Z. Cao,     Chinese Academy of Sciences, P. R. China

8.1 Introduction

8.2 Deposition of stoichiometric Cu₃N

We adopted the radio frequency (27.12 MHz) reactive magnetron sputtering method to grow Cu₃N thin films on Si (100) wafers in a custom-designed system, with a gas mixture of argon and nitrogen. The high-purity (5 N) Cu target is 60 mm in diameter. In order to suppress the nitrogen re-emission by heat, the substrate was held at 60 °C just to achieve adequate adhesion for the deposits. During reactive sputtering growth, the RF power was tuned to a maximum of 150 W, and the working pressure was maintained at 9.0 × 10⁻¹ Pa with a total gas flow rate of 4.0 sccm. The proportion of N₂ in the working gas is another variable for the optimization of film composition, and consequently the structure for this material. Prior to each deposition, the target was sputtered for 30 minutes to obtain a steady material flow to the substrate.

The low substrate temperature provokes less deteriorating effects in the deposits, and consequently samples of prescribed compositions can be well reproduced. A low substrate temperature and higher working pressure can effectively suppress the formation of pure metal phase in the deposits. In order to avoid the destructive effect of the probing particles upon the samples, repeated measurements have been conducted on the equally fresh samples for the analysis of the unstable materials considered here.

With a fixed power supply of 150 W, the copper content in the deposits declines steadily with increasing nitrogen proportion, and at a nitrogen proportion of 90% it approaches the value 76.9% (Fig. 8.2). Further increase of nitrogen in the working gas failed to obtain an expected stoichiometric Cu₃N deposit, even with pure nitrogen at this power supply. This failure in obtaining a strictly stoichiometric copper nitride deposit has nothing to do with the substrate temperature because nitrogen re-emission from compound is unlikely to occur at 60 °C, but rather lies in the nature of magnetron sputtering. Since the reaction with nitrogen can only take place in the region of the plasma ring on the surface of the Cu target, it is the energetic sputtering there that prevents the formation of stoichiometric Cu₃N clusters. The unwanted presence of elemental copper precipitates in the deposits provides an explanation for the discrepancy or even controversy in many characterizations of the film properties, as will be made clear later.

X-ray diffraction confirmed the nanocrystallinity of the deposits containing Cu₃N and Cu biphases. With a Cu content below 78.8%, the reflections from the cubic Cu₃N phase become dominant, as can be seen from Fig. 8.3. The reflections assigned to the {001}- and {002}-planes of the Cu₃N phase remain unchanged, while the reflection from the {111}-plane of the Cu₃N phase shifts irregularly through the sequence (a)–(e). This is quite reasonable since the nitrogen re-emission most preferably occurs on the {111}-planes, thus their spacing is susceptible to the variation of the overall composition of the deposits. The lattice constant calculated from this reflection oscillates between 3.868 and 3.797 Å.

For Cu₃N, the {111}-planes consist exclusively of either Cu atoms or N atoms, hence nitrogen loss most preferably occurs on these planes, which sets a limit on the size of the Cu₃N crystallites (to be discussed below). At the grain boundaries, the Cu atoms agglomerate to form nanocrystals, generally being ~5 nm in size as revealed by transmission electron microscopy. The existence of juxtaposed Cu-rich {111}-planes of Cu₃N and the Cu nanocrystals implies an uncertain factor in determining the optical and, in particlar, electrical properties. By measuring the electrical resistivity, a sudden drop was observed when the content of Cu was over 78.8% (Fig. 8.4), suggesting that a new mechanism for the electrical conductance was now switched on. Writing the composition of this sample in the form Cu_15.2–Cu_63.6 N_21.2, and recalling that the excessive Cu atoms exist both as Cu nanocrystals and on the Cu-rich {111} planes of Cu₃N crystallites, an electrical conduction path solely via Cu chains can be established through the sample, i.e. the percolation mechanism comes into play. The resistivity for the sample with 76.9% Cu is 4.04 × 10⁻⁵ Ωm, just with a little bit more excess Cu it drops to 1.0 × 10⁻⁶ Ωm, falling already into the range for good conductors. The sensitivity of the electrical resistivity to the content of Cu explains the large inconsistency in the reported values of electrical resistivity for copper nitride deposits.

