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Surface & Coatings Technology 412 (2021) 127059

Available online 12 March 2021

0257-8972/© 2021 The Author(s). Published by Elsevier B.V. This is an open access article under the CC BY license (http://creativecommons.org/licenses/by/4.0/).

Phase formation and structural evolution of multicomponent

(CrFeCo)

1-y

N

y

films

Smita G. Rao

*

, Rui Shu , Robert Boyd , Grzegorz Greczynski , Arnaud le Febvrier , Per Eklund

Department of Physics, Chemistry, and Biology (IFM), Link¨oping University, Link¨oping 58183, Sweden

A R T I C L E I N F O Keywords:

Multicomponent nitride Thin film

Magnetron sputtering High entropy alloy

A B S T R A C T

The Cantor alloy (CoCrFeMnNi) and its variants, in bulk as well as thin films, have been extensively studied. They are known to exhibit cubic crystal structures and thermodynamic stability regardless of their complex chemical composition. Therefore, they may find use as hard, wear-resistant, corrosion and oxidation-resistant coatings. The addition of light elements, such as nitrogen, is known to help improve these properties further through processes such as amorphization and nitride compound formation. Here, we investigate the ternary CrFeCo system to study the effects of nitrogen addition. (CrFeCo)1-yNy multicomponent thin films are grown on silicon substrates by DC magnetron sputtering. Changes in crystal structure, morphology, mechanical and electrical properties with gradual increases of nitrogen in the film are described and discussed. Increased addition of ni-trogen from 14 at.% to 28 at.% in the film leads to a transformation from an fcc to a bcc crystal structure, affects both the mechanical and electrical properties. XPS analysis shows the tendency of nitrogen to bond with Cr over other metals. The films display hardness values between 7 and 11 GPa with resistivities values ranging between 28 and 165 μΩ cm.

1. Introduction

The need for materials that enhance life span, performance, and economic viability among various other properties has propelled research in alloy design from traditional binary to more complex sys-tems multicomponent alloys. The concept of multicomponent and high entropy alloys (HEA) was theoretically conceived in the 1980s [1,2]. In 2004, two separate studies by Cantor et al. and Yeh et al. experimentally showed that alloys with multi-principal elements form simple crystal structures [3,4]. As the name suggests, HEAs form stable solid solutions due to low Gibbs free energy (ΔG) resulting in a high entropy of mixing (ΔSmix). The same reasoning is not completely applicable to all multi-component alloys, a concept which includes a broader selection including but not limited to HEAs. Multicomponent alloys are attractive candidates for industrial applications not only because of the formation of stable solid solutions but also due to the combination of properties that can be achieved [5,6]. Trends in current research which are shifting from bulk alloys to thin films brings the requirement of an in-depth knowledge on growth processes, phase formation, physical and chemi-cal properties [3,7,8].

Physical vapor deposition (PVD) techniques such as magnetron

sputtering are well established methods for thin film deposition [9]. Several studies have been conducted using sputtering techniques to describe growth kinetics, structure evolution, mechanical, electrical, and chemical properties of multicomponent thin films [10–12].

The addition of light elements to improve mechanical properties started to gain interest with the work on nitride films of multi-element high-entropy alloys in 2004 [13]. Since then numerous studies have been carried out with focus on 3d transition metal (Cr, Mn, Fe, Co, Ni and Cu) multicomponent nitride films. The reason for interest in nitride films is their alleged improvement in mechanical properties. For re-fractory metal HEAs the increase in hardness, wear resistance etc. can be attributed to the formation of stoichiometric nitride solid solutions [14–17]. However, the same cannot be said for the Cantor system i.e. the 3d transition metals belonging to group 7 and 8. This is because most of these metals are weak nitride formers. In this regard, studies on AlCoCrNiNx amorphous high entropy nitride (HEN) films were carried out by Kim et al. to understand the effect of process pressure on microstructure and mechanical properties while varying the nitrogen flow ratio (RN) from 25% to 100%. They report that increasing the process pressure aids in the densification of the film resulting in improved mechanical properties (H, E and H/E) [18]. Recently, Sha * Corresponding author.

