# Infrared Plasmonics with Conductive Ternary Nitrides

C. Metaxa, S. Kassavetis, J.F. Pierson, D. Gall, P. Patsalas

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C. Metaxa, S. Kassavetis, J.F. Pierson, D. Gall, P. Patsalas. Infrared Plasmonics with Conductive Ternary Nitrides. ACS Applied Materials & Interfaces, 2017, 9 (12), pp.10825-10834. 10.1021/acsmami.6b16343. hal-05032965

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Infrared plasmonics with conductive ternary nitrides

C. *Metaxa*, S. *Kassavetis*, J.F. *Pierson*, D. *Gall*, P. *Patsalas*,*

*Aristotle University of Thessaloniki, Department of Physics, GR-54124 Thessaloniki, Greece

#Institut Jean Lamour (UMR CNRS 7198), Université de Lorraine, Parc de Saurupt, 54011 Nancy

cedex, France

&Rensselaer Polytechnic Institute, Department of Materials Science and Engineering, Troy, NY

12180, USA

*Corresponding author: tel: +30-2310-998298, e-mail: ppats@physics.auth.gr

ABSTRACT: Conductive transition metal nitrides are emerging as promising alternative plasmonic materials that are refractory and CMOS-compatible. In this work, we show that ternary transition metal nitrides of the B1 structure and consisting of a combination of group-IVb transition metal, such as Ti or Zr, and group III (Sc, Y, Al) or group II (Mg, Ca) elements can have tunable plasmonic activity in the infrared range in contrast to Ta-based ternary nitrides, which exhibit plasmonic performance in the visible and UV ranges. We consider the intrinsic quality factors of surface plasmon polariton for the ternary nitrides, and we calculate the dispersion of surface plasmon polariton and the field enhancement at the vicinity of nitride/silica interfaces. Based on these calculations, it is shown that among these nitrides the most promising are TiₓSc₁-ₓN and TiₓMg₁-ₓN. In particular, TiₓSc₁-ₓN can have plasmonic activity in the usual telecom bands at 850, 1300 and 1550 nm. Still, these nitrides exhibit substantial electronic losses that deteriorate the plasmonic field enhancement. This unequivocally calls for improved growth processes that would enable the fabrication of such ternary nitrides of high crystallinity.

KEYWORDS: 1. Surface Plasmon Polariton; 2. Infrared; 3. Conductive Nitrides; 4. Conductor/Dielectric Interfaces.
# I. INTRODUCTION

The emergence of plasmonics has promised radical breakthroughs in various technological sectors, such as microelectronics [1-4], solar energy harvesting [5-9], photo-detection [10-12], optical storage of information [13-16], metamaterials [17-19], and telecommunications [20-24]. The most widely used plasmonic materials are gold and silver, which are not compatible with CMOS processing and their plasmonic response is mostly in the visible spectral range [25,26]. The implementation of plasmonic technology into mainstream telecom applications calls for new plasmonic materials that would be compatible with the established Si and fiber optic technology, i.e. materials that would have plasmonic response at the 850, 1300 and 1550 nm wavelengths [27] and be grown and processed by CMOS-compatible techniques. Conductive transition metal nitrides (TMN), such as TiₓTa₁₋ₓN [28,29], TiₓZr₁₋ₓN (0≤x≤1) [30,31] and TiₓAl₁₋ₓN (0.55<x≤1) [32], provide a combination of unique assets, such as high electron conductivity and mobility [33], high melting point, and CMOS-compatibility [34]. Binary TMNs exhibit strong plasmonic response [35-38], which becomes spectrally tunable for ternary TMNs covering the near UV and the entire visible spectral range [39]. Therefore, alloying of the traditional TMNs, such as TiN, with other elements with two or three valence electrons might shift their plasmonic response to the infrared spectral range. This would combine the assets of plasmonic TMNs with the requirements of the Si telecoms and might establish plasmonics firmly into the telecom technology.

In this work we demonstrate that TiN-based ternary nitrides with rare- (Sc, Y) or alkaline-earth (Mg, Ca) elements can establish a solid platform of conductors whose optical properties and plasmonic performance can be spectrally tunable in the infrared spectral range covering at least the 850 and 1300 nm wavelengths. In particular, we present experimental dielectric function spectra of a variety of continuous opaque polycrystalline films of Ti- or Zr-based TMNs (TiₓRE₁₋ₓN, TiₓAE₁₋ₓN, where RE=Sc, Y and AE=Mg, Ca, and ZrₓY₁₋ₓN; optical data on the latter were completely missing from the
literature) in the infrared range and we correlate their fundamental optical parameters with their chemical composition (x). We also report the optical spectra of epitaxial TiₓSc₁-xN(001)/MgO(001), which are found to show superior plasmonic performance. Finally we complement our study of the optical properties of TiₓAl₁-xN by considering its performance in the infrared, beyond the range reported in Ref. 32. We evaluate their potential plasmonic performance by calculating their intrinsic quality factor (QᵢSPP), the dispersion relations of Surface Plasmon-Polaritons (SPP) at nitride/silica and nitride/silicon interfaces, and the established field enhancement factors (FEF) for SPP [36,40]. We show that the most promising candidate is the TiₓSc₁-xN system, especially in epitaxial form, which can sustain SPP at interfaces with silica and silicon at wavelengths compatible with the current telecom technology (850, 1300, and 1550 nm). A second candidate, which has poorer performance but consists of more abundant elements, is TiₓMg₁-xN. In both cases, the structural defects of the presented samples and their fine grain size, which is mostly due to the low temperature growth (approx. 300 °C) and secondarily due to the used vacuum arc deposition technology [41], reduce their real potential and more efforts should be dedicated to the improvement of the growth of highly crystalline TiₓSc₁-xN and TiₓMg₁-xN on silica and silicon in order to reduce the electron losses and improve both QᵢSPP and FEF.

## II. EXPERIMENTAL

Opaque (thickness >300 nm) polycrystalline TiₓSc₁-xN, TiₓY₁-xN, TiₓMg₁-xN, and TiₓCa₁-xN films were deposited at room temperature on polished stainless steel substrate using a hybrid reactive arc evaporation – magnetron sputtering process. The Ti vapours were created using a cathodic arc source (63-mm diameter and 35-mm thick), while the Sc, Y, Mg or Ca were sputtered from a magnetron cathode (50-mm diameter and 3-mm thick). The deposition of the films was realized in a reactive mixture of argon and nitrogen (40 and 70 sccm, respectively; purity 99.999%). A substrate bias voltage of -100 V was applied during the film growth and the arc current applied to the titanium target was maintained at 90 A, while the composition x was varying by adjusting the pulsed DC current applied to
the sputtering target (Advanced Energy Pinnacle⁺ generator with a 50 kHz frequency and an off time of 4 µs); more details on these experiments are described in Ref. 41. Opaque (1500 nm thick) ZrₓY₁₋ₓN films were grown on floating Si substrates by pulse-DC (Advanced Energy Pinnacle⁺ generator with a 50 kHz frequency and an off time of 4 µs) reactive magnetron sputtering at room temperature using three individual targets (pure Zr, Zr doped with Y 8% or 16%; purity 99.9%) in a mixed Ar/N₂ ambient (30/4 sccm, respectively; purity 99.999%) of a working pressure of 0.44 Pa. Opaque (200 nm thick) epitaxial TiₓSc₁₋ₓN films were grown on 10×10×0.5 mm³ MgO(001) substrates at 750 °C by ultra-high vacuum reactive magnetron sputter deposition in 99.999% pure N₂ discharges at 0.66 Pa. The Sc to Ti ratio was controlled by the relative power applied to two symmetrically positioned magnetrons with 5-cm-diameter Sc (99.9% purity) and Ti (99.999%) targets, respectively. Additionally, spectroscopic data for TiₓAl₁₋ₓN in the infrared range (beyond the range reported in Ref. 32) are also included in this study for comparison purposes.

