
h.c.p. Ti and cubic phases. This is in agreement with XPS results, revealing peaks attributed to
both elemental Ti and/or TiC and TiO in zone I. Anyway, there is a shift in the peak position of
the h.c.p. Ti (ICDD card no. 44-1294) towards lower angles (increase of the lattice), which may
be the result of its doping with some oxygen or carbon atoms.
Regarding the cubic phase, its exact nature is somewhat difficult to identify, although a careful
look at the composition plots, Fig. 1, may give some good indications. The first sample
(prepared without oxygen flow) has only about 20 at.% carbon and the assumption to claim a
TiC-type structure may be difficult due to the huge shift of the diffraction peaks from the TiC
lattice position (ICDD card no. 32-1383), since the main peak is located at about 2θ = 37.2°,
whereas the TiC (111) peak should be located at 2θ = 35.9°. On the other hand, the diffraction
peaks are rather close to the expected position for a f.c.c. TiO-type phase but, since this
sample has only ~ 10 at.% oxygen, the formation of phase does not seem possible. The reason
might be related to the thermochemistry of the reactions that take place. With oxygen present
in the reactor, its addition to growing structures is highly probable since the formation of a
titanium–oxygen bond is energetically more favorable than that of titanium with carbon, as
already seen.
10
Because the low oxygen content may not be enough to form TiO, the f.c.c. lattice is most likely
the result of a highly sub-stoichiometric TiC-type lattice. That is coherent with the significantly
lower lattice parameter in comparison to that of TiC (shift of diffraction peaks to higher angles,
Fig. 10), with oxygen occupying some of the C vacancies, forming a kind of Ti(C,O) solid
solution-type structure. This would mean a TiC-driven structure, with growing additions of
oxygen. The XPS measurements also induced the coexistence of Ti—C and Ti—O bonds in the
samples with small amount of oxygen. Moreover the ratio O/carbide is still quite low in this
zone and increases with the oxygen flow rate. The existence of h.c.p. Ti is also consistent with
the high sub-stoichiometric condition of this sample and the traces of Ti—Ti bonds detected by
XPS in zone I. A similar behavior can be claimed for all the samples prepared with oxygen
partial pressures up to 1.8 × 10− 2 Pa, where both CO/CTi and (CO + CC)/CTi ratios are below 1,
and thus favoring the coexistence of both C and O in the f.c.c. lattice.
It is also interesting to note the change in the lattice parameter evolution for the samples
prepared with atomic ratio (CO + CC)/CTi higher than 1.2. For these samples, the shift of
diffraction peaks corresponds to lattice parameters decreasing (diffraction peaks appearing at
higher diffraction angles, Fig. 10), approaching the value of unconstrained TiO. The main
features in this second set of samples are most likely to be those related to an excess of
oxygen in the lattice (interstitials) or even some C that may remain within the lattice.
Nevertheless, the energetically more favorable Ti—O bond tends to prevail with the increase
of the available O and thus the lattice tends to approach to that of “pure” TiO. Within this
second set of samples (zone II), the possibility to also have some carbon within grain
boundaries must be taken into account as observed in other works focused on C-based films.
Beyond these structural features, Fig. 10 shows also some tendency towards a progressive
change of the preferential growth from < 111 > to < 200 > with increasing metalloid
concentration. Texture change has often been the subject of extensive research, mainly in the
case of transition metal nitrides. The most widely accepted explanation is that the growth of
these structures such as those of transition metal nitrides (TiO and TiC in the case of this work)
by reactive sputter deposition is ruled by the interplay and competition between several
surface reaction and diffusion processes. In the particular case of films growing with reduced
surface mobility (absence of ion bombardment, such as the case in this work), results suggest
that the preferred orientation is determined by the overall lowest energy plane, which comes
from a critical competition between the residual strain and surface energy. Although (200)
planes are the most stable ones in these nitride/carbide/oxide cubic structures, resulting from
the minimum surface energy, the film growth by sputtering is highly dynamically controlled, so
that (111) planes with faster growth rate dominate. The orientation change from < 111 > to <
200 > between zone I and zone II may result from the stoichiometry changes that occur. The
atomic ratios CO/CTi and/or (CO + CC)/CTi tend to 1, for the films where < 111 > growth
occurs, while those where these values are significantly higher than 1 (films prepared with
oxygen partial pressures between 1.9 × 10− 2 and 3.5 × 10− 2 Pa, corresponding to (CO +
CC)/CTi atomic ratios above 1.35) showed a < 200 > preferential growth, as it can be observed
from Fig. 10. For the lowest oxygen partial pressures, the growth rate of (200) grains is
inhibited due to the lattice distortion (lattice defects, vacancies, etc.).