E-mail address: smita.gangaprasad.rao@liu.se (S.G. Rao).

Contents lists available at ScienceDirect

Surface & Coatings Technology

journal homepage: www.elsevier.com/locate/surfcoat

https://doi.org/10.1016/j.surfcoat.2021.127059

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et al. reported high values of hardness, wear resistance and scratch response in (FeMnNiCoCr)Nx HEA coatings with fcc crystal structure at low and bcc structures at higher concentrations of nitrogen [19]. They concluded their study by suggesting that nitrogen segregation towards grain boundaries could help in improving mechanical properties of these films.

We see that processes parameters such as reactive gas flow and process pressure, substrate temperature, bias etc. play a role in phase formation. However, most of these studies involve complex systems with five or more metals, making it difficult to understand the behaviour of the material system while changing process parameters such as reactive gas flow ratio. The present study aims to understand the role played by small amounts of nitrogen on the crystal structure, morphology, and mechanical properties of 3d transition metal (Cr, Fe and Co) multi-component thin films. Towards this, thin films of (CrFeCo)Ny were grown on silicon (100) substrates by DC magnetron sputtering. These metals were chosen for three reasons, as a simplified model of the Cantor system, for their similarity in atomic size (Cr = 1.27 Å, Fe = 1.26 and Co =1.25 Å) and their similar sputter yields.

2. Experimental details

The films were deposited using a lab-scale ultrahigh vacuum magnetron sputtering system evacuated to a base pressure < 9 × 10−7 Pa at room temperature before sputter deposition. Depositions were carried out for 30 min at a substrate temperature of 300 ◦C with a floating bias. The system was equipped with four magnetrons inclined at 30◦ posi-tioned at 14 cm from the substrate holder which was rotated at 12 RPM during deposition [20]. Coatings were co-sputtered onto 50-mm-diam-eter Si(100) substrates using three individual targets in a mixture Ar +N2 atmosphere in which the nitrogen flow ratio varied from 1 to 11%. The total pressure was set to 0.4 Pa (3 mTorr). The DC target power for all three targets was set at 100 W regardless of individual target sputter rates. Prior to deposition, the substrates were cleaned by immersion in acetone and isopropanol for 10 min in an ultrasonic bath and finally blow-dried with nitrogen gas.

The crystal structure analysis was carried out by X-ray diffraction (XRD) using a Panalaytical Empyrean X-ray diffractometer. All scans were performed using Cu-Kα radiation (λ = 1.54060 Å) operated at a voltage of 45 kV and a current of 40 mA in a Bragg-Brenteno configu-ration. In order to reduce background caused by fluorescence radiation due to the presence of Fe, Co, and Cr in the film the pulse height dis-tribution (PHD) settings were changed. The PHD levels of the X-cele-rator 1D detector was adjusted such that the lower level was 58.5% and upper level-76.0% [21].

A scanning electron microscope (Leo 1550 Gemini, Zeiss) operated with an acceleration voltage of 12 kV and in-lens detector was used to observe the surface and cross-section morphology. The chemical com-positions of the films were obtained using an energy dispersive spec-trometer (EDS), integrated with the SEM. The analysis was performed at an operating voltage of 20 kV, at a working distance of 8 mm. The EDS data was analysed using the AZtec Nano-analysis software. Further analysis of composition was obtained via time-of-flight elastic recoil detection analysis (ToF-ERDA). Measurements were carried out in a 5 MV NEC-5SDH-2 Pelletron Tandem accelerator at Uppsala University, Sweden. ToF-ERDA results were detected at a 45◦angle between the 36 MeV 127I9+primary beam incident at 67.5◦and a gas ionization chamber detector. Elemental depth profiles were acquired from ToF-ERDA time and energy coincidence spectra using the software package Potku 2.0 [22]. More information regarding data treatment and composition determination is provided in the supplemental information.