The composition x of the films, the [N]/[Metal+RE] or [N]/[Metal+AE] ratios, as well as their crystal structure were studied by energy dispersive x-rays (EDX), x-ray photoelectron spectroscopy (XPS), and x-ray diffraction (XRD), respectively. EDX were acquired in a JEOL 840A scanning electron microscope equipped with an Oxford Instruments X-Ray, analyzer; XPS experiments were carried out in a Kratos Axis Ultra DLD instrument equipped with a monochromated Al-Kα beam and a hemispherical sector electron analyzer, a pass energy of 20 eV was used that provided a broadening of less than 500 meV for the Ag 3d line; XRD experiments were carried out in θ-2θ mode using a Bruker (Siemens) D500 apparatus with Co radiation (λ=0.179 nm) for Ti-based polycrystalline films, and a Bruker (Siemens) D5000 apparatus with Cu radiation (λ=0.154 nm) for ZrₓY₁₋ₓN.

The complex dielectric function spectra ($\epsilon=\epsilon_1+i\epsilon_2$) of the opaque polycrystalline films were determined by spectroscopic ellipsometry in the spectral range 0.6-6.5 eV (2067-191 nm) with a step of 50 meV at 70° angle of incidence. Optical transmission and reflection spectra of epitaxial TiₓSc₁₋ₓN were obtained using a Perkin-Elmer Lambda 9 spectrophotometer, equipped with integrating sphere for
light collection, in the spectral range 0.496-6.2 eV (2500-200 nm). Spectral intensity distributions were calibrated using reflection spectra from an undoped single-crystal Si(001) wafer; then the dielectric function spectra were determined from the optical reflectivity following the procedures described in the on-line supporting information and were initially reported in Ref. 28.

### III. RESULTS AND DISCUSSION

Ti/Zr and Sc/Y belong to the groups IVb and IIIb of the periodic table of elements, respectively. As such, they have four or three valence electrons, respectively, in *d* and *s* orbitals. Consequently, their nitrides are conductors and semiconductors, respectively. In particular, the electrical conductivity of all TMNs is due to the excess of *d* electrons of the metal [28-32]. Ti/Zr and Sc/Y form bonds with nitrogen atoms (valence electron configuration 2s²2p³) in a stable cubic rocksalt (B1) phase [30,31,41-43]. As a result, the ternary TiₓSc₁₋ₓN system is stable over the entire compositional range in the B1 structure [41, 44,45] in contrast to the TiₓAl₁₋ₓN system, which can be stabilized in the B1 phase only in a limited compositional range [32]. Therefore, it is expected that the plasmonic properties of TiₓSc₁₋ₓN films might be extended into the infrared range, while the plasmonic properties of TiₓAl₁₋ₓN lie in the deep red range [32]. TiₓY₁₋ₓN has similar electron configuration with TiₓSc₁₋ₓN, but the large lattice mismatch between TiN and YN results in TiₓY₁₋ₓN films of low crystalline quality, which have high electron losses due to scattering at the grain boundaries [41]. For this reason, in this study we investigate also the ZrₓY₁₋ₓN system, which has similar electron configuration with TiₓSc₁₋ₓN and TiₓY₁₋ₓN, but smaller mismatch between ZrN and YN [30,41,42]. On the other hand, TiₓAE₁₋ₓN (AE=Mg, Ca) are stable in the B1 structure in a limited compositional range [41], similar to the TiₓAl₁₋ₓN. However, even relatively small concentrations of Mg or Ca may reduce substantially the carrier density of TiN due to their two valence electrons [41]. So for small concentrations of Mg, or Ca the TiₓAE₁₋ₓN can have conduction electron density below the range that can be achieved by the TiₓAl₁₋ₓN system. Therefore, TiₓAE₁₋ₓN films also have rational perspectives for infrared plasmonic behavior. The abundance of Mg
and Ca, in contrast to the shortage of Sc and Y, is also a significant reason for investigating their
plasmonic performance.

The spectra of the real (ε₁) and imaginary (ε₂) parts of the dielectric function of various TMN
films in the visible (400-700 nm), near-infrared (700-1400 nm), and in the most of the short-
wavelength-infrared (1400-2067 nm) spectral ranges are presented in Figures 1 and 2. In particular, Fig.
1 shows the dielectric function spectra of ternary compounds of a Group IIIb (Sc, Y) and a Group IVb
(Ti, Zr) element. In order to evaluate the effect of the morphology and crystallographic features to the
optical properties in the infrared range of the ternary nitrides we also considered the optical reflectivity
spectra (Fig. 2a) of sputtered, epitaxial TiₓSc₁-xN on MgO(001) substrates. The dielectric function
spectra were extracted from the optical reflectivity as described in the supporting information; the
relevant ε₁ spectra are presented in Fig. 2b and they are characteristic of good conductors with varying
absolute values of ε₁ depending on the composition x. Fig. 3 shows the dielectric function spectra of
ternary compounds of a Group IIa (Mg, Ca) element and Ti. Similar dielectric function spectra for the
TiₓAl₁-xN system, but in a narrower spectral range, can be found in Ref. 39. All the studied films act
optically as semi-infinite materials due to their conductive character and large thickness (>200 nm) with
the major exceptions of the Sc-rich (x=0.19 and 0.045) epitaxial TiₓSc₁-xN, which exhibit interference
fringes indicating their relative transparency. Therefore, these samples will not be considered in the rest
of this study. The ε₁ spectra of TMNs have attracted particular attention because they provide a clear
qualitative view of their conductive character that renders negative ε₁ values [28-30, 36-39]. In all cases
it is evident that the enrichment of TiN or ZrN with a group IIIb (Sc, Y) or a group IIa (Mg, Ca) element
results in less negative ε₁ values in the entire infrared range. In the cases of TiₓMg₁-xN and TiₓCa₁-xN
(Fig. 3a,c) the ε₁ values exhibit a broad plateau beyond 1600 nm due to excessive electron losses that
were quantified in Ref. 41. For short wavelengths, ε₁ gets positive values due to interband transitions
and the accompanying dielectric losses, which are more evident in the ε₂ spectra (Fig. 1b,d,f, and 3b,d);
the dielectric losses, though, are not a source of concern in our study due to their spectral separation from the infrared range.

