Nb-B-C thin films for electrical contact applications deposited by magnetron sputtering

The high wear resistance, high chemical inertness, and high electrical conductivity of magnetron-sputtered transition metal diborides make them a candidate material for sliding electrical contacts. However, their high hardness makes it difficult to penetrate surface oxides, resulting in a high electrical contact resistance. In this study, the authors have investigated how the contact resistance can be improved by the formation of softer Nb-B-C films. The Nb-B-C films were deposited by magnetron sputtering and shown to exhibit a nanocomposite microstructure consisting of nanocrystalline NbB2−x grains with a solid solution of C separated by an amorphous BCx phase. The formation of the BCx phase reduces the hardness from 41 GPa for the NbB2−x film to 19 GPa at 36 at. % C. As a consequence the contact resistance is drastically reduced and the lowest contact resistance of 35 mΩ (contact force 5 N) is achieved for a film containing 30 at. % C. However, crack formation and subsequent delamination and fragmentation is observed for the C-containing Nb-B-C films in tribology tests resulting in high friction values for these films.


I. INTRODUCTION
Magnetron-sputtered transition metal diborides (MeB 2 ) exhibits many interesting properties such as high hardness, high wear resistance, and high electrical conductivity.A potential application of MeB 2 films is as contact material in sliding electrical contacts, where low wear rates combined with high conductivity and chemical stability are required.The most widely studied diboride is TiB 2 , [1][2][3][4][5] which has excellent mechanical properties.A disadvantage with TiB 2 , however, is the high friction coefficient (>0.5) generally reported, 1,6 which makes it unsuitable in a sliding contact.An alternative diboride is sputtered NbB 2Àx films, 7 which we have recently demonstrated to exhibit a very low friction coefficient (0.16) against stainless steel.Unfortunately, the contact resistance of NbB 2Àx is far too high.This can be due to the very high hardness (>40 GPa) of the films, which makes it difficult to penetrate surface oxides and limits the deformation of the film in a contact situation.Studies on transition metal carbides 8,9 have demonstrated that single-phase carbide films (e.g., TiC and NbC) with a high hardness exhibit very high contact resistances, which could be reduced dramatically by the formation of a softer a-C matrix phase.This phase reduces the film hardness and increases the ductility, making it possible to break and penetrate thin surface oxides and thereby drastically reduce the contact resistance.A softer film will also deform more easily in a contact situation and thus form a larger contact area, resulting in a reduced contact resistance.For a contact application, the a-C phase must be only a few monolayers thick to maintain a low electrical resistivity in the film. 8,9Thus, a possible route to improve the electrical contact properties of the NbB 2Àx films could therefore be to alloy the films with carbon.The solubility of carbon in NbB 2Àx is low, 10 and it is likely that added carbon will segregate to the grain boundaries forming an amorphous boron-carbon (a-BC x ) containing matrix phase.The aim of this study is to investigate the effect of carbon alloying on mechanical and electrical properties of magnetron sputtered NbB 2Àx films.Nanocomposite films with nc-NbB 2Àx grains in an a-BC x matrix are deposited by magnetron sputtering from NbB 2 and C targets.The microstructure of the nanocomposites is investigated, and their mechanical, tribological, and electrical properties are evaluated with a special emphasis on potential use for sliding electrical contact applications.

