The contact angle measurements were accustomed to analyzing the surface wettability of as-prepared filter membranes, as illustrated in Figure 5. Additionally, the wettability of the studied surface layers was evaluated using MF-0.05 as a test sample. A droplet of water was dropped onto a ZnO-coated filter membrane in the air. It diffused rapidly within 80 ms, which suggested that the coated filter membrane had satisfactory hydrophilicity (Figure 5a–c). Furthermore, the oil droplet swiftly spread on the as-prepared filter membrane in the air, demonstrating a superoleophilicity (Figure 5e–g). However, the oil droplet could not spread on the as-prepared filter membrane in the water, illustrating that the coated filter membrane had an excellent oil repellency once submerged (Figure 5d). A drop of water was also unable to spread on the filter membrane previously prepared, which demonstrated the filter membrane's outstanding repellence towards water underoil (Figure 5h). ZnO coated filter membrane had a static underwater oil CA of 151.4◦ ± 0.8◦ and underoil water CA of 152.7◦ ± 0.6◦ , respectively, exhibiting both superhydrophobicity in oil and superoleophobic properties in water.

#### *3.3. Separation of Oil-in-Water Emulsions*

In contrast to a stratified oil–water mixture, an oil-in-water emulsion containing emulsified oil droplets with various sizes was difficult to be separated. To prevent emulsified droplet penetration and achieve separation of O/W emulsions, the micropore size was one of the important characteristics of separation membranes. Here, the capacity of ZnO coated filter membrane to separate in the O/W emulsion was demonstrated. To analyze the separation efficacy of an as-prepared filter membrane, emulsions of diesel/water, N-decane/water, dodecane/water, and N-tetradecane/water were invented. Figure 2

illustrated that the apparatus utilized separate O/W emulsions. The cleaned water was gravity-fed through the membrane and collected in the beaker underneath, while the as-prepared emulsions were poured into the top tube. The electronic balance recorded the mass of the collected filtrate in real-time. The separation findings of dodecane/water emulsion were illustrated in Figure 6. As shown in the microscopic imaging, the prepared water/dodecane emulsion contained a large amount of micron-sized oil droplets with an average size distribution of 8 µm. The collected water was transparent after filtration across the as-prepared filter membrane and there were no oil droplets revealed when observed under an inverted microscope.

**Figure 6.** The microscopic and optical images of O/W emulsions (**a**) before separation and average size of oil droplets; (**b**) filtered with filter membrane coated with ZnO.

Furthermore, other emulsions such as diesel/water, N-decane/water, and N-tetradecane/ water were successfully separated (Figure 7a,b). As expected, the as-prepared filter membrane successfully separated the four emulsions, but only left minimal residual oil in the filtrate. All the oil-in-water emulsions exhibited the promising permeation fluxes. The highest flux of 3139 L·m−2 ·h −1 was obtained for diesel-in-water emulsion. These findings suggested that the ZnO-coated filter membrane could separate a variety of O/W emulsions. Though the separation efficiency obtained in this study was similar with previous studies using inorganic coating materials (Nano-ZnO and Nano-TiO2) to separate oil-in-water emulsions, we obtained a higher separation flux performance than these studies [41–43]. In addition, the as-prepared filter membrane was evaluated by a ten-times recycling separation of the dodecane/water emulsion to indicate the characteristic of reusability. The membrane was ultrasonic-assisted cleaned with absolute ethyl alcohol for 5 min after each cycle of the experiment (collected filtrate volume was 80 mL), and it was immediately used for the following filtration test. After 10 recycles, the TOC concentration was less than 72 mg/L and the separation efficiency remained higher than 98%, confirming its capability for standing oil–water separation application (Figure 7c). Moreover, our results demonstrated an excellent cyclic stability of the filter membrane coated with ZnO, although the variations of flux in each cycle decreased with the increase in filtration duration (Figure 7d).

**Figure 7.** (**a**) The separation efficiency and TOC in the filtrates of oil-in-water emulsions; (**b**) various oil-in-water emulsion separation fluxes; (**c**) the recycling separation of the as-prepared filter membrane for O/W emulsions; (**d**) separation flux stability in 10-cycle experiments.

#### *3.4. Separation of Water-in-Oil Emulsions*

The ZnO-coated filter membrane could separate O/W emulsion and O/W emulsion. The water content in the filtrate, separation flux, and corresponding separation efficiency were used to analyze the separation performance of the as-prepared filter membrane. As demonstrated in Figure 8a, the appearance of the W/O emulsion was cloudy and contained sufficient micron-sized water droplets with an approximate average size of 10 µm. There were no water droplets in the liquid filtered by the ZnO-coated filter membrane (Figure 8b).

Figure 9a demonstrated that the as-prepared filter membrane was capable of accurately separating various W/O emulsions, such as water/diesel, water/N-decane, water/dodecane, and water/N-tetradecane. The efficiency of separation of the as-prepared filter membrane for various water-in-oil emulsions were more than 98%. Furthermore, the water concentration for the filtrates of different emulsions was less than 80 ppm. The following Equation (2) was used to estimate the filtration flux of the as-prepared filter membrane, and fluxes (L) for different emulsions were evaluated for each minute. The separation fluxes

for water/diesel, water/N-decane, water/dodecane, and water/N-tetradecane emulsions were about 392, 663, 960, and 744 L·m−2 ·h −1 , respectively (Figure 9b). The separation efficiency obtained in this study was similar to existing studies about inorganic coating materials, which were used to separate water-in-oil emulsions. However, the separation flux obtained here was lower than in all these studies [44,45].

**Figure 8.** The microscopic and optical images of surfactant-stabilized water-in-oil emulsions (**a**) before separation average size of water droplets; (**b**) filtered with filter membrane coated with ZnO.

**Figure 9.** The separation efficiency and water content in (**a**) water-in-oil emulsions filtrates; (**b**) filtration fluxes of various water-in-oil emulsions; (**c**,**d**) the recycling separation performance of ZnO coated filter membrane for water-in-oil emulsions.

As an example, the reusable separation capability of the as-prepared filter membrane was accessed by the separation experiment of water/N-decane. The separation of each cycle lasted for 8 min, followed by a 5 min ultrasonically cleaned with absolute ethyl alcohol. Figure 9c showed that even after 10 cycles of use, the efficiency of separation of the as-prepared filter membrane was always greater than 98%, and the water content in the filtrates was lower than 80 ppm, suggesting that the as-prepared filter membrane was more recyclable. Furthermore, during the reusability experiment, the filtration flux did not change appreciably, suggesting the filter membrane had excellent separation efficiency and splendid stability (Figure 9d).

#### *3.5. Demulsification Mechanism in Oil-in-Water and Water-in-Oil Emulsion*

ZnO nanostructures had achieved both a superhydrophobicity in oil and a superoleophobic property in water. As shown in Figure 10a–c, the separation mechanism could be concluded as follows: it was due to the excellently amphiphilic property in air; the as-prepared filter membrane was extensively wetted by the water phase when it interacted with an oil-in-water emulsion, allowing the water phase to easily infiltrate through. The emulsified oil droplets were then intercepted and adsorbed on the filter membrane in the second step. Finally, oil droplets attracted and coalesced with other emulsified droplets, allowing oil-in-water emulsions to be separated efficiently. Moreover, these water droplets captured by the membrane during the separation process were rapidly coalesced and demulsified to create bigger droplets (Figure 10e–g).

**Figure 10.** Mechanism of demulsification process of oil-in-water and water-in-oil emulsions; (**a**–**c**) coalescence and demulsification of emulsified oil droplets; (**e**–**g**) coalescence and demulsification of emulsified water droplets; (**d**,**h**) suppositional separation process of oil-in-water and water-in-oil emulsions.

Despite the above results and analysis, the flux change in separation process deserves further study. The oil droplets constantly agglomerated and adsorbed to the fibers during oil-in-water emulsion separation. As illustrated in Figure 11a,b, the oil phase was intercepted to access the filter membrane's surface. The oil–water flux through the membrane was determined by membrane porosity, pore size, as well as pore distribution. Current study found that relatively large pore size leads to a high separation flux. Some parts of porous structures would be blocked by the absorbed oil droplets on the membrane, which resulted in a decreased porosity and a shortened effective separation area of the membrane. In other words, the oil–water separation was saturated with the duration, which significantly decreased the oil–water separation flux. As envisioned, our results suggest that the variations of permeation flux decreased with the increase in filtration duration. The flux decline was in the range of 45–55% (Figure 11c). The same trend was observed during the separation of a water-in-oil emulsion (Figure 11d–f). In addition, the fluxes in separation process varied with the time (Figure 11c,f). As the accumulation of droplets increased on

the surface, the pore size decreased and eventually changed the influx during the emulsion separation process. Our observations were consistent with the previous studies with similar focuses. For example, the separation flux of membranes with different pore sizes was reported at a range between 250 and 4300 L·m−2 ·h −1 [15,41]. We need to point out that the data obtained from a single moment may not accurately describe the flux change in the separation process. It will be more practical if we focus on a general description of the entire separation process or the average data shown in Figures 7d and 9d. Finally, it was needed to point out that the filter membrane could continue to adsorb emulsified droplets even after a simple cleaning.