Of particular practical importance is the screened plasma energy $E_{ps}=\hbar\omega_{ps}=hc/\lambda_{ps}$ ($c$ being the speed of light in free space), which is the spectral energy at which the $\varepsilon_1$ becomes zero ($\varepsilon_1=0$), a phenomenological parameter that may be affected by both the intraband and interband absorption; despite of the lack of any solid physical meaning of $E_{ps}$, it is associated with the spectral energy at which LSPR occurs [38,39]. It is clear from Fig. 4 that $\lambda_{ps} (=hc/E_{ps})$ varies from 450 to almost 770 nm, stretching the potential limits of the plasmonic behavior of ternary conductive nitrides into the infrared spectral range. In particular, $E_{ps}$ is substantially reduced with decreasing the group-IV element's atomic concentration ($x$) for all the studied cases of conductive ternary nitrides.

In the cases of good lattice match group-III/group-IV ternary nitrides (Ti$_x$Sc$_{1-x}$N, Zr$_x$Y$_{1-x}$N, Ti$_x$Al$_{1-x}$N) there is a linear dependence between $E_{ps}$ and $x$ and the slope of the reduction is quite similar for all cases, although the $E_{ps}$ for sputtered Ti-based nitrides (Ti$_x$Al$_{1-x}$N and epitaxial Ti$_x$Sc$_{1-x}$N) is shifted to lower absolute values, which is a common observation for sputtered nitrides [46]. On the contrary, the structurally unstable Ti$_x$Y$_{1-x}$N (due to the lattice mismatch between the constituent TiN and YN phases [41]) does not exhibit a linear dependence of $E_{ps}$ with $x$. Similar linear dependencies of $E_{ps}$ with $x$ are also observed for Ti$_x$Mg$_{1-x}$N and Ti$_x$Ca$_{1-x}$N, albeit with twice the slope. This clearly implies that the $E_{ps}$ is mostly affected by the conduction electron density, whereas the contribution of dielectric losses to it is of secondary importance; given the three valence electrons of Sc, Y, and Al and the two valence electrons of Mg, and Ca the reduction of conduction electrons is faster when we substitute Ti or Zr by atoms of the later than the former elements. In order to quantify that, we consider the effective number $N_{eff}$ of valence electrons of the constituent metals of the ternary nitrides:

$$N_{eff} = x \cdot V_1 + (1-x) \cdot V_2, \qquad (1)$$
where V₁=4 (for Ti and Zr) and V₂=2 and 3 for Mg, Ca, and Sc, Y, Al, respectively. Indeed, E<sub>ps</sub> is directly scaled with N<sub>eff</sub>, as shown in the inset of Fig. 4, unequivocally proving the association of E<sub>ps</sub> with the conduction electron density for all ternary nitrides.

The potential of any conductor to sustain a SPP close to the wavelength λ<sub>ps</sub> (=hc/E<sub>ps</sub>) can be evaluated by the maximum intrinsic quality factor, which is determined by quantities associated with the volume plasmon, and it is defined as [47]:

$$Q_{iSPP} = \frac{\omega_{pu}^2}{2\gamma_D^2}, \qquad (2)$$

where ω<sub>pu</sub> and γ<sub>D</sub> are the unscreened plasma frequency, i.e. the frequency of light where a bulk plasmon with energy E<sub>pu</sub>=ħω<sub>pu</sub> occurs in the conductor, and the Drude broadening factor of the conductor, respectively. ω<sub>pu</sub> and γ<sub>D</sub> are associated with the conduction electron density N<sub>e</sub> [30,32,38,41]:

$$\omega_{pu} = \sqrt{\frac{N_e e^2}{\epsilon_o m^*}}, \qquad (3)$$

where e is the electron charge, ε<sub>o</sub> is the permittivity of free space and m* is the electron effective mass, in SI units, and with the conduction electron relaxation time τ<sub>D</sub> [30,32,38,41]:

$$\gamma_D = \frac{1}{\tau_D}, \qquad (4)$$

respectively, and they can be determined by fitting the dielectric function spectra by a Drude-2 Lorentz model, which determines the complex dielectric function ε<sub>σ</sub>(ω) of the material with the frequency ω of the incoming light [30,32,38,41]:
$$ \tilde{\epsilon}(\omega) = \epsilon_{\text{inf}} - \frac{\omega_{pu}^2}{\omega^2 - i\gamma_D\omega} + \sum_{j=1}^{2} \frac{f_j \cdot \omega_{oj}^2}{\omega_{oj}^2 - \omega^2 + i\gamma_j\omega} \quad (5) $$

In Eq. (5) $\epsilon_{\text{inf}}$ is a background constant, larger than unity, which is due to high-energy contributions (beyond the experimental spectral range) referring to transitions that are not taken into account by the Lorentz term(s). Each of the Lorentz oscillators is describing an interband transition [41] and it is located at an energy position $E_{0j}=\hbar\omega_{0j}$, with strength $f_j$ and damping (broadening) factor $\gamma_j$. Figure 2c summarizes the variations of $E_{pu}$ and $\gamma_D$ for polycrystalline and epitaxial Ti$_x$Sc$_{1-x}$N for various x values, while the values for all the samples are listed in the on-line supporting information.

In an effort to evaluate the intrinsic plasmonic potential of the considered nitride conductors and associate it with the spectral range where SPP can occur, we correlate and display the $Q_{\text{ISPP}}$ values with $E_{ps}$ and the corresponding value of the wavelength $\lambda_{ps}$ in Fig. 5. The $Q_{\text{ISPP}}$ values of all the considered ternary nitrides are orders of magnitude inferior to the values reported for Ag and Au [47]; however, the reported $Q_{\text{ISPP}}$ values for Ag and Au [47] refer to highly crystalline metals with grain size larger than 300 nm [48,49], while the studied nitride samples consist of ultrafine grains less than 20 nm long [32,41]. Therefore, a fair comparison would be to the Ag and Au of the same grain size; in that case, the ternary nitrides are still inferior but their plasmonic performance is comparable in the same order of magnitude to Au and Ag [38,39], as it is also shown in the on-line supporting information.