II. EXPERIMENT
The Nb-B-C films were deposited in an ultrahigh vacuum chamber (base pressure of 10 À7 Pa) by DC magnetron sputtering from 50 mm circular NbB 2 (99.5%) and C (99.999%) targets in an Ar atmosphere at a constant pressure of 0.4 Pa (3.0 mTorr).The current to the NbB 2 -magnetron was kept constant at 150 mA, while the current to the C-magnetron were varied from 0 to 200 mA in order to achieve films with different C content.The samples were coated on to the following substrates; single-crystal Si(001) (10 Â 10 mm 2 ) and a) Electronic mail: nils.nedfors@kemi.uu.se b) Electronic mail: ulf.jansson@kemi.uu.seAl 2 O 3 (10 Â 10 mm 2 ) for structure analysis and electrical resistivity measurements, Ni-plated bronze (15 Â 15 mm 2 ) for electrical contact resistance measurements, and mirror polished 316 L stainless steel (20 Â 20 mm 2 ) for tribological analysis.A bias of À50 V and a constant temperature of 300 C were used in all experiments.A thin Nb/NbC adhesion film (total thickness $50 nm) were deposited on to the substrates prior to the primary deposition in order to improve the film adhesion.X-ray photoelectron spectroscopy (XPS) depth profiles acquired using a Physical Systems Quantum 2000 spectrometer with monochromatic Al Ka radiation and 2 keV Ar þ -ion sputtering over an area of 1 Â 1 mm 2 were used to determine the elemental composition.Sensitivity factors were determined from reference Nb-B-C samples with compositions acquired by elastic recoil detection analysis (ERDA).The chemical bond state of the films were determined from XPS spectra acquired after 30 min of 200 eV Ar þ -ion sputter etching over an area of 1 Â 1 mm 2 .Grazing incidence x-ray diffraction (GI-XRD) measurements were carried out using a Philips X'pert MRD diffractometer with Cu Ka radiation and parallel beam geometry with a 2 incidence angle.Microscopy studies were carried out on selected films, using a FEI Tecnai G2 TF 20 UT field emission gun transmission electron microscope (TEM) operated at a 200 kV acceleration voltage.The cross sectional TEM specimens were first mechanically polished to a thickness of $50 lm, followed by Ar þ -ion milling, with ion energy of 5 keV.As a final step, the samples were polished using 2 keV Ar þ -ions.Film thicknesses were determined by SEM on fractured cross-sections of the films.
A CSM Instruments nanoindenter XP with a diamond Berkovich tip was used for the assessment of hardness and elastic modulus by applying the Oliver-Pharr method 11 on the load-displacement curves acquired with an indentation depth set to 50 nm and a loading/unloading rate of 1.5 mN/min.The film adhesion was estimated using a CSEM Scratch Tester equipped with a 200 lm radius Rockwell diamond tip loaded from 0 to 70 N at a loading rate of 100 N/min, resulting in a 14 mm scratch path.The critical load of failure is taken at the contact load where an abrupt increase is seen for the acoustic emission.Tribological measurements were performed in ambient atmosphere with 55% relative humidity using a ball-on-disk set-up against stainless steel balls (100Cr6) with a radius of 6 mm, intended for ballbearings.The track radius was 2.5 mm, the sliding speed 0.1 m/s, and the contact force 1 N.A Zeiss Leo 1550 scanning electron microscope (SEM) equipped with an AZtec energy dispersive x-ray spectrometer (EDS) as well as an Olympus AX70 optical microscope was used to investigate the wear tracks.A Veeco WYKO NT1100 optical profiler was used to measure the surface curvature of the films in order to calculate the total residual stress using Stoney's equation corrected for films deposited onto Si(001) substrates. 12An Advanced Instrument Technology CMT-SR2000N four-point-probe was used to determine the film resistivity.The electrical contact resistance, measured between the film surface and an Au-coated probe in a custom built set-up, was acquired at three different contact forces (5, 7, and 10 N) and nine different spots on the film surface.The contact resistance at each contact force was taken as the average value from the nine spots after the lowest and highest values had been removed.