X-ray photoelectron spectroscopy (XPS) spectra were acquired in a Kratos Axis Ultra DLD instrument from Kratos Analytical (UK). Mono-chromatic Al Kα radiation (hv = 1486.6 eV) was used with an anode power of 150 W. A 3 × 3 mm2 area of the sample surface was sputter- etched for 120 s with a 4 keV Ar+ ion beam to remove surface

contamination. The Ar+ion energy was later reduced to 0.5 keV for 600 s to minimize surface damage. Spectra were collected from the 0.3 × 0.7 mm2 area positioned at the centre of the etched crater with electrons emitted along the surface normal. An ISO-certified procedure was used to calibrate the binding energy (BE) scale of the spectrometer. To avoid problems related to the use of the C 1s peak of adventitious carbon for charge referencing [23,24] all spectra are aligned to the Fermi edge cut- off. The analyser pass energy was set to 20 eV which resulted in the full width at half maximum of 0.55 eV for the Ag 3d5/2 peak.

Several samples were selected for (Scanning) Transmission Electron Microscopy [(S)TEM] analysis. Cross-sectional TEM specimens were prepared by manual polishing down to a thickness of 60 μm, followed by Ar+ ion milling at 5 keV, with a 6◦incidence angle, on both sides while rotating the sample in a Gatan precision ion polishing system. All ana-lyses were performed using a FEI Tecnai G2 TF 20 UT instrument operated at 200 kV. STEM images were collected with the annular de-tector spanning the range 80 to 260 mrad.

Nanoindentation and scratch testing was performed in a Hysitron Triboindenter 950 equipped with a 2D transducer. The indentation was carried out using a Berkovich tip (100 nm radius) while keeping the indentation depth below one tenth of the thickness of the coating. The hardness (H) and reduced elastic modulus (Er) of the films were calcu-lated according to the Oliver-Pharr method [25]. Prior to measurements, the instrument was calibrated along the indentation axis using a stan-dard fused silica reference sample. A set of 9 indents was carried out in a displacement-controlled mode by adjusting peak displacement while keeping segment times constant with a delay of 30 s between each event. The electrical resistivity of the samples was determined by multi-plying the sheet resistance measured using a four-point probe (Jandel RM3000 station) of the film by the thickness of the film as measured by SEM. The relevant correction factor for 50-mm-diameter substrates was used.

3. Results

Preliminary quantification of film composition was carried out using EDS (standard deviation of 5 at.%). Due to the difficulty in quantifying light elements by EDS, ToF-ERDA was used to obtain the content of nitrogen and oxygen within the sample. Differentiating between the individual metals (Cr, Fe, Co) in the ToF-ERDA measurement data was not possible due to signal overlap (data treatment is presented in the supplementary information in figure S4). However, the nitrogen/metal and oxygen/metal ratios were obtained. The composition obtained from both EDS and ToF-ERDA are given in Table 1. The notation Me1− yNy is used, i.e., Me = (Cr + Fe + Co).

ToF-ERDA reveals a higher presence of oxygen at the surface origi-nating from a possible post-deposition oxidation of the film. The thick-ness of the oxide layer, as estimated from ERDA depth profiles, was between 2 and 7 nm. The oxygen content within the film ranged be-tween 0.6 and 0.8 at.% except for Me0.86N0.14 which showed a slightly higher value of ~2 at.%. Increasing the percentage of N2 in the gas mixture from 1 to 11% lead to an increase of nitrogen content in the film Table 1

Composition and thickness of (Cr, Fe, Co)1-yNy films estimated by combination

EDS and ERDA. Me1− yNy =1. Percentage of N2 (%)

in gas mix Composition from EDS and ERDA (±2 at. %)

Thickness

(nm) Reduced name Me1− yNy Cr Fe Co N 1 25 26 35 14 640 Me0.86N0.14 3.5 21 27 35 17 580 Me0.83N0.17 6 20 26 35 19 545 Me0.81N0.19 8.4 17 24 33 26 530 Me0.74N0.26 11 16 23 33 28 490 Me0.72N0.28

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from 14 at.% to 28 at.%. Variations are noticeable in the metal content with the Cr content decreasing from 24.8 to 16.3 at.%, Fe content decreasing from 26.4 to 22.6 at.%, while Co remained relatively con-stant when the percentage of nitrogen was increased in the plasma.