**Figure 11.** Image of metal fiber felt surface in the separation process of (**a**) oil-in-water emulsions and (**d**) water-in-oil emulsions; (**b**,**e**) the performance oil–water separation through one hour; (**c**,**f**) variation of separation fluxes with each minute.

#### **4. Conclusions**

Herein, the surfaces with special wettability were achieved by ZnO nanostructures, which could improve oil/water separation efficiency. In this study, ZnO nanostructures were fabricated by a two-step process for the sol-gel process and crystal growth from the liquid phase to achieve both a superhydrophobicity in oil and a superoleophobic property in water. Our study demonstrated an excellent wettability of nanostructured stainless-steel fiber felt. The ZnO coated filter membrane had a static underwater oil contact angle of 151.4◦ ± 0.8◦ and underoil water contact angle of 152.7◦ ± 0.6◦ , respectively.

In oil–water separation, the emulsified impurities in both water-in-oil and oil-in-water emulsion liquid were effectively intercepted to achieve an oil–water separation. The filter material with small pores (~5 µm diameter) could separate diverse water-in-oil and oilin-water emulsions with a separation efficiency higher than 98%. Based on the oil–water separation mechanism, we could predict that the blocking of filter materials was caused by the collection of emulsified impurities. Thus, the effect of filtering quantity on the separation efficacy was investigated to analyze the influence of impurities. The investigation results showed that the filtration flux decreased with the collection of emulsified impurities. However, the filtration flux could restore after cleaning and drying, suggesting the recyclable nature of our strategy. The nanostructured filter material is a promising candidate for separating both water-in-oil and oil-in-water emulsion in industry. The proposed filter materials in this study may be repetitively used without any treatments, secondary pollution, and low energy consumption. Therefore, our methods may have wide application in the future.

**Author Contributions:** Conceptualization, X.L., D.Y. and B.W.; methodology, X.L.; validation, X.L., S.F., C.W. and L.C.; writing—original draft preparation, X.L. and S.F.; writing—review and editing, D.Y., L.C. and B.W.; supervision, D.Y. and B.W. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by National Key R&D Program, grant number 2019YFC1907602; National Natural Science Foundation of China, grant number 52076036; Talent introduction and scientific research start-up funded project of Northeast Petroleum University, grant numbers 2019KQ and<PHONE_NUMBER>; and Northeast Petroleum University Youth Science Foundation of China, grant number 2019YDQ-01.

**Data Availability Statement:** The data presented in this study that support the findings are available on reasonable request from the corresponding author.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **Effect of Cu2O Substrate on Photoinduced Hydrophilicity of TiO2 and ZnO Nanocoatings**

**Maria V. Maevskaya 1 , Aida V. Rudakova 1 , Alexei V. Emeline 1 and Detlef W. Bahnemann 1,2,\***

**Abstract:** The effect of a Cu2O substrate on the photoinduced alteration of the hydrophilicity of TiO2 and ZnO surfaces was studied. It was demonstrated that the formation of heterostructures Cu2O/TiO2 and Cu2O/ZnO strongly changed the direction of the photoinduced alteration of surface hydrophilicity: while both TiO2 and ZnO demonstrate surface transition to superhydrophilic state under UV irradiation and no significant alteration of the surface hydrophilicity under visible light irradiation, the formation of Cu2O/TiO2 and Cu2O/ZnO heterostructures resulted in photoinduced decay of the surface hydrophilicity caused by both UV and visible light irradiation. All observed photoinduced changes of the surface hydrophilicity were compared and analyzed in terms of photoinduced alteration of the surface free energy and its polar and dispersive components. Alteration of the photoinduced hydrophilic behavior of TiO2 and ZnO surfaces caused by formation of the corresponding heterostructures with Cu2O are explained within the mechanism of electron transfer and increasing of the electron concentration on the TiO2 and ZnO surfaces.

**Keywords:** photoinduced hydrophilicity; surface energy; heterostructures; charge transfer; work function; adsorbed water

#### **1. Introduction**

Surface wettability is an important property of modern functional materials [1,2]. The self-cleaning, anti-fogging and anti-corrosive action of photoactive coatings and photocatalytic materials is based on a light-controlled surface hydrophilicity effect [3]. The photoinduced alteration of the surface wettability possesses a number of advantages compared to other methods based on the application of electric potential, mechanical stress, thermal or chemical action on the surface, etc. Photostimulated alteration of the surface hydrophilicity is easily controllable, environmentally friendly, energetically beneficial, safe and non-destructive.

Since the discovery of the effect of photoinduced superhydrophilicity for the titanium dioxide surface [4], fundamental studies of the mechanism and key factors of the process of the photoinduced alteration of the surface hydrophilicity are still in progress. Nowadays, it is established that electronic photoexcitation of the coating material plays the key role in this surface process, which is confirmed by numerous experimental data obtained by different research groups [4–11]. Particularly, it is assumed that the trapping of photogenerated charge carriers by active surface sites leads to further rearrangement of the hydroxyl-hydrated layer. As a result of such reconstruction, the surface free energy (SFE) changes. Then, the hydrophilicity of the surface changes accordingly: the higher the SFE is, the more hydrophilic the surface becomes.

Recently, we proposed a mechanism of photoinduced hydrophilic conversion, which includes elementary steps associated with the photoactivation and photodeactivation of

**Citation:** Maevskaya, M.V.; Rudakova, A.V.; Emeline, A.V.; Bahnemann, D.W. Effect of Cu2O Substrate on Photoinduced Hydrophilicity of TiO2 and ZnO Nanocoatings. *Nanomaterials* **2021**, *11*, 1526. https://doi.org/10.3390/ nano11061526

Academic Editor: Yuanbing Mao

Received: 6 May 2021 Accepted: 31 May 2021 Published: 9 June 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

the active surface sites responsible for the transition of the surface from one hydrophilic state to another [5,6]. It is represented as follows:

$$\mathbf{S} + \mathbf{h}(\mathbf{e}) \to \mathbf{S}^\* \tag{1}$$

$$\mathbf{S}^\* + \mathbf{e}(\mathbf{h}) \to \mathbf{S} \tag{2}$$

$$\mathbf{S}^\* \to \mathbf{H} \tag{3}$$

$$\text{H} + \text{e(h)} \rightarrow \text{S} \tag{4}$$

$$\mathbf{H} \to \mathbf{S} \tag{5}$$

$$k\_{\mathbf{5}\prime} \tag{6}$$

and a solution of the corresponding set of differential kinetic Equations (1)–(5) gives the expression for the rate of the photoinduced hydrophilic conversion (Equation (6)):

$$
\Delta \mathbf{H}(\mathbf{t}) = \mathbf{H}(\mathbf{t}) - \mathbf{H}\_0 = \left(\frac{\mathbf{A}}{\mathbf{B}} \mathbf{S}\_0 - \mathbf{H}\_0\right) \left(1 - \mathbf{e}^{-\left(\frac{\mathbf{C}}{\mathbf{B}}\right)\mathbf{t}}\right) \tag{6}
$$

where A = *k*1·*k*3·*n*1, B = *k*2·*n*2 + *k*3 and C = *k*1·*k*3·*n*1 + *k*2·*k*4·*n*2 2 − *k*3·*k*4·*n*2; *k*i is the rate constant of the i-th stage; S0 is the initial concentration of the surface sites (S) acting as either hole or electron trap; and *n*1 and *n*2 are the surface concentrations of the photocarriers (electrons and holes), participating in activation (1) and deactivation (2) of the active surface sites.

The direction of the photoinduced alteration of the surface hydrophilicity is determined by the sign of the ∆H(t) value: for a positive value, the surface becomes more hydrophilic, and, for a negative value, it becomes more hydrophobic. Analyzing the kinetic Equation (6), one can see that the rate and direction of the process on the surface of the same material is dictated by the ratio of the concentration of photocarriers of opposite sign (*n*1*/n*2) [5,6]. Upon photoexcitation of a solid, the ratio between concentrations of electrons and holes can be changed by two means.

The first method is achieved by varying the spectral composition of the acting light, namely by the photoexcitation of the material in the region of its intrinsic or extrinsic absorption. The change in surface wettability with a change in the spectral composition of the irradiating light has been demonstrated for various coatings [5,7,9,12–15].