11
Regarding the samples deposited with the highest oxygen partial pressures (oxidized
sputtering mode), where the composition analysis showed that the films have a stoichiometry
close to that of TiO2, and thus, as expected, the X-ray diffractograms confirm the oxide phase
formation, revealing the existence of body-centered-titanium oxide phase (anatase-type), aTiO2 (ICDD card no. 21-1272), Fig. 11. There is no clear evidence of rutile-type TiO2
(tetragonal). The formation of TiO2-type compounds is somewhat expected due to the
stoichiometry of the samples.
Further details on the structure of the films were obtained by micro-Raman spectroscopy.
Being a complementary technique to the previous ones, Raman was expected to give a set of
further information, which may allow understanding the overall set of properties of this thin
film system [44], [45] and [46]. Fig. 12 displays the variation of Raman spectra as a function of
the rOC ratio, for samples from each zone, for a spectral region between 50 and 700 cm− 1. It
is worth remembering that compounds with the rock-salt structure like TiC and TiO have no
first order Raman-active vibrational modes, therefore the bands in the spectra are only
induced by disorder. The Raman bands of these binary titanium-based compounds should
appear in the low frequency range (up to 700 cm− 1). For the sample with rOC = 0.38 (without
oxygen flow) and for the sample rOC = 1.02, both belonging to zone I, the spectrum does not
present any significant signal (Fig. 12a) and b) of any structure like TiC, TiO or a mixture of
both, and no evidence of any isolated carbon is detected.
For the samples from zone II, where rOC = 1.35 and rOC = 1.62, (Fig. 12c) and d) respectively),
a large band centered at ~ 270 cm− 1, begins to emerge indicating the formation of a quite
disordered structure. This signal could be attributed to a complex TiCO structure. Indeed, the
TiC Raman bands should appear at 260 cm− 1, 420 cm− 1 and 605 cm− 1, while, in Ti2O3
compounds, the Ag mode corresponding to the Ti—O vibration appears at 269 cm− 1 (Fig.
12b). Raman spectroscopy is consistent with XRD results, as it proves the existence of a
complex structure of Ti(C,O). The C/Ti ratio remains constant between zone I and zone II, so
the appearance of this structure can only be explained by the increasing participation of
oxygen in the complex structure of Ti(C,O). Thus we can conclude that the change observed in
XRD (grain growth, change in lattice parameter, …) should be attributed to the addition of
oxygen atoms in the f.c.c. structure of Ti(C,O), going from a TiC dominant to a TiO dominant
structure, still in a complex structure of the type Ti(C,O).
In the transition between a Ti(C,O)-type structure to a TiO2-type structure, the samples
become transparent, and for the sample with the highest oxygen content, the peaks located at
144 cm− 1, 197 cm− 1, 396 cm− 1, 519 cm− 1 and 639 cm− 1, corresponding to the TiO2
anatase structure are clearly seen (Fig. 12e). This is in complete agreement with the XRD
analysis: for both techniques, TiO2-anatase structure is detected for the highest oxygen partial
pressures, 5.5 × 10− 2 Pa and 5.9 × 10− 2 Pa.