Fig. 1a presents the θ-2θ XRD patterns of Me1− yNy films deposited on silicon (100) for different fractions of N2 in the gas flow. The Si substrate peaks are observed at 2θ values of 32.9◦and 69.3, corresponding to the 002 (“forbidden”) and 004 reflections [26]. Fig. 1b shows the magnified XRD spectra highlighting the area of interest. The metallic films (MeN0) exhibit a hcp crystal structure with minor contributions of fcc (see figure S2 and S3 in supplementary information). The addition of nitrogen causes a change in the structure as well as the lattice parameter of the films. With the lowest concentration of nitrogen, the film (Me0.86N0.14) exhibits a fcc structure with lattice parameter of 3.590 ± 0.002 Å. For the higher nitrogen content films (Me0.74N0.26 and Me0.72N0.28) a single, broad, intense peak is observed. These films were identified to be pref-erentially oriented along the {110} of a bcc structure with a lattice parameter of 2.849 ± 0.002 Å (Me0.74N0.26) and 2.832 ± 0.002 Å (Me0.72N0.28) similar to that of bcc iron [27]. For the intermediate compositions in nitrogen (Me0.83N0.17 and Me0.81N0.19), the XRD results indicated the presence of both phases, fcc and bcc. The fact that the peaks at 43.4◦ and 50.6move together with addition of nitrogen confirm that they belong to the same crystal structure. Fig. 1b indicates the positions of the main peaks of Cr2N. For the film Me0.83N0.17 a shoulder peak at 42.8◦can be identified as the 111 reflection of Cr2N.

Fig. 2 shows SEM surface and cross-section micrographs of the films deposited on silicon with different nitrogen content in the film (from Me0.86N0.14 to Me0.72N0.28). The film with the lowest nitrogen content has highly faceted, small grained, and a columnar structure growth (Fig. 2a, b). Increasing the nitrogen content of the film from 14 to 28 at. % leads to an increase in the grain size from ~20 nm to ~80 nm. An abrupt change can be observed on the cross section for the film Me0.81N0.19 (deposited with 6% nitrogen flow ratio shown in 2c). Me0.81N0.19 film seems to be relatively smoother while the columnar growth which was present for lower nitrogen content films is less pro-nounced which gives a compact and dense aspect to the film. In addition to the change in morphology with nitrogen incorporation, a decrease in film thickness from 600 nm to 490 nm was also observed (listed in

Table 1).

Cross-sectional TEM and STEM micrographs of Me0.86N0.14, Me0.81N0.19 and Me0.72N0.28, as well as their corresponding selected area electron diffraction (SAED) patterns are presented in Fig. 3a–l. The native silicon oxide layer, with a thickness of ~1.5 nm, was observed at the interface between substrate and film as indicated in Fig. 3b. Cross- section TEM confirmed the observations by SEM where the film with lowest nitrogen content has well-define columnar structure with under dense grain boundaries (Fig. 3a). As the nitrogen content was increased, the columnar growth morphology seemed to diminish giving way to a denser film. The SAED of Me0.86N0.14, Me0.81N0.19, and Me0.72N0.28 are presented in Fig. 3d, h and l. All three films were found be poly-crystalline as seen by the rings of diffraction on the SAED patterns. Me0.86N0.14 film presents distinct spots on the SAED pattern all of which can be located on rings centred on the pattern. The spots/rings are identified as the 111, 200, 220 and 311 reflections of a fcc structure with d-spacing values along the [111] and [200], 1.985 ± 0.005, 1.753 ± 0.005, respectively. The cell parameter was calculated to be 3.507 ± 0.005 Å. The SAED pattern of Me0.81N0.19 film although a bit ambiguous presented a combination of a fcc and bcc structure. The spots and rings present on the SAED pattern of the film Me0.72N0.28 are identified as the (110), (200) and (211) reflections of a bcc phase with a d-spacing of 2.001 ± 0.002, 1.438 ± 0.002 and 1.128 ± 0.002 Å respectively, which correspond to a cell parameter of 2.829 ± 0.004 Å. Note, the brightness of the spots located along the growth direction reveals a low degree of texture along the [110] direction visible on Fig. 3l. As the nitrogen content increased, the film structure evolved from a fcc phase (Fig. 3d), to a fcc + bcc mixed phases (Fig. 3h) and finally a majority bcc phase (Fig. 3l). SAED patterns were in line with the XRD phase analysis (Fig. 1). The occurrence of diffused rings and the (211) plane reflection for Me0.81N0.19 and Me0.72N0.28 is an indication of the decreasing domain size as well as the formation of a bcc structured material.