In Fig. 5 there is a general trend that the redshift of $E_{ps}$, and consequently of the spectral position of SPP, comes in expense of the quality factor $Q_{\text{ISPP}}$. Among the considered ternary nitrides Ti$_x$Sc$_{1-x}$N shows the highest quality factor in the entire spectral range, and only Ti$_x$Mg$_{1-x}$N can be comparable, though only in a narrow spectral window. The general trend of reducing $Q_{\text{ISPP}}$ at longer wavelengths is quite well understood considering Eq. (2). Indeed, the reducing $E_{ps}$ for all cases is associated with the reduction of the conduction electron density due to the smaller number of valence electrons of the group
II or group III elements compared to Ti and Zr, thus reducing $\omega_{pu}$ as well [32,41]. Subsequently and in most cases, $\gamma_D$ is increasing as Ti and Zr are gradually replaced by atoms of the group II or group III elements [32,41] either due to the crystal grain refinement (e.g. as in the case of Ti$_x$Y$_{1-x}$N [41]), or due to structural defects (e.g. as in the cases of Ti$_x$Al$_{1-x}$N [39] and Ti$_x$Sc$_{1-x}$N [41]). The former quantity ($\omega_{pu}$) cannot contribute to further enhancement of the quality factor $Q_{iSPP}$, because it is intentionally reduced in order to shift SPP towards the infrared spectral range. Thus, further work should be dedicated for the reduction of $\gamma_D$ in order to improve $Q_{iSPP}$. This can be achieved by elimination of the structural defects by growth at higher temperature and on lattice-matched substrates; indeed, this is clearly indicated by the substantially higher $Q_{iSPP}$ values for the epitaxial Ti$_x$Sc$_{1-x}$N films grown at high temperature (750 °C) on lattice-match MgO(001) substrates as shown in Fig. 5, however, the improvement is less pronounced in the infrared range. Especially the vacuum arc deposition, which was used in this work to deposit TiN, was found to be inferior and superior to the conventional dc-magnetron sputtering [50] and high-power impulse magnetron sputtering [51], respectively, in terms of the grain size and of the structural defects of conductive nitride films. The importance of the growth optimization is also illustrated by the inferior properties of ZrN compared to TiN in this study in contrast to ZrN and TiN grown by multi-cathode sputtering and reported in Ref. 39, where ZrN was found to be less lossy than TiN. Also, the Zr$_x$Y$_{1-x}$N films were not found to have superior performance compared to Ti$_x$Y$_{1-x}$N, as originally anticipated, due to the existence of a substantial amount of structural defects; indeed, the XRD analysis (for details see on-line supporting information) revealed that the studied Zr$_x$Y$_{1-x}$N exhibit micro-strain (which is associated with structural defects [41,45]) and macro-strain (which is associated with hydrostatic stresses in the films [45, 52,53]) in the ranges 4.1-5.8×10⁻³ and 3.7-6×10⁻³, respectively. The microstrain values of the ZrN nitride of this work (5.8×10⁻³) are substantially higher than that of TiN (3.2×10⁻³ [41]), and the microstrain values of Zr$_x$Y$_{1-x}$N are intermediate between those of Ti$_x$Y$_{1-x}$N and Ti$_x$Sc$_{1-x}$N [41]. Indeed, sputtered ZrN films deposited at the same growth conditions (pressure, substrate bias and target power) with TiN, exhibited substantially higher compressive intrinsic stress [54] due to the
backscattered Ar⁺ ions from the Zr target, the reduced thermalization of Zr atoms in the gas phase and the reduced mobility of Zr adatoms on the active growing surface [54, 55], all of them due to the higher atomic mass of Zr compared to Ti; for all these reasons the growth of pure epitaxial ZrN films was achieved quite recently [56], despite of the growth of ZrN films, which include some epitaxial domains since the late eighties [57]. Consequently, ZrN, which is the basis for ZrₓY₁₋ₓN, is the source of the structural defects in sputtered ZrₓY₁₋ₓN mostly for kinetic reasons, despite of the lattice match between ZrN and YN.

The $Q_{iSPP}$ as expressed in Eq. (2) and as displayed in Fig. 5 vs. $E_{ps}$ provides an indication of the intrinsic plasmonic properties of the conductor as it based on quantities associated with bulk plasmons into the corresponding conductors, and cannot describe accurately the SPP that occurs at conductor dielectric interfaces when illuminated via the dielectric. Therefore, a more realistic view of the potential of ternary conductive nitrides for infrared plasmonics can be extracted from the dispersion relation that correlates the frequency $\omega$ with the wave vector $k_x$ in the direction of propagation of SPP, in such a configuration, i.e. along a conductor/dielectric interface. The general equation that defines the aforementioned relation is the following [58]:

$$k_x = \frac{\omega}{c} \sqrt{\frac{\tilde{\epsilon}_{TN} \cdot \tilde{\epsilon}_D}{\tilde{\epsilon}_{TN} + \tilde{\epsilon}_D}}, \qquad (6)$$

where $\tilde{\epsilon}_{TN}$ and $\tilde{\epsilon}_D$ are the complex dielectric functions of each ternary nitride and dielectric (silica or silicon in our case), respectively. Given that the dielectric functions are inherently complex, Eq. (6) should be modified to:

$$k_x = \frac{\omega}{c} \mathrm{Re} \left( \sqrt{\frac{\tilde{\epsilon}_{TN} \cdot \tilde{\epsilon}_D}{\tilde{\epsilon}_{TN} + \tilde{\epsilon}_D}} \right), \qquad (7)$$
when optical attenuation in the conductor exists; this is very important for the conductive nitrides, due to their strong electron losses. A more convenient way of writing the dispersion relation in order to be more easily compared to the experiments is:

$$ \hbar k_x c = \hbar \omega \cdot Re \left( \sqrt{\frac{\tilde{\epsilon}_{TN} \cdot \tilde{\epsilon}_D}{\tilde{\epsilon}_{TN} + \tilde{\epsilon}_D}} \right). \qquad (8) $$

Figure 6 presents the calculated (via Eq. 8) SPP dispersion relations at the interfaces between air, silica or silicon and (a) TiₓSc₁₋ₓN, (b) TiₓY₁₋ₓN, (c) TiₓMg₁₋ₓN, (d) ZrₓY₁₋ₓN, (e) TiₓAl₁₋ₓN, (f) TiₓCa₁₋ₓN, respectively, using the dielectric function spectra presented in Figures 1, 2 and 3 for the conductive nitrides, reported in Ref. 59 for a commercial pure n-type Si(001) wafer after removal of the native oxide and reconstruction of its surface, and reported in Ref. 60 for silica glass. Note that the calculation for silicon was performed only for the transparent region, i.e. for photon energies in free space below the band gap of 1.112 eV (wavelength longer than 1115 nm), which lies exclusively in the infrared spectral range.

The first significant observation is that in all cases the dispersion relations are characteristic of lossy metals and, therefore, involve finite extreme values of $k_x$ [58] due to conduction electron (due to finite $\gamma_D$ values of all ternary nitrides) and dielectric (due to the existence of interband transitions) losses. The consequence of this observation is the existence of quasi-bound modes at which the slope of the dispersion curve is inversed and the existence of a cross-over energy at which the SPP dispersion curve crosses the photon line (black lines in all panels of Fig. 6).