The second way to change the *n*1*/n*2 ratio is to create a layered heterostructured coating, the components of which form a type II heterostructure. In the literature on selfcleaning materials, the formation of composite, or heterostructured, coatings is primarily mentioned as an effective method for improving their self-cleaning properties due to the enhanced photocatalytic oxidative ability [3,16–18]. It is well known that the formation of the type II heterojunctions promotes the charge separation, which reduces electron–hole recombination, thus improving the photocatalytic efficiency of such systems compared to single photocatalysts [3,15,19–21]. At the same time, it has been shown that this approach is also promising and productive both for fundamental studies of photoinduced superhydrophilicity of the surface and for application of self-cleaning coatings with controlled wettability [5,6,22–24].

In this study, the effect of the Cu2O substrate on the photoinduced hydrophilic behavior of TiO2 and ZnO thin films was studied and interpreted using the proposed approach (Equations (1)–(6)). For this purpose, the "layer-by-layer" thin films of TiO2/Cu2O and ZnO/Cu2O, with TiO2 and ZnO layers on the top of the sandwich-like heterostructures, were synthesized. These pairs of photoactive metal oxides were chosen due to suitable relative alignment of their conduction and valence band edges to form type II, staggered heterojunctions: the conduction-band edge and the valence-band edge of cuprous oxide are higher in energy than the corresponding band edges of either titanium dioxide or zinc oxide. As a result, for both TiO2/Cu2O and ZnO/Cu2O heterostructures, electrons are confined in the titanium and zinc oxides, while holes can be accumulated in Cu2O. Due to such an

effective charge separation, the considered composite materials have already demonstrated their high efficiency in both photoelectrochemical and photocatalytic processes [25–37].

In this study, we demonstrated that the formation of TiO2/Cu2O and ZnO/Cu2O heterojunctions drastically changes the surface hydrophilic properties of TiO2 and ZnO surfaces under UV irradiation. In addition, it was shown that selective photoexcitation of Cu2O with visible light also affects the surface hydrophilicity of the TiO2 and ZnO components due to electron transfer in heterostructures.

#### **2. Materials and Methods**

The individual ZnO and Cu2O films were formed by a sol-gel dip coating method (KSV Nima dip coater, Espoo, Finland) on SiO2-coated glass substrates to prevent the diffusion of sodium ions from the glass during thermal treatments. The velocity of withdrawing from solution was 100 mm/min for all coatings.

For the copper oxide sol, 15 mL of diethanolamine (99%, Fluka, Seelze, Germany), used as a stabilizer, were intensively stirred in 150 mL of isopropanol (99.8%, Ecos-1, Moscow, Russia), and then 14.9 g of copper (II) acetate hydrate Cu(OAc)2·H2O (99.0%, Vekton, Saint-Petersburg, Russia) were added at room temperature. A dark blue solution with a concentration of 0.5 M was kept for 24 h before further procedure [38]. To form Cu2O films, the obtained Cu(II)-containing layers were annealed at 300 ◦C on a hot plate in a nitrogen atmosphere for 40 min.

For the zinc oxide sol, 2.5 mL of ethylene glycol (99.5%, LenReactiv, Saint-Petersburg, Russia) and 10 g of zinc acetate dihydrate Zn(OAc)2·2H2O (99.0%, Vekton, Saint-Petersburg, Russia) were mixed in a round-bottomed flask and heated at 100 ◦C for 15 min to obtain a uniform transparent mixture. After cooling down to room temperature, 150 mL of isopropanol (99.8%, Ecos-1, Moscow, Russia) and 6.4 mL of triethylamine (99.5%, PanReac AppliChem, Darmstadt, Germany) were added to the mixture to promote the hydrolysis of the zinc acetate. Then, 0.5 mL of glycerin (99.5%, LenReactiv, Saint-Petersburg, Russia) were added dropwise to improve the film quality. The obtained clear and homogeneous solution previously stirred at 60 ◦C for 1 h was aged for 24 h at room temperature [39]. The formed ZnO layers were annealed at 280 ◦C in ambient atmosphere for 60 min.

The TiO2 thin film was formed by the atomic layer deposition (ALD) method ("Nanosurf" installation produced by "Nanoengineering Ltd.", Saint-Petersburg, Russia, RC "Centre for Innovative Technologies of Composite Nanomaterials", Research Park, Saint-Petersburg State University) on SiO2-coated glass substrate. Titanium tetrachloride (CAS №7550-45-0, quality level MQ200, Merck, Darmstadt, Germany) and deionized water were used as titanium precursor and hydrolysis agent, respectively. The deposition was carried out on the substrate at 200 ◦C.

The "layer-by-layer" TiO2/Cu2O and ZnO/Cu2O systems were formed by deposition of TiO2 and ZnO layers, respectively, on a thin Cu2O film formed as described above.

The surface morphology and film thickness of all synthesized coatings were explored by scanning electron microscopy (Zeiss Supra 40 VP system, Oberkochen, Germany). The smoothness of the film surface was assessed by the AFM method. X-ray diffraction measurements with Bruker "D8 DISCOVER" high-resolution diffractometer (CuKa X-ray radiation, within the angle range of 20◦ ≤ 2θ ≤ 80◦ with a scanning speed 5.0◦/min, Germany) were used for the crystal phase determination. Structural reference data were taken from the ICSD database. The transmittance spectra were recorded in the 250–800 nm spectral range at ambient conditions using Lambda 650S spectrophotometer (PerkinElmer, Inc., Shelton, CT, USA). The XPS spectra were recorded using a Thermo Fisher Scientific Escalab 250Xi spectrometer (Thermo Scientific™, Waltham, MA, USA).

Work function measurements were performed with a scanning Kelvin probe system SKP5050 (KP Technology, Wick, Scotland) versus a golden reference probe electrode (probe area 2 mm2 ). The probe oscillation frequency was 74 Hz, and the back potential was 7000 mV. Work function values were obtained by averaging 50 data points for four different sites of each sample. Estimated experimental error does not exceed ±0.06 eV.

The contact angle values were measured using optical tensiometer (Bioline Theta Lite, Biolin Scientific, Gothenburg, Sweden). The surface energy was calculated by the Owens–Wendt–Rabel–Kaelble (OWRK)/Fowkes approach using the two-liquid method (water contact angle versus methylene iodide contact angle) [40]. Ultrapure water has initial pH of 5.5. An experimental error of contact angle measurements was determined using 5 data points measured at different spots of the coating and did not exceed 2◦ .

The work function and contact angle were measured after each step of the surface treatment procedure. After annealing at 200 ◦C for 30 min, the state of the film surface is designated as "as prepared". After wetting in the ultrapure water with pH of 5.5 and drying at 80 ◦C, the surface state is denoted as either "after wetting" or "initial state". The third step was irradiation of the coatings: the surface state of the samples irradiated by ultraviolet (UV) or visible light is mentioned as "after UV irradiation" or "after Vis irradiation", respectively.

The irradiation of the films by UV or visible light was carried out using 150 W Xenon lamp (LOMO) equipped with a water filter and UV band pass (250 nm < λpass < 400 nm) or Vis cutoff color filter (λcut = 420 nm). The irradiance was 1.19 and 12.3 mW for UV and Vis irradiation, respectively. For all sample surfaces, the kinetics of the photoinduced water contact angle alteration was presented as a dependence of the water contact angle on the irradiation time.

#### **3. Results**

#### *3.1. Surface Characterization*

The analysis of experimental data obtained by XRD, XPS, SEM and AFM methods (see the Supplementary Materials) confirmed the formation of coatings with desired structures. Indeed, XRD patterns of metal oxide coatings indicate the presence of the corresponding crystal phases of TiO2 (anatase), ZnO (zincite) and Cu2O (cuprite) (see Figures S1 and S2).

SEM images demonstrate that coatings were formed by closely packed nanoparticles with average sizes about 50 nm for TiO2 and about 15 nm for ZnO (see Figures S3 and S4). AFM data show that the roughness of the coating surfaces does not exceed ±5 nm (see Figures S5 and S6), which indicates a nice smoothness of the surface, and, therefore, the surface hydrophilicity of the coatings is not significantly affected by the surface profile. All major characteristics of the prepared nanocoating surfaces are summarized in Table 1.

#### *3.2. Electronic Properties of the Nanocoating Components*

XPS spectra recorded in low binding energy region (see Figure 1) were used to determine positions of the valence bands of the components of heterostructured coatings. Corresponding values in vacuum scale are given in Table 2.

**Figure 1.** Low binding energy XPS spectra of the components of heterostructured coatings: 1, Cu2O; 2, TiO2; and 3, ZnO. **Figure 1.** Low binding energy XPS spectra of the components of heterostructured coatings: 1, Cu2O; 2, TiO2 ; and 3, ZnO.

**Table 2.** Positions of the conduction and valence bands and band gap energies of the coating components with respect to vacuum energy level. **Table 2.** Positions of the conduction and valence bands and band gap energies of the coating components with respect to vacuum energy level.