The films obtained above at 150 W seriously deviate from the stoichiometry for Cu₃N. In order to obtain a nearly stoichiometric sample, the combination of power supply and the pressure of working gas (now we employ pure nitrogen gas) should be tested. Figure 8.5(a) and (b) display the XRD patterns for the samples deposited with various RF powers, and under two working pressures of 0.7 Pa and 0.9 Pa for comparison. It can be seen that the films always display the (001) and (002) reflections, and the (111) reflection appears when the supplied power is over 100 W. The presence of pure Cu crystallites is disclosed by the minor Cu-(111) reflection at 2θ = 43.32°. Clearly, the higher pressure and the lower RF power favor the formation of deposits of pure Cu₃N phase in a competitive way. At 0.7 Pa, the Cu-(111) reflection is visible in the deposit prepared with an RF power of 100 W; at 0.9 Pa, however, it remains absent when the RF power was raised to 150 W. This is quite easy to understand since both the higher pressure and lower RF power can reduce the energy of nitrogen ions bombarding the Cu target. By reactive sputtering of Cu target with nitrogen plasma, a very large projectile energy (large power and low pressure) is unfavorable for the formation of Cu–N bonds in the sputtered particles. Generally speaking, to obtain nearly stoichiometrical copper nitrides at a preset low substrate temperature, the gas pressure and the power should be chosen in such a way that the density of nitrogen ions is sufficiently high and the ion energy has to be low so as not to spoil the Cu–N bonded species emitted from the bombarded target. A strict specification of the processing parameters demands a thoughtful mass- resolved analysis of the bombardment products. Remarkably, in addition to the [001]-orientation, the larger RF power also provokes the occurrence of the [111]-orientation for the Cu₃N crystallites. From Fig. 8.5 one can say with confidence that with a power supply at 70 or 50 W, [001] oriented thin films of copper nitride without any trace of the Cu(111) reflection on the XRD pattern can be obtained.

Copper nitride even in the pure Cu₃N phase still reveals a minor deficiency of nitrogen. This is to say that in the stoichiometrical samples we presented here, the Cu content is still somewhat larger than 75.0%. In fact, from the energy-dispersive X-ray analysis and photoelectron spectroscopic data, the lowest Cu content is about 75.6%, to the accuracy of the methods. The nearly stoichiometrical copper nitride sample should be a typical deficit semiconductor, as seen from the temperature dependence of the electrical resistivity (see Fig. 8.6(a)). A slight increase in the content of Cu to ~76.0% can totally destroy this behavior for a deficit semiconductor, as the resistivity first drops irregularly with the decreasing temperature, and below 50 K it turns to increase steeply (not shown). In the Cu-rich sample with 78.8% of Cu, due to the presence of distinct Cu particles (Du et al. 2005), it shows a characteristic metallic behavior that the electrical resistivity increases linearly with the temperature rising from 5 K to 300 K (Fig. 8.6(b)). According to the equation ρ (T) = ρ₀ [1 + α(T– T₀)], the temperature coefficient a is determined to be 7.54 × 10⁻⁴, much less than the value 1/232 for pure Cu. This can be explained by the fact that in the Cu-rich copper nitride, the conduction path is made of both Cu nanoparticles and the Cu₃N nanoparticles enclosed by the Cu-terminated {111}-planes. On the contrary, the stoichiometric Cu₃N sample manifests a typical semiconductor character, the temperature dependence of electrical resistivity (Fig. 8.6(b)) can be approximately fitted by the equation ρ (T) = C exp(− E_a/kT), with the activation energy E_a being approximately 8.59 × 10⁻² eV. For comparison, E_a is 1.47 × 10⁻² eV as reported by Wang et al. (1998).