Fig. 3c, g and k are HAADF-STEM images obtained from films Me0.86N0.14, Me0.81N0.19, and Me0.72N0.28 respectively. Image contrast mechanism in HAADF-STEM images is principally due to variation in mass (composition) and thickness. The contrast variation from dark to light from top to bottom of the image is due to variation in the sample thickness caused by the sample preparation. In the case of Me0.86N0.14

Fig. 1. (a) θ-2θ XRD pattern for (CrFeCo)Ny coatings on Si(100) substrates. Peaks from the metallic film corresponding to hcp phase are marked with ¤. Refer to

supplementary information for more details on MeN0. (b) Magnified XRD pattern of Me1− yNy thin films indicating position of Cr2N peaks in comparison to

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the HAADF-STEM image (Fig. 3c) highlights the under dense grain boundaries. The contrast modulation within the interior is probably the result of compositional variation on the nanoscale. Increasing the ni-trogen content leads to dense films with no grain boundaries as seen for Me0.81N0.19, and Me0.72N0.28 (Fig. 3g and k). This result concurs with SEM cross sectional images which also indicates an increase in density with increasing nitrogen content. Compositional modulations on the nanoscale are observed for Me0.86N0.14 and Me0.72N0.28, very similar to those observed within the grains of Me0.86N0.14.

Fig. 4a–d shows the high-resolution Cr 2p, Fe 2p, Co 2p, and N 1s XPS core-level spectra recorded from Me1− yNy films with varying nitrogen content. Out of all metallic core-levels the largest changes upon N incorporation are observed in Cr 2p spectra, in which case the Cr 2p3/2 spin-split component originally present at 574.19 eV broadens signifi-cantly and moves towards higher binding energy to reach 574.44 eV in the case of Me0.72N0.28 film. Not much change is observed in the Fe 2p spectra upon nitrogen incorporation, the Fe 2p3/2 peak is observed at 706.95 eV for all samples. The Co 2p3/2 peak on the other hand shifts towards lower BE with nitrogen incorporation from 778.40 eV for the MeN0 film to 778.20 eV in the case of Me0.72N0.28 sample [28].

Fig. 5a shows a typical load displacement curve from a nano-indentation measurement carried out on Me0.81N0.19. The penetration depth was set to 50 nm resulting in a maximum load of 1.4 mN. The films undergo a plastic deformation on indentation. Fig. 5b shows the resulting hardness values from Me1− yNy thin film. The films exhibit hardness values between ~7.0 GPa to ~11.0 GPa comparable to hard-ness values obtained for multicomponent metallic as well as nitride films of the 3d transition metal family [29–31]. The film Me0.81N0.19 with the mixed fcc-bcc phase structure exhibited the maximum hardness in the series of samples with values around 11 ± 0.7 GPa.

Fig. 6 shows the change in electrical resistivity of Me1− yNy thin films with increasing nitrogen content in the films. The majority of films exhibited electrical resistivities values between 112 μΩ cm and 165 μΩ

cm similar to other multicomponent metallic and nitride films of the 3d transition metals [13,32,33]. The film Me0.81N0.19 exhibited an elec-trical resistivity of one order of magnitude lower than the other samples with an electrical resistivity of 28 μΩ cm.