\* The Ebg values were determined using transmittance spectra of nanocoatings and corresponding Tauc plots (see Figures S7 and S8). \* The Ebg values were determined using transmittance spectra of nanocoatings and corresponding Tauc plots (see Figures S7 and S8).

The Kelvin probe method was applied to measure work function characteristics of the nanocoatings to estimate the positions of the corresponding Fermi levels. The results of the measurements are given in Table 2. Based on the obtained parameters, one can sketch the energetic diagrams of the components forming heterostructured nanocoating, particularly, assuming that energy level of the bottom of conduction band (ECB) is a sum of the energies corresponding to the top of the valence band (EVB) and optical band gap energy (Ebg) of the corresponding components: The Kelvin probe method was applied to measure work function characteristics of the nanocoatings to estimate the positions of the corresponding Fermi levels. The results of the measurements are given in Table 2. Based on the obtained parameters, one can sketch the energetic diagrams of the components forming heterostructured nanocoating, particularly, assuming that energy level of the bottom of conduction band (ECB) is a sum of the energies corresponding to the top of the valence band (EVB) and optical band gap energy (Ebg) of the corresponding components:

$$\mathbf{E\_{CB}} = \mathbf{E\_{VB}} + \mathbf{E\_{bg}} \tag{7}$$

ECB = EVB + Ebg (7) The corresponding energy diagrams of the semiconductor components forming the The corresponding energy diagrams of the semiconductor components forming the heterostructured nanocoatings including Fermi level positions are presented in Figure 2.

heterostructured nanocoatings including Fermi level positions are presented in Figure 2. According to the energy diagrams shown in Figure 2, the formation of Cu2O/TiO2 and Cu2O/ZnO heterostructures should result in electron transfer from Cu2O to either TiO2 or ZnO, both in the dark and under photoexcitation.

2, TiO2; and 3, ZnO.

components with respect to vacuum energy level.

responding Tauc plots (see Figures S7 and S8).

energy (Ebg) of the corresponding components:

**Figure 2.** Energy diagrams of the positions of the top of the valence bands and the bottom of the conduction bands and Fermi level positions of the individual components of the heterostructured coatings.

**Figure 1.** Low binding energy XPS spectra of the components of heterostructured coatings: 1, Cu2O;

**Table 2.** Positions of the conduction and valence bands and band gap energies of the coating

**Sample Ebg \*, eV EVB, eV ECB, eV EF, eV**  TiO2 3.2 −7.4 −4.2 −5.24 ZnO 3.3 −7.8 −4.5 −5.14 Cu2O 2.2 −6.2 −4.0 −5.04 \* The Ebg values were determined using transmittance spectra of nanocoatings and cor-

The Kelvin probe method was applied to measure work function characteristics of the nanocoatings to estimate the positions of the corresponding Fermi levels. The results of the measurements are given in Table 2. Based on the obtained parameters, one can sketch the energetic diagrams of the components forming heterostructured nanocoating, particularly, assuming that energy level of the bottom of conduction band (ECB) is a sum of the energies corresponding to the top of the valence band (EVB) and optical band gap

The corresponding energy diagrams of the semiconductor components forming the heterostructured nanocoatings including Fermi level positions are presented in Figure 2.

ECB = EVB + Ebg (7)

#### *3.3. Effect of Light Irradiation on the Photoinduced Hydrophilicity of Heterostructured Nanocoating Surfaces*

Electron transfer from Cu2O to either TiO2 or ZnO in the dark caused by formation of the corresponding heterostructures should lead to the establishment of new equilibrium states of the heterostructures characterizing by new Fermi level positions located between the Fermi level positions of the individual components. Experimental data of new Fermi level positions of Cu2O/TiO2 and Cu2O/ZnO heterostructures indicate a successful formation of heterostructures (see Figure 2 and the experimental data in Table 3). The effect of photoinduced alteration of the surface hydrophilicity of nanocoatings was studied for strongly hydrated surface. As noted in the Experimental Section, such state of the surface is characterized by good reproducibility and considered as the "initial" surface state for studies of photostimulated alteration of the surface hydrophilicity.

**Table 3.** Work function values (WF) after different treatments for all studied coatings.

Particularly, strong surface wetting leads to alteration of the work function values due to the formation of the multi-layered hydroxyl-hydrated structure of adsorbed water on the surfaces. As evident from experimental data, the effect of surface wetting decreases as follows: TiO2 > TiO2/Cu2O > ZnO > ZnO/Cu2O > Cu2O. That indicates that the most significant interaction between adsorbed water and metal oxide surface resulting in formation of the surface dipole moment affecting work function takes place on the TiO2 surface, while the Cu2O surface is practically unaffected by wetting.

Irradiation of the "initial" state of the coating surfaces with either UV or visible light results in alteration of the surface hydrophilicity (see Figures 3 and 4). To observe the effect of Cu2O substrate on photoinduced alteration of the surface hydrophilicity of TiO2 and ZnO, the kinetics of water contact angle alteration measured for heterostructured coating are compared with the corresponding kinetics for pristine TiO2 and ZnO coatings.

**WaterContact Angle (grad)**

*Nanomaterials* **2021**, *11*, x FOR PEER REVIEW 7 of 12

**Figure 3.** Kinetics of alteration of the water contact angle on the surfaces of TiO2 (1 and 3) and TiO2/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation. **Figure 3.** Kinetics of alteration of the water contact angle on the surfaces of TiO2 (1 and 3) and TiO2/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation. TiO2/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation.

*2*

*4 3 2*

**Irradiation time (min) Figure 4.** Kinetics of alteration of the water contact angle on the surfaces of ZnO (1 and 3) and **Figure 4.** Kinetics of alteration of the water contact angle on the surfaces of ZnO (1 and 3) and ZnO/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation. **Figure 4.** Kinetics of alteration of the water contact angle on the surfaces of ZnO (1 and 3) and ZnO/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation.

*1*

ZnO/Cu2O (2 and 4) coatings under UV (1 and 2) and visible (3 and 4) light irradiation. As evident from the presented dependencies, both TiO2 and ZnO single component coatings demonstrate a surface conversion to superhydrophilic states under UV irradiation and no significant alteration of the surface hydrophilicity under visible light irradiation. However, the presence of the Cu2O substrate in the heterostructured coatings drastically changes the photoinduced hydrophilic behavior of both TiO2 and ZnO surfaces. Indeed, under UV irradiation, both surfaces do not demonstrate any tendency to trans-As evident from the presented dependencies, both TiO2 and ZnO single component coatings demonstrate a surface conversion to superhydrophilic states under UV irradiation and no significant alteration of the surface hydrophilicity under visible light irradiation. However, the presence of the Cu2O substrate in the heterostructured coatings drastically changes the photoinduced hydrophilic behavior of both TiO2 and ZnO surfaces. Indeed, under UV irradiation, both surfaces do not demonstrate any tendency to trans-As evident from the presented dependencies, both TiO2 and ZnO single component coatings demonstrate a surface conversion to superhydrophilic states under UV irradiation and no significant alteration of the surface hydrophilicity under visible light irradiation. However, the presence of the Cu2O substrate in the heterostructured coatings drastically changes the photoinduced hydrophilic behavior of both TiO2 and ZnO surfaces. Indeed, under UV irradiation, both surfaces do not demonstrate any tendency to transform to superhydrophilic state.

form to superhydrophilic state. Conversely, both surfaces show decrease of the surface hydrophilicity: stronger for TiO2 and weaker for ZnO. At the same time, a pronounced decrease of the surface hydrophilicity is also observed under visible light irradiation for both surfaces when neither TiO2 nor ZnO is photoexcited but Cu2O only (see data on their band gaps in Table 2). In general, a main reason and a driving force for alteration of the surface hydrophilicity is the alteration of the surface energy. Therefore, we estimated the total surface energy and its polar and dispersion components using the "two-liquid" approach. The corresponding values of the total surface free energy and its polar and dispersion components for the "initial" surface states of the heterostructured coatings and ZnO and TiO2 coatings and after UV and visible light irradiation are presented in Table 4. The graphical presentation form to superhydrophilic state. Conversely, both surfaces show decrease of the surface hydrophilicity: stronger for TiO2 and weaker for ZnO. At the same time, a pronounced decrease of the surface hydrophilicity is also observed under visible light irradiation for both surfaces when neither TiO2 nor ZnO is photoexcited but Cu2O only (see data on their band gaps in Table 2). In general, a main reason and a driving force for alteration of the surface hydrophilicity is the alteration of the surface energy. Therefore, we estimated the total surface energy and its polar and dispersion components using the "two-liquid" approach. The corresponding values of the total surface free energy and its polar and dispersion components for the "initial" surface states of the heterostructured coatings and ZnO and TiO2 coatings and after UV and visible light irradiation are presented in Table 4. The graphical presentation Conversely, both surfaces show decrease of the surface hydrophilicity: stronger for TiO2 and weaker for ZnO. At the same time, a pronounced decrease of the surface hydrophilicity is also observed under visible light irradiation for both surfaces when neither TiO2 nor ZnO is photoexcited but Cu2O only (see data on their band gaps in Table 2). In general, a main reason and a driving force for alteration of the surface hydrophilicity is the alteration of the surface energy. Therefore, we estimated the total surface energy and its polar and dispersion components using the "two-liquid" approach. The corresponding values of the total surface free energy and its polar and dispersion components for the "initial" surface states of the heterostructured coatings and ZnO and TiO2 coatings and after UV and visible light irradiation are presented in Table 4. The graphical presentation of the tendencies of SFE alteration caused by irradiation is shown in Figures S9–S11.

of the tendencies of SFE alteration caused by irradiation is shown in Figures S9–S11.