The host material Cu₃N in the cubic anti-ReO₃ lattice is a typical semiconductor with an indirect band-gap of ~1.9 eV (Moreno-Armenta et al. 2004, Nosaka et al. 2001, Maruyama and Morishita 1996). Due to the weak Cu–N bonding that facilitates nitrogen re-emission from the compound, the copper nitride samples generally show some deficiency of nitrogen; consequently they display electron-like conductivity as confirmed by the negative Hall coefficient. By careful control of the processing parameters, nearly stoichiometric, single-phased Cu₃N thin films can be obtained with a carrier density brought down to below 3.0 × 10¹⁸/cm³.

With the stoichiometrical Cu₃N sample at hand, a high-resolution TEM image can be obtained, which requires enormous expertise since the energetic electron beam in a TEM can readily destroy the Cu₃N lattice. Figure 8.7 displays the first TEM image for the stoichiometric Cu₃N, exhibiting well-aligned {001} fringes, which confirms the high crystallographic quality of the deposit. The planar distance for the {001}-planes, i.e., the lattice constant, measures 0.383 nm in the TEM image, which is in good agreement with the XRD measurements.

Due to the nitrogen deficiency, the deposits consist of distinct nanosized particles instead of being a single crystalline as claimed by some authors. This can be judged easily from the line width of the XRD patterns. From the SEM images (Fig. 8.8) we see that the sample with 76.9% Cu is sharply contrasted, where the size of the crystallites ranges from 40 to 60 nm. For the sample with 75.6% Cu, the crystallites are roughly 40 nm in size, and the image seems smeared. This is obvious since the nearly stoichiometric sample is an insulator – it causes a serious charging effect under the illumination of the electron beam in an SEM.

With the nearly stoichiometrical samples, many inconsistencies concerning the physical properties of Cu₃N can be resolved. First, the pure, nearly stoichiometric Cu₃N deposit is a typical wide-gap semiconductor, whose electrical resistivity at room temperature measures 2 × 10⁻² Ωm. Smaller values for this quantity originate in the incorporation of Cu particles or Cu-rich {111} surfaces of the Cu₃N crystallites. It has a band-gap larger than 1.8 eV, as determined from the photoreflectance spectrum even for the slightly substoichiometric copper nitride thin film with 76.9% Cu (Du et al. 2005). Another important feature is the decomposition temperature. To study the thermal stability of the as-deposited films, the samples were annealed in the same chamber under the nitrogen environment at 0.9 Pa for 20 minutes. in a temperature range from 220 °C to 420 °C. The temperature at which Cu(111) reflection becomes noticeable on the XRD pattern of the post-annealed samples is taken as the decomposition temperature. Aannealing for 20 minutes. at 220 °C and 300 °C does not provoke any discernible changes to the diffraction pattern. Only at 350 °C do two tiny peaks emerge at 2θ ≈ 43.32° and 26 ≈ 50.55° (arrowed in Fig. 8.9) which can be assigned to the (111) and (002) reflections of the pure Cu phase (Fig. 8.9). The primarily absent Cu (111) and (002) reflections unambiguously signify the onset of thermal decomposition of Cu₃N. At 420 °C, only strong Cu peaks were observable, the Cu₃N completely dissolved. Although from the current experimental data the precise value of the decomposition temperature cannot be fixed, we can say with confidence that the Cu phases emerge from the decomposed stoichiometrical Cu₃N, revealed by XRD, only at a temperature of ~350 °C by annealing under the protection of nitrogen at a pressure of 0.9 Pa.

8.3 Nitrogen re-emission

The thermal instability of copper nitrides constitutes essential difficulties for their synthesis and structural characterization. By reactive magnetron sputtering deposition of copper nitride films, the effect of nitrogen re-emission on the film morphology may be directly observable, from which the effect on the growth process and on the microscopic structure of deposits can be inferred. The investigation of this topic is also helpful in understanding the failure mechanism of copper nitrides in usage, which may cause serious instability problems in practical applications.