III. RESULTS AND DISCUSSION
The chemical composition of the as-deposited films is summarized in Table I.All films have a B/Nb ratio of 1.7-1.8,although they are sputtered from a NbB 2 target (see our previous study on NbB 2Àx films 7 ).It should be noted that NbB 2 exhibits an unusually wide homogeneity range.The composition of about NbB $1.8 for the binary film in Table I, is thus close to the lower limit of NbB $1.86 determined by Nunes et al. 13 Figure 1 shows GI-XRD diffractograms of the different films with peak positions from a reference NbB 2 bulk sample 14 included in Fig. 1.All the diffraction peaks can be assigned to NbB 2 with the hexagonal AlB 2 -type structure.With increasing C content, the peaks become broader and less intense, but clear indications of a NbB 2 phase can be seen also in the most C-rich film.The small peak at $34.5 seen for the films containing !12 at.% C and the small peak at $38.5 seen for the film with the highest C-content are assigned to the NbC and Nb adhesion layers, respectively.Furthermore, the position of some diffraction peaks shift with composition, which can be attributed to either stresses or a solid solution of carbon into the boride structure.Table II shows the calculated cell parameters and residual stresses for all films.As can be seen, the a-axis increases slightly from about 3.12-3.14Å in the most C-rich film.In contrast, the c-axis shows a stronger dependence on composition with a possible maximum at lower carbon contents followed by a slightly reduced cell-axis in the most carbon-rich film.The cell volume increases from 29.07 Å3 in the C-free film to 29.70 Å3 at 12 at.% C, see Table II.Mayrhofer et al. 2 observe for a TiB 2.4 film a shift of 0.17 for the 001 XRD peak when the residual stress of the film is reduced from À2.72 GPa to a nearly unstrained state.The residual stresses in our films differ by less than 1 GPa between the films in this study, suggesting that the peak shifts mainly can be attributed to a solid solution of carbon into the NbB 2Àx phase.At equilibrium, the solubility of C into NbB 2 is very low (<3 at.%). 10 However, it is well known that magnetron sputtering can produce thin films with highly supersaturated concentrations of dissolved elements.One example is Ti-Fe-C, 15 where only about 1 at.% Fe can be dissolved into the TiC phase at equilibrium, while >20 at.% Fe easily can be dissolved on the Ti sites in magnetron sputtered films.The observation of such a solid solution of C into the boride structure is in agreement with observations in the Ti-B-C system, where the solid solution leads to a TiB x C y phase (see, e.g., Refs.16 and 17).The exact concentration and the position of the carbon atoms in the NbB 2Àx structure are unknown, but most likely they are placed at vacancies in the boron sublattice.Finally, the peak broadening in Fig. 1 indicates a reduction in size of the NbB 2 grains with the increase in C content, in a similar way to what is seen for Ti-B-C films. 17Applying Scherrer's formula we can estimate that the NbB 2 grain size is reduced from about 10 to 2 nm as the C content increases in the films, see Table II.
Figure 2 shows cross-sectional TEM images of a Nb-B film and Nb-B-C films with 21 and 36 at.% C in a, b, and c, respectively.The pure NbB 2Àx film has thin (5-10 nm) columnar grains elongated in the growth direction.As 21 at.% C is added to the films the NbB 2 diffraction rings in the selected area electron diffraction (SAED) pattern become broader indicating a reduction in the size of the NbB 2 grains, see Fig. 2(b).The dark field image, obtained using segments of the 001 and 100 diffraction rings, shows equiaxed crystalline 3-5 nm large grains [as measured in high resolution TEM (HRTEM)] surrounded by an amorphous structure.The Z-contrast image taken with the high angle annular dark field (HAADF) detector in scanning TEM mode show brighter regions surrounded by darker regions, which can be attributed to Nb-rich regions (bright areas) surrounded by a Nb-deficient phase.For the film with 36 at.% C [see Fig. 2(c)], an amorphouslike structure is seen in the low magnification dark field image.SAED pattern for this film, however, shows broad diffraction rings, indicating that there exist small crystallites in the film.The diffraction ring positions coincide with the pattern seen for the film containing 21 at.% C and thus confirms the existence of a NbB 2 phase also in this film.The high-resolution image confirms the mainly amorphous structure with only some occasional grains with size less than 3 nm.The HAADF image in Fig. 2(c) shows a two-phase structure, similar to what is seen at 21 at.% C, with brighter Nb-rich regions with a diameter of $2 nm surrounded by darker Nb-deficient regions.
Figure 3 shows the B1s and C1s XPS spectra from all samples.The B1s spectra are clearly composed of several peaks.In a previous study 7 we demonstrated that the B1s spectrum from the binary NbB 2Àx film can be separated into four peaks originating from B-Nb in NbB 2Àx (188.8,187.2, and 188.0 eV) and from B-B bonds at 187.5 eV in the tissue phase between the NbB 2Àx grains.The three B-Nb peaks originate from boron in the bulk of the grains (B-Nb b ), surface boron (B-Nb s ), and defects (B-Nb d ), and their origin have been described in detail in a study by Aizawa et al. 18 on single-crystal NbB 2 .For the films with !12 at.% C, an additional peak at 189.1 eV appears and increases in intensity with increasing C content.This peak therefore presumably originates from B-C bonds and has previously been reported 19 at this binding energy for an amorphous BC x phase in Ti-B-C films.It can thus be concluded that C segregates to the grain boundaries forming an a-BC x phase.The XPS of B-C films are however disputed, and a survey by Jacobsohn et al. 20 concludes that both B1s and C1s spectra are composed of several chemical environments.The C1s spectra of the different films are shown in Fig. 3(b).Good fit of the spectra is achieved using three peaks located at: 282.8, 283.7, and 284.5 eV.The peak at 284.5 eV can be assigned to C-C bonds in a-C and has previously been observed in magnetron sputtered Ti-B-C films. 