**Table 4.** Total (t), polar (p), and dispersive (d) surface free energies (SFE) for all coatings studied

of the tendencies of SFE alteration caused by irradiation is shown in Figures S9–S11.

**Table 4.** Total (t), polar (p), and dispersive (d) surface free energies (SFE) for all coatings studied

after different treatments.

after different treatments.

**Table 4.** Total (t), polar (p), and dispersive (d) surface free energies (SFE) for all coatings studied after different treatments.

*Nanomaterials* **2021**, *11*, x FOR PEER REVIEW 8 of 12

Thus, as evident from the presented data, the transformation of TiO2 and ZnO surfaces to superhydrophilic state caused by UV irradiation is accompanied by an increase of the total surface free energy with a major impact from its polar component for TiO2 and from its dispersive component for ZnO, while visible light irradiation induced only minor alteration of the total surface free energy and its components, which results in very weak changes in the surface hydrophilicity. At the same time, the surfaces of both Cu2O/TiO2 and Cu2O/ZnO heterostructured coatings demonstrate a significant decay of the total surface energy caused by both UV and visible light irradiation, resulting in a decrease of the surface hydrophilicity. Note that the impact of the alteration of polar and dispersive components on changes of the total surface energy and surface hydrophilicity for heterostructured coatings is different for ZnO and TiO2 surfaces and depends on the spectral region of photoexcitation. Indeed, UV irradiation of the Cu2O/TiO2 system results in decay of both polar and dispersive components of TiO2 total surface energy, while, for Cu2O/ZnO, the major impact on the surface energy decay originates from the decrease of the dispersive component, which completely compensates the slight increase of the polar component. At the same time, visible light irradiation (resulting in photoexcitation of Cu2O only) causes mainly the decrease of the dispersive component for Cu2O/TiO2 heterostructure and the decay of the polar component for Cu2O/ZnO system. ZnO/Cu2O 75.3 44.2 31.1 73.7 47.2 26.5 71.4 39.6 31.8 Thus, as evident from the presented data, the transformation of TiO2 and ZnO surfaces to superhydrophilic state caused by UV irradiation is accompanied by an increase of the total surface free energy with a major impact from its polar component for TiO2 and from its dispersive component for ZnO, while visible light irradiation induced only minor alteration of the total surface free energy and its components, which results in very weak changes in the surface hydrophilicity. At the same time, the surfaces of both Cu2O/TiO2 and Cu2O/ZnO heterostructured coatings demonstrate a significant decay of the total surface energy caused by both UV and visible light irradiation, resulting in a decrease of the surface hydrophilicity. Note that the impact of the alteration of polar and dispersive components on changes of the total surface energy and surface hydrophilicity for heterostructured coatings is different for ZnO and TiO2 surfaces and depends on the spectral region of photoexcitation. Indeed, UV irradiation of the Cu2O/TiO2 system results in decay of both polar and dispersive components of TiO2 total surface energy, while, for Cu2O/ZnO, the major impact on the surface energy decay originates from the decrease of the dispersive component, which completely compensates the slight increase of the polar component. At the same time, visible light irradiation (resulting in photoexcitation of Cu2O only) causes mainly the decrease of the dispersive component for Cu2O/TiO2 heterostructure and the decay of the polar component for Cu2O/ZnO system.

In turn, the alteration of the work function values induced by either UV or visible light irradiation (see Table 3 and Figure S12) demonstrates strong decrease of the work function for Cu2O/TiO2 coating and practically no significant alteration for Cu2O/ZnO system. In turn, the alteration of the work function values induced by either UV or visible light irradiation (see Table 3 and Figure S12) demonstrates strong decrease of the work function for Cu2O/TiO2 coating and practically no significant alteration for Cu2O/ZnO system.

#### **4. Discussion**

Both Cu2O/TiO2 and Cu2O/ZnO systems represent type II heterostructures, which provides a condition for charge separation at the corresponding heterojunctions, particularly, electron transfer from Cu2O to either TiO2 or ZnO (see Figures 5 and 6). **4. Discussion**  Both Cu2O/TiO2 and Cu2O/ZnO systems represent type II heterostructures, which provides a condition for charge separation at the corresponding heterojunctions, particularly, electron transfer from Cu2O to either TiO2 or ZnO (see Figures 5 and 6).

**Figure 5.** Schemes of the electron transfer in heterostructured coatings under UV light photoexcitation: (**a**) Cu2O/TiO2 and (**b**) Cu2O/ZnO. **Figure 5.** Schemes of the electron transfer in heterostructured coatings under UV light photoexcitation: (**a**) Cu2O/TiO2 and (**b**) Cu2O/ZnO.

**Figure 6.** Schemes of the electron transfer in heterostructured coatings under visible light photoexcitation: (**a**) Cu2O/TiO2 and (**b**) Cu2O/ZnO. **Figure 6.** Schemes of the electron transfer in heterostructured coatings under visible light photoexcitation: (**a**) Cu2O/TiO2 and (**b**) Cu2O/ZnO.

Consequently, both UV and visible light irradiation of heterostructures results in the increase of the electron concentration compared to the concentration of holes in both TiO2 and ZnO. According to our proposed model described in the Introduction (see Equations (1)–(6)), the alteration of the ratio between electrons and holes at the surface, *n*1/*n*2, can result in a significant alteration of the surface hydrophilicity and turn it, particularly, from a more hydrophilic state to a less hydrophilic state. Therefore, we may conclude that a higher concentration of electrons on the TiO2 and ZnO surfaces induced by electron transfer from Cu2O results in the surface transformation to a less hydrophilic state. Consequently, both UV and visible light irradiation of heterostructures results in the increase of the electron concentration compared to the concentration of holes in both TiO2 and ZnO. According to our proposed model described in the Introduction (see Equations (1)–(6)), the alteration of the ratio between electrons and holes at the surface, *n*1/*n*2, can result in a significant alteration of the surface hydrophilicity and turn it, particularly, from a more hydrophilic state to a less hydrophilic state. Therefore, we may conclude that a higher concentration of electrons on the TiO2 and ZnO surfaces induced by electron transfer from Cu2O results in the surface transformation to a less hydrophilic state.

Remarkably, the formation of Cu2O/TiO2 and Cu2O/ZnO heterostructures also strongly affects the surface free energy behavior under their photoexcitation. Indeed, UV irradiation of the individual TiO2 and ZnO surfaces leads to an increase of the SFE and, particularly, its polar component, which is a main reason for the surface transformation to the superhydrophilic state. However, photoexcitation of heterostructures and electron transfer from Cu2O to either TiO2 or ZnO results in the decay of the total SFE and its polar component. This observation may indicate that excess of electrons at the coating surfaces can destroy the order of the hydroxyl-hydrated multi-layer structure of adsorbed water and, thus, decrease the total surface energy of the system, resulting in a decreasing of the surface hydrophilicity. This conclusion was confirmed by the effect of visible light irradiation when only the Cu2O component of the heterostructures was excited, and, therefore, only electron transfer from Cu2O to TiO2 and ZnO components could affect both SFE and surface hydrophilicity. Photoinduced destruction of the order in the hydroxyl-hydrated multi-layer structure of adsorbed water is also (indirectly) confirmed by the decrease of the work function of the heterostructures under irradiation, which might be caused by a decrease of the surface dipole moment due to the disorder of the adsorbed water structure in the hydroxyl-hydrated multilayer. Remarkably, the formation of Cu2O/TiO2 and Cu2O/ZnO heterostructures also strongly affects the surface free energy behavior under their photoexcitation. Indeed, UV irradiation of the individual TiO2 and ZnO surfaces leads to an increase of the SFE and, particularly, its polar component, which is a main reason for the surface transformation to the superhydrophilic state. However, photoexcitation of heterostructures and electron transfer from Cu2O to either TiO2 or ZnO results in the decay of the total SFE and its polar component. This observation may indicate that excess of electrons at the coating surfaces can destroy the order of the hydroxyl-hydrated multi-layer structure of adsorbed water and, thus, decrease the total surface energy of the system, resulting in a decreasing of the surface hydrophilicity. This conclusion was confirmed by the effect of visible light irradiation when only the Cu2O component of the heterostructures was excited, and, therefore, only electron transfer from Cu2O to TiO2 and ZnO components could affect both SFE and surface hydrophilicity. Photoinduced destruction of the order in the hydroxyl-hydrated multi-layer structure of adsorbed water is also (indirectly) confirmed by the decrease of the work function of the heterostructures under irradiation, which might be caused by a decrease of the surface dipole moment due to the disorder of the adsorbed water structure in the hydroxyl-hydrated multilayer.