To demonstrate the effect of nitrogen re-emission, copper nitride thin films were grown with mixed nitrogen (of 60% to 80% in flow rate) and argon as working gas, maintained at a pressure of 1.1–1.3 Pa. The power supply was set at 150 W. Under the given conditions, the deposits are substoichiometric with copper contents exclusively below 77.6%. Under SEM, the films grown for 15 minutes, which are roughly 0.3 μm thick, display a morphology with isolated round blisters (Fig. 8.10). These blisters are typically about 10 microns in dimension. The coverage by these protruding features is about 20%. In films deposited for 30 minutes, the protruding features turned into a ramified rosette-like structure. Such copper nitride rosettes are distributed homogeneously across the sample surface, and show a striking size and shape uniformity, with a lateral dimension around 23 microns (Fig. 8.10). By close inspection of the individual rosettes, it is observed that the rays bifurcate at an average angle of ~74°, slightly larger than the wedge angle of 71.53° for a tetrahedron. In some well-developed ones, fivefold symmetry can be confirmed at the center of the rosettes, though not as strict as in the geometrical sense.

At first glance, this morphology reminds us of island formation as frequently encountered in growing crystalline films, which can be generally modeled with an atomistic view. But we found that this is not the case. Both the smaller round caps and the larger rosettes are hollow, as confirmed by the presence of some cracked entities. When probing through the crack of a broken rosette, the energy-dispersive X-ray spectrometry detected a dominantly strong signal only for silicon and a weak signal for copper – the nitrogen is completely absent. This is to say that the protruding features are due to a complete local detachment of the film; and in that process some tiny amount of copper is left behind on the Si substrate (Ji et al. 2005, 2006a). Nevertheless, such protruding structures cannot have developed from a previously well-deposited coating due to the accumulated stress. This point can be clarified by briefly examining the necessary area expansion of the material for the formation of such hollow features. As a conservative estimation, consider a round cap 20 μm in lateral dimension and 1 μm high. This corresponds to an area expansion of ~5.0% with regard to the underlying circular base. Such a large area expansion coefficient is prohibitive for rigid materials such as the ionic copper nitride. Moreover, such a morphology is formed in the growth stage, post-growth degradation of an originally flat film due to nitrogen re-emission leaves behind a flat, but Cu-rich surface. The aging of films with a compact morphology in the ambient will not invoke any rosette structure.

The puzzle of the enormously expanded area of the rosettes with regard to the underlying base becomes immediately resolved with the aid of scanning electron micrographs at an enlarged magnification. We see that the deposits with rosette structures are composed of distinct crystallites, around 45 nm in dimension, just like a compact film, but the SEM images taken directly on a rosette reveals the peculiarity in the manner of crystallite stacking. The surface of the rosette structures displays ragged steps and terraces, and it is thinner than the flat portion of the deposits. The typical width of a terrace is about 100 nm (Fig. 8.11). Such a morphology has never been reported, to the best knowledge of the author. We are inspired to speculate that the crystallites in a rosette have experienced a rearrangement process. The area expansion in due course of relief formation is sustained through the gliding of nanocrystals, which is facilitated when the nanocrystals are clothed with the amorphized Cu-terminated {111}-planes; and the ongoing fast growth prevents the film from cracking.