17,21,22The other peaks denoted I and II are more difficult to identify.The XRD results suggest a solid solution of carbon into the NbB 2 structure.Niobium carbide films have a C1s binding energy of 282.8 eV, 9 the peak denoted II at 282.8 eV can therefore possibly be attributed to C atoms forming a solid solution in the NbB 2 phase.However, studies on Ti-B-C films 17,22   suggest that C dissolved in TiB 2 may exhibit a chemical shift toward higher binding energies.This shift is about 1 eV and the similar effect should give rise to a feature at about 283.8 eV, more or less at the position for peak I in Fig. 3(b).Furthermore, as stated previously, the C1s spectra in B-C films are composed of several different chemical environments and C1s peaks at both $283 eV and $284 eV are reported. 20The main peak at 282.II.All films have a compressive residual stress of 0.9-1.6GPa.Studies by Mayrhofer et al. 2 and Berger et al. 4 have shown that no significant correlation between residual stress and hardness exists for TiB 2 films.The local maxima for the hardness at 12 at.% C seen in Fig. 4 can be attributed to the reduced boride grain size from 10 nm for the Nb-B film to 6 nm for the C-containing film, as estimated by Scherrer's formula, rather than the increase in residual stress.Upon addition of C to the NbB 2Àx film, some C will segregate to the interfaces modifying the B-rich tissue phase.This will lead to a grain size reduction further improving the hardness.At the higher C contents, where the boride grains are 2-3 nm in size, the volume fraction of the a-BC x phase will increase and plastic deformation will operate by grain boundary sliding rather than dislocation slip. 23,24As a consequence, the hardness should decrease, as seen in Fig. 4. A similar trend has been observed for nc-NbC/a-C films 25 of different C contents with NbC x grain sizes in the range of 3-5 nm and an a-C matrix thickness that is not varying with C content.The drop in elastic modulus from 580 GPa for the C-free film to 340 GPa for the film containing 12 at.% C can be connected to the formation of the a-BC x matrix, exhibiting a lower elastic modulus (250-300 GPa has been reported for a-B 4 C films 26 ) compared to the NbB 2Àx phase.As more C is added to the films, the volume fraction of the a-BC x phase increases, resulting in a further reduction of the elastic modulus.Scratch tests were performed in order to evaluate film adhesion and critical load of film failure is given in Table II.Frequent crack formation starting already at a load of 6-8 N was observed in the scratch paths, which agrees with the rather small residual stresses observed for all films.No indication of film delamination was seen for any of the films, and thus, it is not possible to connect changes in the appearance of the scratch paths to the sudden increase in the acoustic emission seen at the critical load of film failure.These findings indicate a good film adhesion to the steel substrate for all C contents.The friction properties were evaluated by ball-on-disk measurements against stainless steel balls at relative humidity of 55%.A low and steady coefficient of friction of 0.16 is seen for the reference Nb-B film.Also, the film containing 21 at.% C showed in a first measurement a stable coefficient of friction of 0.16.However, in a second run, a coefficient of friction of about 0.15 is seen for the first $200 laps, followed by a drastic increase to 1.0.For the other films, the coefficient of friction is about 0.3 in the beginning of the test and then almost directly (after 50-100 laps) raise to 0.8-1.0. Figure 5 shows optical photographs of the wear track from the Nb-B film, which show a low and stable coefficient of friction [Fig.5(a)] and the film with 30 at. % C, where the friction drastically increased to 1.0 [Fig.5(b)].The wear track connected to a low friction is 135 lm wide and show an even distribution of dark dots 1-2 lm in size, which probably are surface contaminants.An increased amount of oxygen in the wear track was found by SEM EDS-mapping (not shown).The wear track in Fig. 5(b) is representative for the other films, where drastic increases in coefficient of friction after 50-100 laps were observed.The wear track is much wider (240 lm) and contains dark areas hundreds of micrometers in size.The image taken at high magnification shows the formation of cracks in the film.Especially the cracks seen in connection to the dark areas indicate that the film has burst open and detached in these areas.Wear particles can also be seen along the edges of the wear track.Furthermore, a strong Fe signal, but no Nb and B signals, was detected in these dark areas by SEM EDS-mapping (not shown).The formation of wear debris in the wear track and metal to metal contact between the counter surface and the exposed steel substrate will cause the drastic increase in the coefficient of friction for the C-containing films.
As seen in Table II, the resistivity of the films increases linearly with C content, from 100 lX cm for the reference Nb-B film to 412 lX cm for the film with 36 at.% C. The reduction in boride grain size with C content results in a higher fraction of grain boundaries (i.e., electron scattering centers) as well as a higher volume fraction of a-BC x phase, and consequently an increase in resistivity.The electrical contact resistance measured at different contact forces are plotted in Fig. 6.The two films containing 12 at.% C have a wide spread in contact resistance values in the range of 650-15 000 mX and are therefore far outside the range of values that are of most interest in Fig. 6.The contact resistance of the film containing 21 at.% C decreases drastically with the contact load from 5200 mX at 5 N to 174 mX at a contact force of 10 N. The two films containing 30 and 36 at.%C have much lower contact resistance values that are less sensitive to the contact force and varies in the range of 35-62 mX and 55-79 mX, respectively.Contact resistance of a Nb-C sample 25 has been included in Fig. 6 as a reference as well as the contact resistance for gold against gold (0.4 mX).An interesting observation is also that the trend for the electrical contact resistance is opposite to the resistivity with the lowest contact resistance of 35 mX obtained for the C-rich film containing 30 at. % C.This can be explained by the fact that most C-rich films are softer and more ductile than the films with lower carbon content due to the formation of the a-BC x phase.This makes it easier to penetrate surface oxides upon a mechanical load enabling an electrical current to pass through the contact junction.Similar effects have been well-established for nc-MeC/a-C nanocomposite films. 8,9A softening of the film can also facilitate deformation in a mechanical contact resulting in larger contact area and thus a lower contact resistance.