Secondly, the maximum $\hbar k_x c_{max}$, which defines the SPP point, is scaled to the screened plasma energy $E_{ps}$ for all the ternary nitrides in the infrared spectral range; this is clearly illustrated in Fig. 7 where the free space wavelength of light at SPP (i.e. at $\hbar k_x c_{max}$) is plotted vs. $E_{ps}$ for the silica dielectric. However, these two quantities do not coincide and the wavelength of photons in free space that
corresponds to $\hbar k_x c_{max}$ is longer than the wavelength $\lambda_{ps}$ that corresponds to $E_{ps}$ and fulfills the requirement for operation at 850 nm for silica fiber optics when $\text{Ti}_{0.32}\text{Sc}_{0.68}\text{N}$ is used, while $\text{Ti}_{0.73}\text{Mg}_{0.27}\text{N}$ and $\text{Ti}_{0.75}\text{Ca}_{0.25}\text{N}$ are also pretty close to fulfill this requirement possibly with a small further enrichment with Mg or Ca. This is also illustrated in Fig. 8a, where the SPP wavelength is depicted vs. the free space wavelength of light at the SPP point (i.e. at $\hbar k_x c_{max}$); the SPP wavelength is shorter than the wavelength of light, as expected, although their difference is not very large (as in the case of Ag) due to the substantial electron losses.

Thirdly, air can hardly sustain a SPP mode for all cases, while clear SPP is observed for silica albeit at shorter wavelengths than the 850 nm telecom band in most cases. The inability to meet the requirement of 850 nm wavelength is not a deficit of the experimental range of the current work, but it is also due to physical limitations for various ternary conductive nitrides; for example further enrichment with Al or Y in order to extend the range of operation of $\text{Ti}_x\text{Al}_{1-x}\text{N}$ and $\text{Ti}_x\text{Y}_{1-x}\text{N}$ would cause a phase transition from the B1 phase to pseudo-wurtzite [32] or further grain refinement due to the lattice mismatch between TiN and YN [41], respectively, while $\text{Ti}_x\text{Ca}_{1-x}\text{N}$ is characterized by inherent excessive electron loss [41]. The only family of conductive nitrides that can sustain a SPP for the silicon telecom bands (1300 nm and 1550 nm) is the Sc-rich $\text{Ti}_x\text{Sc}_{1-x}\text{N}$, while for the silica fiber band (850 nm) both $\text{Ti}_x\text{Sc}_{1-x}\text{N}$ and $\text{Ti}_x\text{Mg}_{1-x}\text{N}$ might be considered. The later is inferior in terms of the quality factor $Q_{iSPP}$ but given the abundance of Mg, in contrast to the shortage of Sc, it is worth dedicating efforts for the improvement of its growth and plasmonic performance in order to implement it industrially.

Fourthly, the differences between the wavenumbers $k_x$ of the SPP and of the photon in the same dielectric (black lines in all panels of Fig. 6) can be used as a measure of the electronic loss; the bigger the difference the smaller the loss and the better the plasmonic performance of the corresponding conductor. This is quantified in Fig. 8b, where the difference $\lambda_{SPP}-\lambda_{photon}$ is plotted vs. the free space wavelength of light; $\lambda_{SPP}$ and $\lambda_{photon}$ are the wavelengths of SPP (at $\hbar k_x c_{max}$) and of the photons in the silica dielectric, respectively. This difference is reduced at longer wavelengths for all the considered
ternary nitrides proving that the redshift of the plasmonic performance comes in expense of the electron loss, making the manifestation of SPP quite ambiguous [note however that the manifestation of LSPR is more pronounced due to its finite lineshape; indeed, LSPR is clearly observed for the studied nitrides - more details are presented in the on-line supporting information - thus proving the plasmonic potential of these materials]. The same behavior is also observed for $Q_{iSPP}$ vs. the free space wavelength of light (Fig. 8c), meaning that the intrinsic characteristics ($Q_{iSPP}$) define to a great extend the extrinsic plasmonic performance ($\lambda_{SPP}-\lambda_{photon}$). Therefore, the considerations expressed in the previous paragraphs regarding the improvement of $Q_{iSPP}$ via improving the growth process of $Ti_xSc_{1-x}N$ and $Ti_xMg_{1-x}N$ would establish a guide towards the improvement of infrared plasmonic devices incorporating such nitrides, as well.

For most applications the enhancement of the electric field at the vicinity of the conductor's surface due to SPP is crucial and, therefore, should be considered for the evaluation of the studied ternary nitrides. According to Ref. 40 and 61 the field enhancement factor (FEF) can be calculated using the dielectric function spectra of the conductor (in our case the ternary nitrides) and of the dielectric (e.g. silica) via the equation:

$$FEF = \frac{\varepsilon_{1,TN} + \varepsilon_{1,d}}{\varepsilon_{1,TN} \cdot \varepsilon_{1,d}} \times \frac{\varepsilon_{1,TN}^2}{\varepsilon_{2,TN}}, \qquad (9)$$

where $\varepsilon_{1,TN}$ and $\varepsilon_{2,TN}$ are the real and imaginary parts of the dielectric function of each ternary nitride, and the $\varepsilon_{1,d}$ is the real part of the dielectric function of the adjacent dielectric (in our case $\varepsilon_{1,d}=\varepsilon_{1,Silica}=1.9$). Figure 9 displays the calculated FEF for the most promising cases of ternary nitrides with plasmonic activity in the infrared range, namely $Ti_xSc_{1-x}N$ and $Ti_xMg_{1-x}N$ with solid symbols/continuous lines, and open symbols/dashed lines, respectively. The parentheses show the equivalent $N_{eff}$ for each x value. Note that there is no field enhancement (FEF<1) for $Ti_{0.73}Mg_{0.27}N$. The
FEF confirms the better plasmonic performance of TiₓSc₁₋ₓN, which is, however, still inferior to that of TiN or Au [40]. It also validates the gradual deterioration of the plasmonic activity with increasing enrichment with Sc. We attribute again this behavior to the crystal grain refinement and microstrain enhancement (which is due to structural defects) with increasing Sc or Mg content in the ternary nitride films. Therefore, the perspectives of these nitrides in infrared plasmonics are promising but it is essential to improve the crystalline quality of such refractory compounds. This would alter our perception and priorities in the field of the growth of conductive nitrides, which was so far mostly dictated by the requirements of ultra-fine grains in nanostructured nitride films, which enhance their mechanical performance [62-69] albeit in expense of their optical properties.