Summing up, we may conclude that the electron transfer realized in Cu2O/TiO2 and Cu2O/ZnO heterostructures under photoexcitation leads to an increase of the concentration of electrons on the TiO2 and ZnO surfaces that causes a higher disorder in the hydroxyl-hydrated multi-layer structure of the adsorbed water, thus resulting in the decay of the SFE and the surface hydrophilicity. Summing up, we may conclude that the electron transfer realized in Cu2O/TiO2 and Cu2O/ZnO heterostructures under photoexcitation leads to an increase of the concentration of electrons on the TiO2 and ZnO surfaces that causes a higher disorder in the hydroxylhydrated multi-layer structure of the adsorbed water, thus resulting in the decay of the SFE and the surface hydrophilicity.

#### **5. Conclusions**

**5. Conclusions**  We demonstrated the effect of Cu2O substrate on the effect of the photoinduced alteration of the surface hydrophilicity of TiO2 and ZnO surfaces under both UV and visible light irradiation. Particularly, it was shown that under UV light both surfaces can be transformed to the superhydrophilic state, while both UV and visible light irradiation of We demonstrated the effect of Cu2O substrate on the effect of the photoinduced alteration of the surface hydrophilicity of TiO2 and ZnO surfaces under both UV and visible light irradiation. Particularly, it was shown that under UV light both surfaces can be transformed to the superhydrophilic state, while both UV and visible light irradiation of the heterostructured coating results in a decrease of the surface hydrophilicity. This effect of

the heterostructured coating results in a decrease of the surface hydrophilicity. This effect

the Cu2O substrate on the photoinduced hydrophilic behavior of TiO2 and ZnO surfaces is explained in terms of charge separation at corresponding heterojunctions and enrichments of TiO2 and ZnO surfaces with electrons transferred from Cu2O. Consequently, we may conclude that an increase of the free electron concentration on the surface of both TiO2 and ZnO leads to the decay of their surface hydrophilicity. Thus, the purposeful formation of a certain type heterostructures may help to control the surface hydrophilicity at the desired level and the direction of its photostimulated alteration.

**Supplementary Materials:** The following are available online at https://www.mdpi.com/article/ 10.3390/nano11061526/s1, Figure S1: XRD patterns of Cu2O, TiO2 and TiO2/Cu2O heterostructured coatings, Figure S2: XRD patterns of Cu2O, ZnO and ZnO/Cu2O heterostructured coatings, Figure S3: SEM images of the TiO2 surface and TiO2/Cu2O cross-section of the TiO2/Cu2O coating, Figure S4: SEM images of the TiO2 surface and TiO2/Cu2O cross-section of the TiO2/Cu2O coating, Figure S5: AFM image of the TiO2 surface and roughness profile of the TiO2/Cu2O coating, Figure S6: AFM image of the ZnO surface and roughness profile of the ZnO/Cu2O coating, Figure S7: Transmittance spectra of: (a) TiO2/Cu2O heterostructured coating and its components; and (b) ZnO/Cu2O heterostructured coating and its components, Figure S8: Tauc plots for 1 – Cu2O, 2 – ZnO, and 3 – TiO2 ., Figure S9: Alteration of the total SFE caused by irradiation with visible and UV light, Figure S10: Alteration of the SFE polar component caused by irradiation with visible and UV light, Figure S11: Alteration of the SFE dispersive component caused by irradiation with visible and UV light, Figure S12: Alteration of the work function of nanocoatings induced by wetting and UV and visible light irradiation.

**Author Contributions:** Conceptualization, A.V.R. and D.W.B.; methodology, A.V.E.; software, M.V.M.; validation, A.V.R. and A.V.E.; formal analysis, M.V.M.; investigation, M.V.M. and A.V.E.; resources, D.W.B.; writing—original draft preparation, M.V.M. and A.V.E.; writing—review and editing, A.V.E., A.V.R. and D.W.B.; visualization, M.V.M.; supervision, D.W.B.; project administration, A.V.E.; and funding acquisition, A.V.E. and D.W.B. All authors have read and agreed to the published version of the manuscript.

**Funding:** The reported study was funded by RFBR, project number 19-32-90111.

**Acknowledgments:** AVE and MVM are grateful to Russian Foundation for Basic Research (RFBR) for grant N 19-32-90111 to support this research. Experimental studies were performed in laboratory "Photoactive nanocomposite materials" supported by Saint-Petersburg State University (ID: 73032813). We are thankful to the Resource Center (RC) "Nanophotonics", RC "Centre for Innovative Technologies of Composite Nanomaterials", RC "X-ray Diffraction Studies", RC "Centre for Physical Methods of Surface Investigation" and RC "Nanotechnology" of the Research Park at the Saint-Petersburg State University for helpful assistance in the preparation and characterization of the samples.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**

## *Article* **Investigation of the Effects of Rapid Thermal Annealing on the Electron Transport Mechanism in Nitrogen-Doped ZnO Thin Films Grown by RF Magnetron Sputtering**

**Simeon Simeonov 1 , Anna Szekeres 1,\* , Dencho Spassov 1 , Mihai Anastasescu 2 , Ioana Stanculescu 3,4 , Madalina Nicolescu 2,\* , Elias Aperathitis 5 , Mircea Modreanu 6 and Mariuca Gartner 2**

**Abstract:** Nitrogen-doped ZnO (ZnO:N) thin films, deposited on Si(100) substrates by RF magnetron sputtering in a gas mixture of argon, oxygen, and nitrogen at different ratios followed by Rapid Thermal Annealing (RTA) at 400 ◦C and 550 ◦C, were studied in the present work. Raman and photoluminescence spectroscopic analyses showed that introduction of N into the ZnO matrix generated defects related to oxygen and zinc vacancies and interstitials. These defects were deep levels which contributed to the electron transport properties of the ZnO:N films, studied by analyzing the current–voltage characteristics of metal–insulator–semiconductor structures with ZnO:N films, measured at 298 and 77 K. At the appliedtechnological conditions of deposition and subsequent RTA at 400 ◦C n-type ZnO:N films were formed, while RTA at 550 ◦C transformed the n-ZnO:N films to p-ZnO:N ones. The charge transport in both types of ZnO:N films was carried out via deep levels in the ZnO energy gap. The density of the deep levels was in the order of 1019 cm−3 . In the temperature range of 77–298 K, the electron transport mechanism in the ZnO:N films was predominantly intertrap tunneling, but thermally activated hopping also took place.

**Keywords:** RF magnetron sputtering; ZnO:N thin films; Raman spectroscopy; photoluminescence spectroscopy; electrical characteristics; charge carrier transport properties

#### **1. Introduction**

Zinc oxide possesses remarkable optical and semiconductor properties, such as a direct wide gap around 3.3 eV at room temperature and a large exciton binding energy of about 60 meV [1]. Because of these properties, ZnO has huge prospects in applications such as optoelectronic devices [2], homojunction LEDs [3,4], solar cells [5], sensors [6,7], and other devices and structures. However, the use of ZnO films in these applications requires the ability to control the majority-carrier type and concentration. An asymmetry exists between n- and p-type doping of ZnO thin films. While it is possible to prepare stable n-type ZnO films even without the introduction of donor dopants during the ZnO film deposition, obtaining p-type ZnO film requires thermal activation of acceptor dopants incorporated into the ZnO film during its deposition [8]. In some cases, an additional

**Citation:** Simeonov, S.; Szekeres, A.; Spassov, D.; Anastasescu, M.; Stanculescu, I.; Nicolescu, M.; Aperathitis, E.; Modreanu, M.; Gartner, M. Investigation of the Effects of Rapid Thermal Annealing on the Electron Transport Mechanism in Nitrogen-Doped ZnO Thin Films Grown by RF Magnetron Sputtering. *Nanomaterials* **2022**, *12*, 19. https:// doi.org/10.3390/nano12010019

Academic Editors: Christian Mitterer and Aurora Rizzo

Received: 30 October 2021 Accepted: 18 December 2021 Published: 22 December 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

problem is the subsequent transformation of the p-type ZnO films into n-type ZnO films [9]. This is why producing stable and qualitative p-type ZnO films is still a challenge facing the technologies of the above-listed devices and structures.