IV. CONCLUSIONS
We have shown that C in Nb-B-C films segregates to the NbB 2Àx interfaces forming an amorphous a-BC x phase during deposition by magnetron sputtering at 300 C. XPS and XRD results suggest that some carbon is also dissolved into the boride structure.The presence of an a-BC x phase drastically reduces the hardness and increases the ductility of the film.As a consequence, the electrical contact resistance is strongly reduced.However, the poor mechanical properties result in crack formation during the friction test followed by delamination and fragmentation of the C-containing films and thus a high coefficient of friction of 0.8-1.0.The crack formation needs to be suppressed before Nb-B-C films can be a candidate material for sliding electrical contact applications.

FIG. 1 .
FIG. 1. GI-XRD diffractograms of the Nb-B-C films with different C contents acquired with a 2 incidence angle.Vertical lines indicate diffraction peak positions for a reference bulk NbB 2 sample (Ref.13).
8 eV (II) as well as the intermediate peak at 283.7 eV (I) can therefore also have contributions from C in a-BC x .Consequently, without more detailed experimental and theoretical studies, it is at present impossible to determine the contributions from C-Nb and C-B bonds, in peak I or II, respectively.However, if both peaks I and II are considered to only have contribution from C in a-BC x , i.e., no solid solution of C in the NbB 2Àx phase, a much higher intensity would be expected for the B-C peak at 189.1 eV in the B1s spectra [Fig.3(a)].Thus, it can be concluded that parts of the C most likely form a solid solution with the NbB 2Àx phase.The mechanical properties of the films obtained from nanoindentation are plotted in Fig.4as a function of C content.The Nb-B film has a hardness of 41 GPa and a maximum in hardness of 46 GPa is seen at a C content of 12 at.%.As the C content is increased further the hardness is reduced and the lowest hardness of 19 GPa is achieved for the film containing 36 at.% C. No such maximum is seen for the elastic modulus, which decreases from 580 GPa for the Nb-B film to 290 GPa for the Nb-B-C film containing 36 at.% C. The residual stresses of the different films estimated from the curvature of the Si substrates are presented in Table

FIG. 2 .
FIG. 2. Cross-sectional dark field TEM images of (a) Nb-B film, (b) Nb-B-C film containing 21 at.% C, and (c) Nb-B-C film with 36 at.% C. The dark field images are obtained using segments of the 001 and 100 diffraction rings.Insets show SAED (top), HRTEM (middle) and Z-contrast HAADF (bottom).A boride grain is marked in each HRTEM image.

FIG. 5 .
FIG. 5. Optical microscope images, taken at two different magnifications, of the wear tracks after ball-on-disk test against stainless steel for (a) the Nb-B film and (b) the Nb-B-C film containing 30 at. % C.

TABLE I .
Current applied to the C magnetron during deposition and the chemical composition of the different films acquired from XPS depth profiles.The given compositions are normalized to the sum of the Nb, B, and C contents.

TABLE II .
Structural parameters, resistivity, and critical load to film failure for the different films.Grain sizes calculated using Scherrer's formula on the 001 and 101 diffraction peaks.