In order to fully understand the formation mechanism for such protruding features, we take a close look at the morphological profile of the rosettes. They have the shape of a starfish with rays, but these rays generally do not meet at the same point in the central disk. Occasionally, a symmetrically developed rosette could be found, in which a perfect pentagram is discernible at the center, as shown in Fig. 8.12(a). The rays show further bifurcations, and the offshoots extend generally at an angle of about 74°, a little larger than 72° as required by fivefold rotational symmetry, which in turn is a little larger than the wedge angle of a tetrahedron (α = 70.53°). This cannot be a coincidence, bearing in mind that the tetrahedral crystallites do have the tendency to juxtapose at an angle of about 72° with a [110] orientation in order to fill the space, giving rise to a loosely fivefold symmetry. Fivefold symmetry evolving from this construction mechanism has been confirmed in nanostructures of cubic metals such as Au, Ag and Cu.

As we know, the {111}-planes of the cubic Cu₃N lattice comprise exclusively either Cu atoms or N atoms. When nitrogen re-emits from a Cu₃N crystallite, it leaves behind crystallites with a soft clothing of amorphous Cu layer in a habit, not necessarily tetrahedral, but with facets joining at the wedge angle (α ≈ 70.53°). The clothing of Cu₃N crystallites by Cu {111}-planes also builds a particular electrical conductance path, as verified from the percolation transition in the variation of its electrical resistivity versus the Cu content in the film. When driven by a sufficiently large stress, here initiated by the re-emitted nitrogen gas from beneath, the nanocrystals will undergo a gliding motion against each other along the {111}-planes, and conglomeration of a large quantity of such nanocrystals with facets joining at α ≈ 70.35° results in fivefold symmetry at larger scales. Displayed in Fig. 8.12(b) is a piece of Cu₃N crystal in a specimen prepared with mechanical milling for transmission electron microscopic investigation – it is pentagonal with an edge length of 660 nm.

Now the picture becomes clear. An ‘orogenic movement’ model can be formulated to illustrate the formation of relief rosettes in a growing Cu₃N thin film (Fig. 8.5). Copper nitride keeps decomposing while film growth proceeds. The re-emitted nitrogen atoms recombine into molecules and together with entrapped nitrogen they agglomerate to form gas bubbles at the film–substrate interface by virtue of a large lateral mobility, since solubility of nitrogen in both Si and Cu₃N is low. Adhesion to the local substrate is undermined where a gas bubble forms, and finally detachment occurs (Fig. 8.5(b)). By a delicate balance between expansion via nanocrystal glide and fast growth, the total area and the lateral dimension of the delaminated part grow in pace without cracking. As can be expected, at reduced growth rates a compact morphology of the film ensues. Gliding motion of the crystallites into ragged steps and terraces – similar to the movement of rocks in mountain building, i.e. the orogenic movement – provides the path for a complete strain relaxation; this explains the negligible residual stress in the sample despite the presence of the relief features, as concluded from the transmission electron microscopic and X-ray diffraction data about the lattice constant (Fig. 8.1). The possibility to glide along the Cu-terminated {111}-planes for the pyramidal Cu₃N crystallites is prerequisite for the formation of the relief morphology; this also explains its absence in the films of nearly stoichiometrical or heavily Cu-rich samples. As to the quintilobed structure of the rosettes that ramify at a branching-off angle of ~74°, they are the consequence of the preferred packing and rearrangement driven by an anisotropic stress of about two million {111}-faceted Cu₃N nanocrystals (Ji et al. 2007). Unlike in the nanorods of cubic metals where the core is fivefold twin crystals textured along the [110]-axis, the fivefold symmetry here is quite ‘soft’ in the sense that it arises rather as an integral effect at large scales.

8.4 Doping of Cu₃N by co-sputtering

The structure of Cu₃N is, in itself, rather interesting. It is an open structure that can, theoretically, accommodate one more metal ion in each unit cell. The incorporation of an excess metal atom can significantly alter the band structure of the material that continuous semiconducting-to-(semi)metallic transition can be expected, through which various interesting phenomena can be measured in the ternary compounds (Lovett 1977). Here, through the co-sputtering of Cu and Pd targets, we demonstrate the deposition of ternary Cu₃NPd_x thin films, where a vanishing temperature coefficient of resistivity (TCR) in the range from 240 K down to 5 K could be measured in Cu₃NPd_0.238 (Ji et al. 2006b). Such a phenomenon has not been found in any other single solids.