## IV. CONCLUSIONS

In this work we investigated the potential of a wide variety of conductive ternary nitrides based on TiN for plasmonic applications in the infrared range. In particular we considered the alloying of TiN with group II and group III elements to from ternary compounds such as TiₓSc₁₋ₓN, TiₓY₁₋ₓN, TiₓMg₁₋ₓN and TiₓCa₁₋ₓN; we also considered TiₓAl₁₋ₓN and ZrₓY₁₋ₓN for comparison. We calculated the intrinsic quality factors for SPP of these nitrides, based on their unscreened plasma and conduction electron relaxation time via the Drude broadening. Their plasmonic performance lies in the infrared range and is substantially redshifted with reducing the effective number of valence electrons of the constituent elements. Among them, the most promising are TiₓSc₁₋ₓN and TiₓMg₁₋ₓN due to their stability and their inherent better microstructure; especially, Sc-rich TiₓSc₁₋ₓN can meet the spectral requirements for plasmonics in the usual telecom bands (850, 1300, 1550 nm). Still, the present samples are characterized by substantial electron losses, which deteriorate the field enhancement at the vicinity of the nitride/dielectric interface. This calls for improved growth technologies for epitaxial nitride films, like those presented in Ref. 44,or highly-crystalline nitrides films on silicon and silica conductive nitrides, which contrast the trends of the last years regarding the development of nanostructured nitride coatings.
## ACKNOWLEDGEMENTS

This work has been partially supported by the French National Research Agency (ANR) under the contract ANR-07-MAPR-0007 (COCTAIL) and by the US National Science Foundation under grant No. 1309490. Dr. D. Pilloud and Dr. A. Hodroj (Institut Jean Lamour) are acknowledged for their contribution in the nitride thin films deposition.
FIGURE CAPTIONS

Fig. 1: Spectra of the real ($\epsilon_1$) and imaginary ($\epsilon_2$) parts of the dielectric function of opaque films (>300 nm thick) measured in the Visible-infrared ranges by ellipsometry for (a),(b) $Ti_xSc_{1-x}N$, (c),(d) $Ti_xY_{1-x}N$, and (e),(f) $Zr_xY_{1-x}N$.

Fig. 2: (a) Optical reflectivity spectra of epitaxial $Ti_xSc_{1-x}N$ films grown on MgO(001) substrates, (b) $\epsilon_1$ spectra of the same films, (c) the variation of the Drude parameters $E_{pu}$ and $\gamma_D$ with x for polycrystalline and epitaxial $Ti_xSc_{1-x}N$.

Fig. 3: Spectra of the real ($\epsilon_1$) and imaginary ($\epsilon_2$) parts of the dielectric function of opaque films (>300 nm thick) measured in the Visible-infrared ranges by ellipsometry for (a),(b) $Ti_xMg_{1-x}N$, (c),(d) $Ti_xCa_{1-x}N$.

Fig. 4: The correlation of the screened plasma energy $E_{ps}$, and of the associated wavelength ($\lambda_{ps}=hc/E_{ps}$) with the composition x for various conductive ternary nitrides in the B1 structure; the inset demonstrates the scaling of $E_{ps}$ with effective number $N_{eff}$ of the valence electrons of the constituent metals.

Fig. 5: The correlation of the maximum intrinsic quality factor for SPP ($Q_{iSPP}$) vs. the $E_{ps}$ and the equivalent wavelength $\lambda_{ps}$ for the various nitrides studied in this work. The solid black line is a guide to the eye for the optimal case of $Ti_xSc_{1-x}N$.

Fig. 6: SPP dispersion relations at the nitride/dielectric interfaces for air and the two most widely used dielectrics in telecom technologies, i.e. silica and silicon. The black lines are the dispersion relations of the light into the dielectrics. Horizontal grey lines indicate the three bands used for telecommunications with today's technology.

Fig. 7: The correlation of the free space wavelength of light at SPP, which corresponds to the maximum $\hbar k_x c_{max}$, with the screened plasma energy $E_{ps}$ for the bare conductive nitrides, as well as of the equivalent wavelength $\lambda_{ps}$, respectively.
Fig. 8: The variations of the (a) SPP wavelength, (b) the difference between the wavelengths of SPP and photon at the interface and in the silica dielectric, respectively, and (c) of the intrinsic $Q_{iSPP}$ for the various nitrides studied in this work.

Fig. 9: The field enhancement factor for SPP at the Ti$_{x}$Sc$_{1-x}$N/silica (solid symbols/continuous lines) and Ti$_{x}$Mg$_{1-x}$N/silica (open symbols/dashed lines) interfaces for various x; in parentheses are the equivalent $N_{eff}$ for each case. Note that there is no field enhancement ($FEF<1$) for Ti$_{0.73}$Mg$_{0.27}$N.
REFERENCES

1. Dionne, J.A.; Diest, K.; Sweatlock, L.A.; Atwater, H.A. PlasMOStor: A Metal–Oxide–Si Field Effect Plasmonic Modulator. *Nano Lett.* **2009**, *9*, 897–902.

2. Wang, L.; Hu, W.; Wang, J.; Wang, X.; Wang, S.; Chen, X.; Lu, W. Plasmon resonant excitation in grating-gated AlN barrier transistors at terahertz frequency. *Appl. Phys. Lett.* **2012**, *100*, art. no. 123501.

3. Heeres, R.W.; Kouwenhoven, L.P.; Zwiller, V. Quantum interference in plasmonic circuits. *Nature Nanotech.* **2013**, *8*, 719-722.

4. Sheldon, M.T.; De Groep, J.V.; Brown, A.M.; Polman, A.; Atwater, H.A. Plasmoelectric potentials in metal nanostructures. *Science* **2014**, *346*, 828-831.

5. Atwater H.A.; Polman A. Plasmonics for Improved Photovoltaic Devices. *Nature Mater.* **2010**, *9*, 205-213.

6. Aubry A.; Lei D.Y.; Fernandez-Dominguez A.I.; Sonnefraud Y.; Maier S.A.; Pendry J.B. Plasmonic Light-Harvesting Devices over the Whole Visible Spectrum. *Nano Lett.* **2010**, *10*, 2574-2579.

7. Gan, Q.; Bartoli, F.J.; Kafafi, Z.H. Plasmonic-enhanced organic photovoltaics: Breaking the 10% efficiency barrier. *Adv. Mater.* **2013**, *25*, 2385-2396.

8. Beliatis, M.J.; Henley, S.J.; Han, S.; Gandi, K.; Adikaari, D.; Stratakis, E.; Kymakis, E.; Silva, S.R.P. Organic solar cells with plasmonic layers formed by laser nanofabrication. *Phys. Chem. Chem. Phys.* **2013**, *15*, 8237-8244.

9. White, T.P.; Catchpole, K.R. Plasmon-enhanced internal photoemission for photovoltaics: Theoretical efficiency limits. *App. Phys. Lett.* **2012**, *101*, art. no. 073905.

10. Sobhani, A.; Knight, M.W.; Wang, Y.; Zheng, B.; King, N.S.; Brown, L.V.; Fang, Z.; Nordlander, P.; Halas, N.J. Narrowband photodetection in the near-infrared with a plasmon-induced hot electron device. *Nature Comm.* **2013**, *4*, art. no. 1643.
11. Konstantatos, G.; Sargent, E.H. Nanostructured materials for photon detection, *Nature Nanotech.* **2010**, *5*, 391-400.

12. Chalabi, H.; Schoen, D.; Brongersma, M.L. Hot-electron photodetection with a plasmonic nanostripe antenna. *Nano Lett.* **2014**, *14*, 1374-1380.

13. Montelongo, Y.; Tenorio-Pearl, J.O.; Williams, C.; Zhang, S.; Milne, W.I.; Wilkinson, T.D. Plasmonic nanoparticle scattering for color holograms, *Proc. Nat. Acad. Sci. USA*, **2014**, *111*, 12679-12683.