2.4. Electrical properties
Titanium oxycarbide thin films prepared with oxygen flow rate higher than qO2 = 5.0 sccm
(zone III) are too resistive to be measured by the van der Pauw method. For film sputter
deposited up to qO2 = 4.5 sccm (zones I and II), the electrical conductivity σ as a function of
12
temperature T was measured for various oxygen flow rates, which corresponds to a rising
oxygen atomic concentration (Fig. 13).
A typical exponential behavior σ = A exp (− E/kT), where A is a constant and E is the activation
energy, is observed versus the temperature and for oxygen content higher than 38 at.%.
Without oxygen injection, electrical conductivity is higher than 3.40 × 105 S m− 1 and exhibits
a metallic-like characteristic behavior since conductivity decreases with temperature. The
corresponding temperature coefficient of resistance is TCR300K = 1.40 × 10− 6 K− 1, which is a
low value for a metallic compound (order of 10− 3 K− 1 for bulk metals). This is mainly due to
the quite high oxygen content (10 at.%) in the film. In addition, such a metallic-like behavior
can also be correlated with the presence of h.c.p Ti as suggested from XRD pattern (Fig. 10)
and the traces of Ti—Ti bonds detected from XPS. Conductivity is significantly reduced down to
σ300K = 5.23 × 104 S m− 1 as the oxygen flow rate reaches qO2 = 2.0 sccm ([O] = 45 at.%).
Similarly, a semi-conducting behavior is observed. This range of high electrical conductivities
well corresponds to the zone I previously defined from compositional and structural analyses.
This slow decrease of conductivity versus oxygen concentration cannot be due to changes of
the crystallographic structure since XRD analyses showed that Ti(C,O) coatings produced with
oxygen flow rate lower than qO2 = 4.5 sccm (PO2 = 3.5 × 10− 2 Pa and 67 at.% of oxygen)
always adopt an f.c.c. structure. In addition, the rise of the grain size as a function of rOC
should produce an enhancement of the transport properties due to larger crystallite grain sizes
and thus a reduction of the number of grain boundaries per electron mean free path.
However, a reverse evolution is measured from results presented in Fig. 13 and Fig. 14.
Therefore, this slight reduction of conductivity in zone I can be attributed to the coexistence of
Ti—O and Ti—C bonds with some amounts of oxygen, as proven from XPS results. The
occurrence of Ti(C,O) solid solution, which is consistent with the values of the lattice
parameters, explains that the order of magnitude of the conductivity is in the metallic range in
zone I. Since rOC is higher than 1, substitution of C by O atoms is not the only phenomenon
occurring as the oxygen supply rises. The amounts of remaining carbon and oxygen in the film
are then significant to favor the resistive character of the material.
For oxygen concentrations higher than 60 at.%, DC electrical conductivity abruptly drops to
σ300K = 2.68 × 102 S m− 1 (Fig. 13 and Fig. 14), and similarly, the semi-conducting behavior
prevails. Activation energy is strongly increased and tends to 80 meV as the oxygen
concentration reaches 67 at.%. This corroborates XPS results well. From Ti2p signals, the
transition from zone I to zone II corresponds to the evolution of monoxide TiO-type bonds to
TiO2-type ones. In addition, films in zone II tend to be less crystallized. As a result, the
transport properties in this second zone (Ti(C,O) tend to behave as semi-conductors) cannot be
assigned to some phase transformation, it is rather linked to major changes of the chemical
bonding states. The coexistence of Ti—C, Ti—O and TiO2-type bonds may occur in zone II, but
TiO2-type predominates as oxygen content rises.
A further increase of oxygen supply leads to TiO2-like compound. Electrical conductivity
cannot be measured by the van der Pauw method. The films are insulator as suggested by the
sudden drop of conductivity (end of zone II for oxygen contents higher than 67 at.%.) and the
activation energy should be higher than few hundreds meV.
13
2.5. Optical properties
Typical reflectance spectra in the UV, visible and near IR wavelength range (250–2500 nm) of
Ti(C,O) thin film sputter deposited on glass also exhibit significant changes according to the
oxygen concentration (Fig. 15).