Much effort has been made to obtain p-type ZnO films by applying different advanced deposition methods [10–14] and introducing different elements into ZnO as additives [15]. Among the acceptor dopants, nitrogen is considered themost promising [9]. The nitrogen atom possesses three 2p valence electrons, while the oxygen atom has four 2p valence electrons. The size of the nitrogen atom and the energy of the N 2p valence electrons are closest to the corresponding values of the oxygen atom [16]. Therefore, one might expect the replacement of some O anions in the anion sublattice with N anions and the formation of shallow acceptor levels in the ZnO energy gap. However, to prepare a stable p-type ZnO film by N doping is rather difficult due to the low solubility of nitrogen, and thus the low concentration of holes in the films. Another obstacle to effective p-type doping via the replacement of O by N in the anion sublattice is the simultaneous generation of shallow and deep donor-type defects in the ZnO energy gap, such as oxygen vacancies in the O sublattice, VO, and zinc interstitials, Zni [17–19], which leads to self-compensation of N acceptors in ZnO [20]. Difficulties in the preparation of p-type ZnO films have provoked interest in the studies of ZnO heterojunctions, mainly ZnO–Si heterojunctions. The purpose of these investigations is to expand the applications of ZnO thin films in photoelectron devices and structures, especially in solar cell structures.

Zinc oxide is also the subject of our research. As a pathway for ZnO thin film deposition, we applied a radio frequency (RF) magnetron sputtering technique. This deposition technique combines the relative easiness of both nitrogen doping of the growing layers by the simple introduction of N2 into the process chamber with control of the deposition parameters (such as gas pressure, substrate temperature, scattering power, and deposition rate). The ability to precisely maintain these parameters during film deposition determines the reproducibility and efficiency of this method. Because of this, RF magnetron sputtering is one of the most widely used methods for thin-film deposition. In order to improve the properties of our ZnO:N films, we applied postdeposition rapid thermal annealing (RTA) in a nitrogen atmosphere at different temperatures. The choice of RTA was due to the fact that it is more technologically relevant than conventional thermal treatments that take much longer and are more highly energy consuming. Moreover, during the RTA process, nitrogen release from the ZnO:N layer is less probable. It has been shown that the RTA method is a particularly suitable method for improving the structural, optical, and electrical properties of ZnO films [21,22].

In order to use nitrogen-doped ZnO films in ZnO homo- and heterojunctions and/or as transparent conductive oxide films, a detailed characterization of their structural, optical, and electrical properties is required. It is particularly important to understand, and hence to tailor, the charge transfer processes in these materials. Accordingly, detailed structural studies including X-ray Diffraction (XRD), Transmission Electron Microscopy (TEM), Atomic Force Microscopy (AFM), and Scanning Electron Microscopy (SEM) were preliminarily conducted on the RF-sputtered N-doped ZnO films [23–25]. We established that RTA treatment in N2 at moderate temperatures essentially improves the crystalline state of the ZnO films, magnetron sputtered onto crystalline silicon [23], or fused silica [24] substrates, yielding a polycrystalline columnar structure with nanocrystallites of 9–13 nm, preferentially oriented in the (002) direction. In these studies, the ZnO films were deposited in a gas mixture of Ar:O2:N2 = 50:40:10 and at a total pressure of 0.66 Pa (5 m Torr). Further, we collected the same measurements as mentioned above on ZnO films deposited under increased amounts of N2 and a constant amount of Ar. It was established that variation of the Ar:O2:N2 mixture to 50:25:25 and 50:10:40 weakly affected the films' crystallinity and surface morphology (the root-mean-square roughness values being lower than 5 nm), but influenced the optical properties [25], and hence alteration in the electrical properties is also expected.

The main purpose of the present research was to investigate the electrical properties of RF magnetron-sputtered nitrogen-doped ZnO films and to elucidate the charge transport mechanism in these films. There are few data in the literature concerning the charge carrier transport through N-doped ZnO films and, peculiarly, the charge transport parameters. For this purpose, we prepared Al–ZnO:N–Si–Al structures and their electrical properties were studied in detail. In addition, Raman and Photoluminescence (PL) measurements were carried out, aiming to establish the origin and nature of the deep levels associated with inherent defects in the polycrystalline ZnO:N films. This could shed more light on electron transitions and charge carrier transport involving N-induced defective centers. The final goal of the research was elucidation of the electron transport properties of zinc-oxidebased films, which would expand their application area in optoelectronics, for example as transparent ZnO:N n-p homojunctions.

#### **2. Materials and Methods**

#### *2.1. Film Preparation*

The nitrogen-doped ZnO thin films were deposited by RF magnetron sputtering employing Nordiko RFG-2500 equipment (Nordiko Technical Service Ltd., Havant, Hampshire, United Kingdom). Commercially available zinc nitride target (Testbourne Ltd., Basingstoke, Hampshire, UK), Zn:N = 1:1, purity 99.9%, 6 in diameter ×0.25 in thick) was sputtered in a gas mixture of Ar, O2, and N2 onto p-type Si(100) substrates with specific resistivity over 100 Ω cm. In general, the substrates never occupied an area bigger than ~12 cm2 on the holder. The distance between the target and substrates was 11 cm.

The substrates were ultrasonically cleaned in acetone and isopropanol, rinsed in deionized water, and dried in flowing nitrogen gas. The native oxide on the Si surface was etched in HF solution before the substrates were introduced into the sputtering system. After reaching the base pressure of 1.33 × 10−5 Pa (10−7 m Torr) and prior to deposition, the target was presputtered for at least 15 min (Ar plasma, 0.66 Pa (5 m Torr), 100 W RF power) to remove any contaminants from the target surface and to enable equilibrium conditions to be reached. The RF power was kept at 100 W. The sputtering was carried out at a total pressure of 0.66 Pa (5 m Torr) in a gas mixture of Ar:O2:N2, keeping Ar constant at 50% and varying the ratio of O2 and N2 percentages: (i) O2:N2 = 40:10, (ii) O2:N2 = 10:40, and (iii) O2:N2 = 25:25. During deposition, no external heating was applied to the substrates. At the given technological conditions, the deposition rate of the ZnO films was 0.83 nm/min for the films made in the oxygen-rich plasma and 0.32 nm/min for the films made in oxygen-deficient plasma.

After deposition, and following our previous investigations concerning postdeposition annealing conditions as mentioned above [23–25], the ZnO:N samples were subjected to RTA at temperatures of 400 ◦C and 550 ◦C for 1 min in N2 atmosphere. The annealing temperature of 550 ◦C was not exceeded in order to avoid any possible release of nitrogen from the structure of the films [26].

The denotations of the ZnO:N–Si samples prepared under different technological conditions are presented in Table 1. In the table, we have included the thicknesses of the corresponding films, which were determined with an accuracy of ±0.2 nm from the spectral ellipsometric measurements. The scatter of the as-deposited films' thickness resulted mainly from the low deposition rates and different sputtering durations. It should be pointed out that the thickness uniformity of the films deposited on given Si samples was ~96%, as established by the ellipsometric mapping [25].

**Table 1.** Technological conditions and corresponding sample number of ZnO:N films with the given layer's thickness.

The bandgap energy values of ZnO:N films deposited under the same technological condition as were used in the present work have been determined from the ellipsometric data analysis published elsewhere [25]. It was found that for the as-deposited ZnO:N films, the bandgap energy value was 3.25 ± 0.025 eV, while for the 400 ◦C and 550 ◦C RTA samples there was a slight reduction toward 3.20 ± 0.025 eV. These optical bandgap energy values are in good agreement with the literature reports [17,19,27,28].

#### *2.2. Characterization Methods*

Raman scattering spectra of the films were recorded on a complex configuration Bruker Vertex 70 FTIR/FT-Raman instrument (Bruker, Ettlingen, Germany). The spectrometer equipped with a RAM II module using a Nd:YAG laser (1064 nm) with variable power (1–500 mW) and LN2 cooled Ge detector. The Fourier-transform Raman spectra were registered between 200 and 600 cm−1 with 512 scans and 4 cm−1 spectral resolution, and a laser power of 1 mW.

The photoluminescence (PL) measurements were carried out on a Carry Eclipse (Agilent Technologies, Melbourne, Australia), fluorescence spectrometer, with the slits set at 20 nm. The photoluminescence at room temperature was excited by the 325 nm line from a Xe bulb lamp. The PL spectra were taken with a scan rate of 120 nm/min and spectral resolution of 0.5 nm.