The inclusion of another metal atom in each unit cell of the Cu₃N lattice results in compounds like Cu₄N or the ternary Cu₃NPd (Hahn and Weber 1996). Remarkably, Cu₃NPd is a semimetal due to the intersection of the energy bands at the Fermi level. Consequently, a semiconducting-to-semimetallic transition is anticipated in the off-stoichiometric Cu₃NPd_x when ‘x’ increases continuously from zero to unity.

We prepared thin films of Cu₃N and Cu₃NPd_x (0 < x < 0.350) on the Si (001) wafers, a low substrate temperature (~100 °C) in combination with a small power supply (< 150 W) and a high working pressure (> 0.7 Pa) is applied. The interposition of metal atoms at the cell centers of the cubic Cu₃N lattice severely deteriorates the covalent bonding in the resulting materials such that the compounds Cu₄N and Cu₃NPd only exist in theory. In fact, the largest reported value of ‘x’ never exceeds 0.989 in Cu₃NPd_x whereby the compounds exist in fine powders (Jacobs and Zachwieja 1991). We found that those Cu₃NPd_x thin film samples with x > 0.5 show a very poor crystallinity; therefore we restrict our discussion to the samples with 0.0 ≤ x ≤ 0.350 which are well crystallized as confirmed by both X-ray diffraction pattern (Fig. 8.13) and transmission electron micrograph (Fig. 8.14). For samples containing only a tiny amount of Pd, e.g., Cu₃NPd_0.048, the X-ray diffraction pattern remains unchanged, like that of the host Cu₃N lattice showing only the (001) and (002) reflections. When the Pd concentration is large enough (x ≥ 0.175), the (001) reflection was replaced by the strong (111) reflection which deteriorates both in intensity and in spectral profile. We can say that the Cu₃NPd_x thin films here concerned with x < 0.35 are crystalline. The addition of Pd atoms results in a slightly enlarged lattice constant which increases from a = 0.383 nm for the pure Cu₃N to a = 0.385 nm for Cu₃NPd_0.175 (Fig. 8.14). The Cu₃NPd_x samples obviously have a better electrical conductivity which in turn results in an improved image quality for the transmission electron micrograph (cf. Fig. 8.7). The fact that the Cu₃NPd_x is metallic at sufficiently large ‘x’ values, while Cu₃N is semiconducting, may suggest that the interposition of Pd atoms alter the energy bands near the Fermi surface.

Figure 8.15 displays the temperature dependence of the electrical resisitivty for the nearly stoichiometric Cu₃N and for the different ternary Cu₃NPd_x compounds. Upon cooling from room temperature, the electrical resistivity in the intrinsic Cu₃N film rises rapidly (Fig. 8.15(a)). This observation is consistent with the wide band-gap nature of this semiconducting material. The insertion of Pd atoms confers Cu₃NPd_x an improved electrical conductivity owing to the narrowed band-gap. Even with a tiny amount of Pd insertion as in the samples with x = 0.048 and x = 0.071, the rapid onset of resistivity upon cooling is shifted to below 50 K, indicating a narrowed band-gap for the current samples (Fig. 8.15(b) and (c)). A further addition of Pd atoms initiates the semiconducting-to-semimetallic transition such that the TCR changes its sign. For Cu₃NPd_0.175, its electrical resistivity rises roughly linearly with decreasing temperature in a range as wide as 270 K, and the TCR is –0.00039 K⁻¹ (Fig. 8.15(d)). Taking T = 280 K as the reference point, the resistance over such a broad temperature range changes only by ~5.0%. The resistivity for Cu₃NPd_0.349 also displays a good linearity from 50 K to 280 K, but now the TCR is 0.00117 K⁻¹ (Fig. 8.15(f)), indicating unambiguously that now the compound has changed to be metallic. In between these two cases of opposite tendencies for the variation of the electrical resistivity with temperature, we may expect a vanishing TCR over a considerable temperature range at some ‘x’ value. This occurs in the sample Cu₃NPd_0.238. From 240 K down to 5 K, the TCR is nearly zero (< 3.0 × 10⁻⁶/K), and the relative variance of the electrical resistivity in this temperature range measures as small as 1.42 × 10⁻⁵ (Fig. 8.15(e)).