4. Discussion

4.1. Composition

Increasing the content of nitrogen in the film from 14 to 28 at.% affects not only the thickness of the film but also the metallic ratios and morphology. The decrease in the thickness could be a result of two factors, an increase in density of the films with nitrogen addition as seen in TEM images Fig. 3a, e and i and/or a decrease in the sputter yield as a result of target poisoning [34]. From EDS and ERDA results shown in Table 1, we notice that the amount of Cr in the film is reduced from 24 to 16 at.%, Fe reduced from 26 to 22 at.%, while Co remained relatively constant as nitrogen flow ratio is increased, suggesting that the Me1− yNy series films are affected by nitrogen compound formation on the target surface. This effect is particularly evident in Cr as it is more prone to form nitride compounds in comparison to Co or Fe. It is important to note here that the nitrogen-containing films are not nitrides; they consist of a metallic-based film with dissolved nitrogen.

4.2. Phase-formation analysis

Multicomponent bulk alloys of 3d transition metals (Cr, Mn, Fe, Co, Ni and Cu) tend to form fcc structures at high temperatures regardless of their composition as it is the thermodynamically favoured phase [35–38] however from the phase diagram of the Fe-Cr-Co bulk system we notice that it is not possible to form a fcc solid solution [39]. Similar behaviour has been seen for the metallic film MeN0, where the films exhibit a hcp crystal structure (Fig. 1) [40–42]. SAED patterns taken Fig. 2. Top view and the corresponding cross-section SEM micrographs of Me1− yNy thin films with nitrogen content increasing from a to e. The SEM images of the

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from the cross-section sample of MeN0 confirms the same and also provides indications of a fcc structure (figure S3 in supplementary in-formation). The formation of fcc on addition of nitrogen in an originally hcp structured film has been observed before and is common for most transition metals. This type of behaviour can be explained by H¨agg’s model [43]. In the present study a similar change has been confirmed at a low percentage of nitrogen (1–3.5%) where the fcc lattice parameter was estimated to be ~3.6 Å (Me0.86N0.14). At these low amounts, the nitrogen is dissolved in the fcc structure but starts to form Cr2N as the concentration is increased as in the case of Me0.83N0.17. Bulk Cantor alloys as well as their variants have been known to form such precipitate Cr2N phases due to the limited solubility of N in the metal matrix [44–46]. As the amount of nitrogen is increased (Me0.74N0.26 and Me0.72N0.28) a phase transformation into a bcc structure occurs along with a change from an under dense columnar to a more dispersed grain structure (Fig. 1). Cross sectional TEM images (Fig. 3a, e, i) show the same. The adatom mobility of certain crystallographic orientations is affected as nitrogen is inserted and hence low degree of texture or preferred orientation persists. This can also be the reason for the faceted growth morphology observed in Fig. 2d and e. The increasing addition of nitrogen may cause distortions to the fcc lattice therefore promoting the nucleation of bcc structured crystallites. Sha et al. reported comparable results on a similar high entropy system (FeMnNiCoCr). At ~6 at.% of nitrogen the films exhibited a fcc structure which evolves into a bcc on increasing the nitrogen content (~15–22 at.%) [19,41].

One possible explanation for the formation of bcc structured in the CrFeCo system could be based on the atomic radii difference. Zhang et al. proposed that the atomic radii difference (δ) plays a role in the

stability and phase formation of multicomponent systems [48]. For example, a study carried out by the combinatorial sputtering of quinary high entropy alloys showed that introducing elements of larger atomic radii such as Al promoted the stabilization of the bcc structure. The reason for this is the lower packing factor of a bcc (68%) in comparison to a fcc structure (74%). The less dense structure allows larger atoms to occupy the lattice without causing distortions [49]. Atomic radii dif-ference (δ) is calculated using the equation,