Reflectance spectra of titanium oxycarbide thin films measured at 10° of incidence in the
wavelength range 250–2500 nm. The influence of the oxygen flow rate is shown and the
changes of the optical reflection also correlate with the different zones I, II and III.
Titanium oxycarbide thin films produced without oxygen supply exhibit the highest reflectance
for the full range of investigated wavelengths, especially in the near IR region where
reflectance R10° is over 70% at 2 μm (qO2 = 0 sccm in Fig. 15). This is typical of a conductinglike behavior and in agreement with the metallic characteristic observed from electrical
conductivity versus temperature (Fig. 13). For films prepared in zones I and II (e.g. qO2 = 2.0
and 4.0 sccm, respectively), a high reflectance is still measured in the near IR wavelengths.
Optical properties of the films in zones I and II show no transmission (not shown here) and
reflectance spectra are consistent with those commonly recorded for metallic materials.
However, a minimum of reflectance is observed close to 400 nm and is shifted from UV to
visible region as the oxygen concentration is increased. At the highest wavelengths, the
spectra exhibit the Hagen–Rubens behavior (1-R proportional to E1/2) [47] and [48], which is
characteristic of metallic conductors and attributable to free carriers (Dingle plateau well
described by a Drude model). At lower wavelengths, the drop of reflectance is observed with a
minimum value of the reflectance for all films located in zones I and II. It is worth noting that
there is a shift of this minimum to higher wavelengths as a function of the oxygen
concentration. This minimum correlates with an energy threshold for interband transitions.
Such kinds of transitions are noticeable in both zones where a broad peak appears close to 250
nm. Shifts of the energy threshold as well as peak due to interband transitions have to be
associated with carbon and oxygen vacancies created in the f.c.c. Ti(C,O) lattice as the oxygen
concentration increases in zones I and II. In addition, from compositional analyses, it was
clearly shown that rOC increases with oxygen concentration. As a result, an increasing of the C
and O vacancy concentrations takes place, which reduces the charge carriers density [49] and
consequently, the optical reflectance and the DC electrical conductivity (as previously
observed in Fig. 13 and Fig. 14).
For Ti(C,O) films corresponding to zone III (qO2 higher than 5.0 sccm), reflectance spectra
show classical interference fringes. It confirms the dielectric character of Ti(C,O) films with
oxygen concentrations higher than 67 at.% (rOC around 2). It is also in agreement with XRD
investigations (amorphous or anatase phase), XPS analyses (chemical states close to those of
TiO2-like compound) and DC electrical conductivity (activation energy higher than few
hundred meV). A further increase of the oxygen flow rate from qO2 = 5.0 to 7.0 sccm, leads to
a slight reduction of fringe amplitude (not shown here). The slight reduction may be attributed
to a lesser amount of defects in the TiO2 lattice or carbon remaining in the films structure at
the grain boundaries or in the anatase phase.
3. Summary and conclusions
14
It is clearly shown that the oxygen flow rate and thus the composition of the atmosphere in
the deposition chamber control the composition of the titanium oxycarbide thin films obtained
by reactive sputtering. The increase of the oxygen flow rate was correlated to the composition
of the films obtained by Rutherford Backscattering Spectroscopy. Three zones appeared to
characterize the evolution of the composition. Moreover, XPS was carried out to study the
chemical changes in the films. XPS results showed the changes in the nature of the chemical
bonds developed by carbon, oxygen and titanium. Through the three zones, there is an
important and progressive change of the Ti—C bonds towards those of Ti—O.
Structural characterization results revealed an undefined f.c.c. structure attributed to a Ti(C,O)
solid solution-type structure for low oxygen contents and a TiO2-anatase structure for high
oxygen contents. Micro-Raman experiments confirmed the existence of a Ti(C,O) disordered
structure. Electrical conductivity measurements of the films revealed a gradual transition from
metallic to semi-conducting and to insulating behaviors as a function of the increase of the
oxygen flow rate. Regarding reflectance variations, it can be observed there is a gradual
change of the optical measurements, as a function of the oxygen flow rate, which is clearly
related to the semi-conducting behavior of the films.