The temperature dependence of the electrical conductivity for solids is susceptible to many influencing factors; no hard-and-fast rules are available. For a semiconductor, the electrical conductivity can be formally written as σ = e(n_eμ_e + n_hμ_h), where e and h denote electron and hole, while n and μ stand for the carrier density and carrier mobility, respectively. Both the carrier mobility and the carrier density, or the number of the carriers, are temperature dependent, yet the temperature dependence of the electrical conductivity is dominated by that of the latter, since the mobility changes only in a mild way with the lattice temperature while the number of carriers can be an exponential-like function of 1/T. With decreasing temperature, the number of carriers decreases while the mobility generally increases, but in different manners; consequently, the resulting electrical resistivity is exclusively temperature dependent. For narrow band-gap semiconductors and semimetals, the situation is more complicated than in the conventional semiconductors like Si or Ge. The carrier distribution function obeys the Fermi–Dirac statistics which in this case cannot be approximated by the Maxwell–Boltzmann formula. Moreover, the conduction band is Kane-type that the effective mass of electron, thus the mobility of electron, changes with the carrier density (Lovett 1977). The vanishing TCR occurs when this compound has narrowed its band-gap to be a semimetal for which the temperature dependence of the number of the carriers becomes rather mild-mannered in contrast to the situation in a semiconductor of a definite bandgap. It results from a delicate balance between the opposite changes of the number of carriers and of the carrier mobility with temperature, and reasonably it appears only at temperatures below 240 K – at higher temperatures the increasing number of carriers can easily compensate for the loss in carrier mobility to give rise to a rapidly decreasing electrical resistivity. The current discovery, though lacking a clear scenario of detailed microscopic processes due to the difficulty in calculating the band structure for the off-stoichiometric compounds such as Cu₃NPd_0.238 and in performing Hall measurement for some supplementary information, does demonstrate the possibility that a balanced change of carrier density and carrier mobility over a wide temperature range is in principle possible in a class of semimetallic materials. Considering the unlikelihood of a temperature-independent electrical resistivity for solids, the current result also bears some significance for solid state physics.

8.5 Conclusions

By carefully adjusting the working parameters including the applied power, gas pressure and the ratio of nitrogen in the working gas, and substrate temperature, thin films of noble-metal nitrides such as Cu₃N and Cu₃NPd_x can be successfully prepared by using the reactive magnetron sputtering method. For stoichiometric Cu₃N, the films are found to comprise distinct crystallites in a typical size of 40–60 nm. The peculiarity of growing noble-metal nitride films lies in the fact that the reaction of sputtered metal atoms or clusters with nitrogen occurs most probably only at the target surface, hence the matching of the density and energy of nitrogen ions with the sputtering yield for a chosen target material is the key factor to realizing stoichiometry in the resulting copper nitride films. A large ion energy is unfavorable since the higher sputtering yield of the metal atoms and the weakened combination with energetic nitrogen ions will give rise to a metal-rich deposit. A low substrate temperature is critical since at slightly elevated temperatures nitrogen re-emission occurs, which may even result in microsized protruding features. Semiconducting to (semi)metallic transition can be observed in doped Cu₃N structures wherein the excessive atoms are expected to sit at the cell centers of the Cu₃N lattice. A vanishing temperature coefficient of resistivity over a temperature of 200 K has been realized in Cu₃NPd_0.238. Many other surprising properties are anticipated in the noble-metal nitrides, which await discovery in the future.

8.6 References