δ = 100 ̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅̅ ∑n i=1 ci ( 1 − ri r )2

where, ci and ri are the atomic percentage and atomic radius of the ith

element, respectively and r the average atomic radius [48]. Calculating δ for (under stochiometric) nitride films is not as straight forward as with metallic films. We require knowledge of the spin states, oxidation state, their coordination and must consider the presence of an anion in the structure. However, it is possible that as nitrogen is added the overall δ increases, inducing the transformation from fcc to bcc. Theoretically calculated phase diagrams of Fe–Co and Fe-Co-Cr metallic systems also demonstrate the competition between hcp, fcc and bcc phases where the stabilization of a certain phase can depend on the concentration of the elements. This is excluding the effects of local magnetic moments [50,51]. Adding nitrogen into such a system could induce distortions or influence the electronic concentration. Further, the sputter deposition process may also play a role in the stabilization of phases which are in non-equilibrium conditions.

XPS spectra shown in Fig. 4 reveal two major changes upon N incorporation, the broadening and shift of the Cr 2p3/2 peak to 574.22 Fig. 3. TEM (a, e, i), HRTEM (b, f, j) and HAADF- STEM (c, g, k) images taken from cross section samples of Me0.86N0.14, Me0.81N0.19 and Me0.72N0.28 along with their

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eV binding energy and change in the shape of N 1s spectra (see Fig. 4d). In the latter case, for lower N content films a double peak with com-ponents present at 397.8 and 397.1 eV is observed. Such N 1s spectrum is characteristic of Cr2N suggesting that N bonds preferentially to Cr rather than to the other metals. This can be explained based on the heat of formation (ΔHf) values for all related nitrides. Table 2 lists the ΔHf of the concerned metal nitrides calculated at ambient temperature and pressure [52,53]. Theoretical studies by H¨aglund et al. have suggested

that the formation enthalpy of Co–N compounds is higher that Fe–N compounds [54]. Clearly from Table 2, Cr–N compounds are charac-terized by the highest heat of formation and, hence, are more likely to form over other metal nitrides. Correlating with XRD results, the XRD diffractogram of Me0.83N0.17 also presents a shoulder peak at 42.84◦ which corresponds to the 111 reflection of Cr2N. This suggests the for-mation of Cr2N at lower dilution of N into the metal matrix [46,47]. Fig. 4. High-resolution Cr 2p, Co 2p, Fe 2p, and N 1s core-level spectra recorded from Me1− yNy films. Reference line in d correspond to extreme maxima recorded on

the series.

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4.3. Mechanical and electrical properties

Mechanical tests revealed that all films display hardness values on par with bulk materials as well as thin films of similar composition [28–30,45,46]. The film with composition Me0.81N0.19 however shows a higher hardness value (11 ± 2 GPa) which might be an effect of grain boundary, solid solution or phase interfacial strengthening, as this film is composed of a mix of bcc and fcc crystal structure [55-57]. Coinciden-tally, the same film also showed a drop in the electrical resistivity to 28

±5 μΩ cm.

To explain the changes in the mechanical properties, we consider the films Me0.86N0.14, Me0.81N0.19, and Me0.72N0.28. From Fig. 3a we see that the under dense columnar structure is the most probable reason for low hardness of Me0.86N0.14. As the amount of nitrogen increases so does the density, as seen for Me0.81N0.19. Me0.72N0.28 which appears fully dense by SEM, with no apparent grain boundaries or voids similar to Me0.81N0.19, exhibited lower mechanical properties. This decrease in the hardness of Me0.72N0.28 could be attributed to the change in the crystal structure from fcc + bcc to bcc. Also, from the diffraction pattern of Me0.72N0.28 texturation along 〈110〉 can be observed. The fact that indentation takes place along the direction of texture could contribute to the drop in hardness.