Acknowledgments
This research is sponsored by FEDER funds through the program COMPETE-Programa
Operacional Factores de Competitividade and by national funds through FCT-Fundação para a
Ciência e a Tecnologia, under the project PTDC/CTM-NAN/112574/2009.
15
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19
Figure captions
Figure 1. Variation of the: a) atomic concentration of the TiCxOy thin films and; b) atomic
ratios, rO = CO/CTi, rC = CC/CTi and rOC = (CO + CC)/CTi; as a function of the oxygen partial
pressure. For comparison purposes, the correspondent gas flows were also added at the top
axis (tick labels in this axis do not vary linearly: each space corresponds to 0.5 sccm).
Figure 2. XPS spectra of Ti2p3/2 and Ti2p1/2 lines recorded on a TiCxOy sample prepared with
an oxygen gas flow of 6 sccm (PO2 = 3.7 × 10− 2 Pa), for different sputter cleaning times.
Figure 3. XPS spectra of Ti2p3/2 and Ti2p1/2 lines recorded on selected samples prepared
within each of the three zones, after the steady state has been reached by sputter cleaning. A
TiO2 reference sample is also included.
Figure 4. XPS spectra of O1s line recorded on selected samples prepared within each of the
three zones and after the steady state has been reached by sputter cleaning.
Figure 5. XPS spectra of C1s line recorded on selected samples prepared within each of the
three zones and after the steady state has been reached by sputter cleaning.
Figure 6. XPS calculated elemental concentrations (open symbols) as a function of the oxygen
partial pressure. The results were obtained after the steady state has been reached by sputtercleaning. For comparison purposes, the values obtained by RBS (plotted in Fig. 1a) are also
included (full symbols).
Figure 7. Evolution of the types of carbon bonds in the analyzed coatings.
Figure 8. Evolution of the concentration ratio O/carbides in the analyzed coatings.
Figure 9. XPS valence band spectra of TiCxOy representative samples (deposited on Si
substrates) within the different compositional zones.
Figure 10. X-ray diffraction patterns of titanium oxycarbides deposited on silicon, with
different rOC ratios lower than 1.62. The patterns were vertically shifted for the sake of clarity.
Figure 11. X-ray diffraction patterns of titanium oxycarbides deposited on silicon, with
different rOC ratios higher than 2.18. The patterns were vertically shifted for the sake of
clarity.
Figure 12. Raman spectra of titanium oxycarbides for a) rOC = 0.38, b) rOC = 1.02, c) rOC =
1.35, d) rOC = 1.62 and e) rOC = 2.32. f) rOC = 2.36 (Raman intensity divided by 1.2.
Figure 13. Electrical conductivity σ versus reverse of temperature 1000/T measured on
titanium oxycarbide thin films for various oxygen concentrations in the films. The gradual
change of conductivity correlates with zones I and II previously defined as a function of the
oxygen mass flow rate.
Figure 14. DC electrical conductivity σ300K and activation energy Ea as a function of the
oxygen concentration in the films. Transition from zone I to zone II is brought to the fore as the
oxygen concentration reaches 50 at. %.
20
Figure 15. Reflectance spectra of titanium oxycarbide thin films measured at 10° of incidence
in the wavelength range 250–2500 nm. The influence of the oxygen flow rate is shown and the
changes of the optical reflection also correlate with the different zones I, II and III.
21
Figure 1
22
Figure 2
23
Figure 3
24
Figure 4
25
Figure 5
26
Figure 6
27
Figure 7
28
Figure 8
29
Figure 9
30
Figure 10
31
Figure 11
32
Figure 12
33
Figure 13
34
Figure 14
35
Figure 15
36