14. Siozios A.; Koutsogeorgis D.C.; Lidorikis E.; Dimitrakopoulos G.P.; Kehagias T.; Zoubos H.; Komninou P.; Cranton W.M.; Kosmidis C.; Patsalas P. Optical Encoding by Plasmon-Based Patterning: Hard and Inorganic Materials Become Photosensitive. *Nano Lett.* **2012**, *12*, 259-263.

15. Zijlstra P.; Chon J.W.M.; Gu M. Five-Dimensional Optical Recording Mediated by Surface Plasmons in Gold Nanorods. *Nature* **2009**, *459*, 410-413.

16. Kumar K.; Duan H.; Hegde R.S.; Koh S.C.W.; Wei J.N.; Yang J.K.W. Printing Colour at the Optical Diffraction Limit. *Nature Nanotech.* **2012**, *7*, 557-561.

17. Pendry, J.B.; Martin-Moreno, L.; Garcia-Vidal, F.J. Mimicking Surface Plasmons with Structured Surfaces. *Science* **2004**, *305*, 847-848.

18. Boltasseva, A. Empowering plasmonics and metamaterials technology with new material platforms. *MRS Bulletin* **2014**, *39*, 461-468.

19. Boltasseva, A.; Shalaev, V.M. Fabrication of optical negative-index metamaterials: Recent advances and outlook. *Metamaterials* **2008**, *2*, 1-17.

20. Sandtke, M.; Kuipers, L. Slow guided surface plasmons at telecom frequencies. *Nature Phot.* **2007**, *1*, 573-576.

21. Burgos, S.P.; Lee, H.W.; Feigenbaum, E.; Briggs, R.M.; Atwater, H.A. Synthesis and characterization of plasmonic resonant guided wave networks. *Nano Lett.* **2007**, *14*, 3284-3292.
22. Shin, W.; Cai, W.; Catrysse, P.B.; Veronis, G.; Brongersma, M.L.; Fan, S. Broadband sharp 90-degree bends and T-splitters in plasmonic coaxial waveguides. *Nano Lett.* **2013**, *13*, 4753-4758.

23. Li, X.; Jiang, T.; Shen, L.; Deng, X. Subwavelength guiding of channel plasmon polaritons by textured metallic grooves at telecom wavelengths. *Appl. Phys. Lett.* **2013**, *102*, art. no. 031606.

24. Holmgaard, T.; Bozhevolnyi, S.I.; Markey, L.; Dereux, A.; Krasavin, A.V.; Bolger, P.; Zayats, A.V. Efficient excitation of dielectric-loaded surface plasmon-polariton waveguide modes at telecommunication wavelengths. *Phys. Rev. B* **2008**, *78*, art. no. 165431.

25. Lenci, S.; De Jaeger, B.; Carbonell, L.; Hu, J.; Mannaert, G.; Wellekens, D.; You, S.; Bakeroot, B.; Decoutere, S. Au-free AlGaN/GaN power diode on 8-in Si substrate with gated edge termination. *IEEE Electr. Dev. Lett.* **2013**, *34*, 1035-1037.

26. Babicheva, V.E.; Kinsey, N.; Naik, G.V.; Ferrera, M.; Lavrinenko, A.V.; Shalaev, V.M.; Boltasseva, A. Towards CMOS-compatible nanophotonics: Ultra-compact modulators using alternative plasmonic materials. *Opt. Expr.* **2013**, *21*, 27326-27337.

27. Nick, M. Fiber Optic Telecommunication (Module 1.8), in *Fundamentals of Photonics*, Roychoudhuri, C., Ed.; SPIE the international society for optics and photonics: Bellingham- Washington, **2000**; pp.298.

28. Matenoglou, G.M.; Lekka, Ch.E.; Koutsokeras, L.E.; Karras, G.; Kosmidis, C.; Evangelakis, G.A.; Patsalas, P. Structure and electronic properties of conducting, ternary TiₓTa₁-ₓN films. *J. Appl. Phys.* **2009**, *105*, art. no. 103714.

29. Koutsokeras, L.E.; Abadias, G.; Lekka, Ch.E.; Matenoglou, G.M.; Anagnostopoulos, D.F.; Evangelakis, G.A.; Patsalas, P. Conducting transition metal nitride thin films with tailored cell sizes: The case of δ-TiₓTa₁-ₓN, *Appl. Phys. Lett.* **2008**, *93*, art. no. 011904.

30. Matenoglou, G.M.; Koutsokeras, L.E.; Patsalas, P. Plasma energy and work function of conducting transition metal nitrides for electronic applications. *Appl. Phys. Lett.* **94**, art. no. 152108.
31. Matenoglou, G.M.; Koutsokeras, L.E.; Lekka, Ch.E.; Abadias, G.; Kosmidis, C.; Evangelakis, G.A.; Patsalas, P. Structure, stability and bonding of ternary transition metal nitrides. *Surf. Coat. Technol.* **2009**, *204*, 911-914.

32. Kassavetis, S.; Abadias, G.; Vourlias, G.; Bantsis, G.; Logothetidis, S.; Patsalas, P. Optical properties of TiₓAl₁-xN thin films in the whole compositional range. *Surf. Coat. Technol.* **295**, 125-129.

33. Evangelou, E.K.; Konofaos, N.; Aslanoglou, X.A.; Dimitriadis, C.A.; Patsalas, P.; Logothetidis, S.; Kokkoris, M.; Kossionides, E.; Vlastou, R.; Groetschel, R. Characterization of magnetron sputtering deposited thin films of TiN for use as a metal electrode on TiN/SiO₂/Si metal-oxide-semiconductor devices. *J. Appl. Phys.* **2000**, *88*, 7192-7196.

34. Dimitriadis, C.A.; Lee, J.I.; Patsalas, P.; Logothetidis, S.; Tassis, D.H.; Brini, J.; Kamarinos, G. Characteristics of TiN/x-n-Si Schottky diodes deposited by reactive magnetron sputtering. *J. Appl. Phys.* **1999**, *85*, 4238-4242.

35. Cortie, M.B.; Giddings, J.; Dowd, A. Optical properties and plasmon resonances of titanium nitride nanostructures. *Nanotechnology* **2010**, *21*, art. no. 115201 (2010).

36. Naik, G.V.; Shalaev, V.M.; Boltasseva, A. Alternative Plasmonic Materials: Beyond Gold and Silver. *Adv. Mater.* **2013**, *25*, 3264-3294.

37. Naik, G.V.; Saha, B.; Liu, J.; Saber, S.M.; Stach, E.A.; Irudayaraj, J.M.K.; Sands, T.D.; Shalaev, V.M.; Boltasseva, A. Epitaxial superlattices with titanium nitride as a plasmonic component for optical hyperbolic metamaterials. *Proc. Nat. Acad. Sci. USA* **2014**, *111*, 7546-7551.

38. Patsalas, P.; Kalfagiannis, N.; Kassavetis, S. Optical Properties and Plasmonic Performance of Titanium Nitride. *Materials* **2015**, *8*, 3128-3154.