HR-TEM images of Me0.81N0.19 and Me0.72N0.28 are shown below in Fig. 7. The images are focused on grain boundaries in both films. In Me0.81N0.19 a dense interface layer is sandwiched between two differ-ently oriented grains as shown by their respective FFT patterns (Fig. 7a). The ring like FFT pattern from the interface layer indicates it is poly-crystalline. This interface acts a compact boundary inhibiting the mo-tion of dislocamo-tions which can result in a higher hardness. For HRTEM image of Me0.72N0.28 the lattice fringes appear distorted in region. This along with presence of Morley fringes, indicating an overlap of two grains (Fig. 7b), is consistent with the direct attachment between the two grains with no apparent interface layer. The drop in the hardness may a result of this difference in the grain interfaces.

Resistivity can be influenced by multiple factors such as, composi-tion, phase constitucomposi-tion, thickness, and stress in the film. As the amount of nitrogen is increased the resistivity is seen to increase suggesting more nitride phases. This trend is followed by all samples except Me0.81N0.19 for which the resistivity drops to 28 μΩ cm. The drop in the resistivity could be an effect of the mixed crystal structure (fcc + bcc), the density of the film or the changing growth morphology.

5. Conclusions

Nitrogen containing multicomponent CrFeCo thin films were grown on silicon substrates by DC magnetron sputtering. Nitrogen content in the films was gradually increased from 14% to 28%. Crystal structure of the films was seen to evolve from a fcc at lower nitrogen at.% (Me0.86N0.14) to a bcc dominant structure at higher nitrogen at.% (Me0.72N0.28). This was also accompanied by a change in microstructure from a columnar to a more dispersed grain structure. SAED patterns reveal that all films were polycrystalline. XPS analysis helped in un-derstanding the bonding mechanism of nitrogen to the metals. N was Fig. 6. Resistivity values of Me1− yNy thin films obtained from sheet resistance.

Error bar of ±5 μΩ cm is hidden within the points. Table 2

Heat of formation values (ΔHf) of binary nitrides calculated at ambient

con-ditions. Values for Co–N compounds are from the materials project and have been theoretically calculated.

Binary metal nitride Heat of formation, ΔHf (kJ/mol)

Cr2N −125.52 CrN − 117.5 Fe3N − 40.00 ± 9 Fe4N −12.17 ± 20 Co2N − 5.88 CoN − 8.20

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found to bond preferentially to Cr rather than to the other metals. The films showed promise in terms of mechanical and electrical properties exhibiting hardness and electrical resistivity values on par with multi-component bulk alloys and thin films of the 3d transition metal series. The highest hardness of 11 ± 0.7 GPa and lowest electrical resistivity of 28 ± 5 μΩ cm were recorded in the film with mixed phase (fcc + bcc) crystal structure.

CRediT authorship contribution statement

Smita G. Rao: Conceptualization, Investigation, Data curation,

Formal analysis, Writing – original draft. Rui Shu: Investigation, Formal analysis. Robert Boyd: Investigation, Formal analysis. Grzegorz

Greczynski: Investigation, Formal analysis. Arnaud le Febvrier:

Conceptualization, Investigation, Supervision, Formal analysis, Writing – review & editing. Per Eklund: Project administration, Conceptuali-zation, Supervision, Writing – review & editing, Funding acquisition.

Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgments

This work is supported by the VINNOVA Competence Centre FunMat-II (grant no. 2016-05156), the Swedish Government Strategic Research Area in Materials Science on Functional Materials at Link¨oping University (Faculty Grant SFO-Mat-LiU No. 2009 00971). The Knut and Alice Wallenberg Foundation is acknowledged for support through the Wallenberg Academy Fellows program (P.E.) and the Electron Micro-scopy Laboratory at Link¨oping University. Daniel Primetzhofer at Uppsala University is acknowledged for assistance with ERDA mea-surements at the Tandem Laboratory infrastructure, financed by the Swedish Research Council via VR-RFI contract C0514401 and by the Swedish Foundation for Strategic Research via SSF contract RIF14-0053. Babak Bakhit at Linkoping University is acknowledged for his assistance with XPS measurements.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi. org/10.1016/j.surfcoat.2021.127059.

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References

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