39. Kassavetis, S.; Bellas, D.; Abadias, G.; Lidorikis, E.; Patsalas, P. Plasmonic spectral tunability of conductive ternary nitrides. *Appl. Phys. Lett.* **2016**, *108*, 263110.
40. West, P.R.; Ishii, S.; Naik, G.V.; Emani, N.K.; Shalaev, V.M.; Boltasseva, A. Searching for better plasmonic materials. *Laser Phot. Rev.* **2010**, *4*, 795-808.

41. Kassavetis, S.; Hodroj, A.; Metaxa, C.; Logothetidis, S.; Pierson, J.F.; Patsalas, P. Optical and Electronic Properties of Conductive Ternary Nitrides with Rare- or Alkaline-Earth Elements. *J. Appl. Phys.* **2016**, *120*, art. no. 225106.

42. Stampfl, C.; Mannstadt, W.; Asahi, R.; Freeman, A.J. Electronic structure and physical properties of early transition metal mononitrides: Density-functional theory LDA, GGA, and screened-exchange LDA FLAPW calculations. *Phys. Rev. B* **2001**, *63*, art. no. 155106.

43. Gregoire, J.M.; Kirby, S.D.; Turk, M.E.; van Dover, R.B. Structural, electronic and optical properties of (Sc,Y)N solid solutions. *Thin Solid Films* **2009**, *517*, 1607-1609.

44. Gall, D.; Petrov, I.; Greene, J.E. Epitaxial Sc<sub>1-x</sub>Ti<sub>x</sub>N(001): Optical and Electronic Transport Properties. *J. Appl. Phys.* **2001**, *89*, 401-409.

45. Gall, D.; Stoehr, M.; Greene, J.E. Vibrational Modes in Epitaxial Sc<sub>1-x</sub>Ti<sub>x</sub>N(001) Layers: An Ab-initio Calculation and Raman Spectroscopy Study. *Phys. Rev. B.* **2001**, *64*, art. no. 174302.

46. Koutsokeras, L.E.; Abadias, G.; Patsalas, P. Texture and microstructure evolution in single-phase Ti<sub>x</sub>Ta<sub>1-x</sub>N alloys of rocksalt structure. *J. Appl. Phys.* **2011**, *110*, art. no. 043535.

47. Blaber, M.G.; Arnold M.D.; Ford, M.J. A review of the optical properties of alloys and intermetallics for plasmonics. *J. Phys.: Condens. Matter.* **2010**, *22*, art. no. 143201.

48. Johnson P.B.; Christy, R.W. Optical Constants of the Noble Metals. *Phys. Rev. B* **1972**, *6*, 4370-4379.

49. Adamov, M.; Perović, B.; Nenadović, T. Electrical and structural properties of thin gold films obtained by vacuum evaporation and sputtering. *Thin Solid Films* **1974**, *24*, 89-100.

50. Gautier, C.; Moussaoui, H.; Elstner, F.; Machet, J. Comparative study of mechanical and structural properties of CrN films deposited by d.c. magnetron sputtering and vacuum arc evaporation. *Surf. Coat. Technol.* **1996**, *86-87*, 254-262.
51. Purandare, Y.; Ehiasarian, A.; Santana, A.; Hovsepian, P. ZrN coatings deposited by high power impulse magnetron sputtering and cathodic arc techniques. *J. Vac. Sci. Technol. A* **2014**, 32, art. no. 031507.

52. Abadias, G. Stress and preferred orientation in nitride-based PVD coatings. *Surf. Coat. Technol.* **2008**, 202, 2223-2235.

53. Patsalas, P.; Gravalidis, C.; Logothetidis S. Surface kinetics and subplantation phenomena affecting the texture, morphology, stress, and growth evolution of titanium nitride films. *J. Appl. Phys.* **2004**, 96, 6234-6246.

54. Abadias, G.; Koutsokeras, L.E.; Guerin, Ph.; Patsalas, P. Stress evolution in magnetron sputtered Ti–Zr–N and Ti–Ta–N films studied by in situ wafer curvature: Role of energetic particles. *Thin Solid Films* **2009**, 518, 1532–1537.

55. Koutsokeras, L.E.; Abadias, G. Intrinsic stress in ZrN thin films: Evaluation of grain boundary contribution from in situ wafer curvature and ex situ x-ray diffraction techniques. *J. Appl. Phys.* **2012**, 111, art. no. 093509.

56. Mei, A.B; Howe, B.M.; Zhang, C.; Sardela, M.; Eckstein, J.N.; Hultman, L.; Rockett, A.; Petrov I.; Greene, J.E. Physical properties of epitaxial ZrN/MgO(001) layers grown by reactive magnetron sputtering. *J. Vac. Sci. Technol. A* **2013**, 31, art. no. 061516

57. Barnett, S.A.; Hultman, L.; Sundgren, J.-E.; Ronin, F.; Rohde, S. Epitaxial growth of ZrN on Si(100). *Appl. Phys. Lett.* **1988**, 53, 400-403

58. Maier, S.A. *Plasmonics: Fundamentals and Applications*, Springer Science and Business Media LLC: New York, USA; **2007**; pp 65-87.
59. Patsalas, P.; Logothetidis, S. Interface properties and structural evolution of TiN/Siand TiN/GaN heterostructures. *J. Appl. Phys.* **2003**, *93*, 989-998.

60. Malitson, I.H. Interspecimen Comparison of the Refractive Index of Fused Silica. *J. Opt. Soc. Am.* **1965**, *55*, 1205-1208.

61. Shalaev, V.M. Fractal nano-composites: Giant local-field enhancement of optical responses. *AIP Conf. Proc.* **2001**, *560*, 237-258.

62. Vepřek, S. The search for novel, superhard materials, *J. Vac. Sci. Technol. A* **1999**, *17*, 2401-2420.

63. Musil, J. Hard and superhard nanocomposite coatings, *Surf. Coat. Technol.* **2000**, *125*, 322-330.

64. Sproul, W.D. New routes in the preparation of mechanically hard films, *Science* **1996**, *273*, 889-892.

65. Musil, J.; Hrubý, H. Superhard nanocomposite Ti₁-xAlxN films prepared by magnetron sputtering. *Thin Solid Films* **2000**, *365*, 104-109.

66. Patscheider, J. Nanocomposite hard coatings for wear protection. *MRS Bulletin* **2003**, *28*, 180-183.

67. Sproul, W.D. Physical vapor deposition tool coatings. *Surf. Coat. Technol.* **1996**, *81*, 1-7.

68. Mayrhofer, P.H.; Mitterer, C.; Hultman, L.; Clemens, H. Microstructural design of hard coatings, *Progr. Mater. Sci.* **2006**, *51*, 1032-1114.

69. Mayrhofer, P.H.; Hörling, A.; Karlsson, L.; Sjölén, J.; Larsson, T.; Mitterer, C.; Hultman, L. Self-organized nanostructures in the Ti-Al-N system, *Appl. Phys. Lett.* **2003**, *83*, 2049-2051.
<figure><img src="image_2.png" /><figcaption>Fig. 1</figcaption></figure>
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