The electrical measurements were conducted on metal–ZnO:N–silicon capacitors (further denoted as metal–insulator–semiconductor (MIS) structures) formed by vacuum thermal evaporation of Al dots with 1.96 × 10−3 cm2 area onto the ZnO top surface through a metal mask, while continuous Al film was evaporated as a contact to the silicon backside.

The electrical properties of the MIS structures were examined from the current–voltage (I–V) characteristics, measured in two ways. In the first way, the I–V characteristics were automatically measured at room temperature in a dark chamber using a Keithley 4200 Semiconductor Characterization System with a ramp rate of 0.1 V/s. In the second way, the I–V curves were measured point by point with a cycle sequence starting from 0 V toward negative or positive voltages, with maximal amplitude applied to the top Al–dot contact followed by a voltage reversal toward zero voltage. The first measurement cycle starting from zero up to maximal applied voltage, Va, is denoted further as the initial stage. The duration of each measurement cycle was approximately 20–25 min. This kind of I–V measurement was applied to samples 2.1 and 2.3 at temperatures of 298 and 77 K. In this way, the presence and behavior of deep levels in the ZnO:N films could be examined and their concentrations determined.

In addition, the impedance of the MIS structures with ZnO layers 2.1 and 2.3 was measured at room temperature with a Tesla BM-507 impedance meter (TESLA, Praha, Czechoslovakia) applying test voltage frequencies in the range of 0.5–500 kHz. Herein, we consider only the parallel conductance, Gm, values, calculated from the measured |Zm| and ϕm quantities using the expression Gm = cos(ϕm)/|Zm|.

#### **3. Results and Discussion**

The structural and morphological properties of the ZnO:N films (Table 1), as well as their chemical compositions, were reported in our previous papers [23–25]. XRD investigations revealed that all ZnO:N films were crystallized in a hexagonal ZnO wurtzite phase (JCPDS data card 36-1451), (002) oriented, and exhibited an improvement in crystallinity after RTA. There was no noticeable change in the XRD patterns induced by the variation of the nitrogen content in the sputtering reactor atmosphere during ZnO:N film deposition, but the values of the crystallite size, estimated using the Scherrer formula, increased with the RTA temperature and with N2 content in the sputtering atmosphere, from 14.4 nm (10% N2) to 15.7 nm (40% N2) for the ZnO:N films after RTA at 550 ◦C [25]. From a topographic point of view, all ZnO:N–Si samples (Table 1) showed a homogeneous distribution of small and rounded superficial grains, with an average diameter in the range of 40–70 nm, all surfaces being smooth, with a root-mean-square (RMS) roughness of lessthan 5 nm. From TEM investigations, it was found that the morphology of the ZnO:N films was columnar, with the grain column axis oriented nearly parallel with the hexagonal <001> axis of the ZnO structure. The columnar structures had variable diameters from 10 nm at the bottom to 50 nm at the top (for the films sputtered under 40% nitrogen), so that the surface had larger grains which were well faceted, as seen by AFM. The columns of the as-deposited films contained a lot of defects and pores which diminished after RTA. The presence of the nitrogen inside the ZnO:N films was not detected by XRD analysis due to its very small amount, but was evidenced in Energy-dispersive X-ray spectroscopy (EDX) observations conducted in a ratio of 1/6 (nitrogen/oxygen) for the films deposited on Si, regardless of the RTA treatment. However, the presence of nitrogen was clearly pointed out by X-ray photoelectron spectroscopy (XPS) in the RTA samples only as zinc nitride and diluted zinc oxynitride, as reported previously [23,24]. Nevertheless, the uniformity of thickness deposition, determined by ellipsometric mapping, was found to be around 96% for the ZnO:N films deposited on silicon [25].

#### *3.1. Analysis of FT-Raman Spectra*

In Figure 1, as a representative illustration, the Raman spectra of the ZnO:N film deposited in gas mixture of Ar:O2:N2 = 50:10:40, (sample 2.1) and RT annealed at 550 ◦C (sample 2.3) are given. The nonsymmetrical shape of the Raman peaks is due to the multiple contributions from the polycrystalline nature of the ZnO:N [25]. The vibrational bands of Si–Si bonds are also indicated, appearing as a strong peak at 523 cm−1 and as a weak and broad peak in the 303–308 cm−1 spectral range. The kind and position of the observed vibrational modes of ZnO:N films in dependence on the oxygen/nitrogen content and RTA temperature are summarized in Table 2.

Pure, crystalline ZnO has eight sets of optical phonon modes at the Γ point in the Brillouin zone, Γopt, expressed as Γopt = A1+ E1 (IR, R) + 2E2 (R) + 2B1 [29]. The A1 and E1 modes are both Raman and infrared active and split into transverse optical (TO) and longitudinal optical (LO) phonon modes. The E2 modes are Raman active and nonpolar at low- (E2 (low)) and high-frequency (E2 (high)) phonon modes. The E2 (low) mode involves mainly Zn sublattice motion while E2 (high) is associated with the vibration of oxygen atoms [30]. There are also two B1 modes, B1 (low) and B1 (high), which are inactive Raman modes but can be activated by introducing defects or by doping with other elements [29,31,32].

The recorded FT-Raman spectra exhibited the state of molecular Zn–O bonds in the ZnO:N films under changing technological conditions. The increase of the peaks' intensity after RTA treatment implies the increase of defects related to O and Zn vacancies and interstitials. The bands that appeared at 230–240 cm−1 and 330 cm−1 correspond to (2TA and 2E2 (low)) and E2 (high)-E2 (low), respectively, and are assigned to the contribution of second-order vibrations, two-phonon modes which become activated in thin films. The peak around 377 cm−1 was identified as the A1(TO) mode, which indicates the displacement of Zn2+ and O2− ions parallel to the c-axis, near the center of the Brillouin zone (Γ point) [30,33,34]. In comparison to the pure crystalline ZnO, the position of the wide

E2 (high) peak at 410–420 cm−1 was significantly shifted in the doped ZnO:N films, indicating that nitrogen doping mainly affected the oxygen bonds. The peak at 550 cm−1 was assigned to the inactive A1(LO) mode, activates by introducing defects or by doping with other elements [25,31–33]. The observed peak at 590 cm−1 was attributed to the E1(LO) mode that appeared due to oxygen vacancies VO, zinc interstitials Zni , and free carriers [35,36] existing in the analyzed ZnO films. The Raman study proved that the incorporation of nitrogen into the films caused disordering of the ZnO lattice. Considering that the spectra were taken with a resolution of 4 cm−1 , the close positions of the corresponding peaks indicated similar structure and chemical bonding. The exceptions were the peaks around 402–412 cm−1 , confirming that the change in the oxygen/nitrogen ratio in the Ar:O2:N2 mixture predominantly affectedthe oxygen bonds in the ZnO lattice. *Nanomaterials* **2022**, *12*, x FOR PEER REVIEW 6 of 18

**Figure 1.** Representative Raman spectra of ZnO:N films, deposited in a gas mixture of Ar:O2:N2 = 50:10:40 (sample 2.1) and after RTA at 550 °C (sample 2.3). **Figure 1.** Representative Raman spectra of ZnO:N films, deposited in a gas mixture of Ar:O2 :N2 = 50:10:40 (sample 2.1) and after RTA at 550 ◦C (sample 2.3).

#### ZnO:N films under changing technological conditions. The increase of the peaks' intensity *3.2. Photoluminescence Analysis*

after RTA treatment implies the increase of defects related to O and Zn vacancies and interstitials. The bands that appeared at 230–240 cm−1 and 330 cm−1 correspond to (2TA In general, in the PL spectra of ZnO films, two emission bands can be observed: one is in the UV region within the 361–369 nm range and the second band is in the visible range of 450–550 nm [32,37]. As the excitation source (Xe bulb lamp) was not intense enough to

**Table 2.** Position of the Raman phonon modes in the ZnO:N thin films.

**RTA Temperature (**°**C)**

and 2E2(low)) and E2(high)-E2(low), respectively, and are assigned to the contribution of secondorder vibrations, two-phonon modes which become activated in thin films. The peak

Zn+2 and O−2 ions parallel to the c-axis, near the center of the Brillouin zone (Γ point) [30,33,34]. In comparison to the pure crystalline ZnO, the position of the wide E2(high) peak at 410–420 cm−1wassignificantly shifted in the doped ZnO:N films, indicating that nitrogen doping mainly affected the oxygen bonds. The peak at 550 cm−1wasassigned to the inactive A1(LO) mode, activates by introducing defects or by doping with other elements [25,31– 33]. The observed peak at 590 cm−1wasattributed to the E1(LO) mode that appeared due to oxygen vacancies VO, zinc interstitials Zni, and free carriers [35,36] existing in the analyzed ZnO films. The Raman study proved that the incorporation of nitrogen into the films caused disordering of the ZnO lattice. Considering that the spectra were taken with a resolution of 4 cm−1, the close positions of the corresponding peaks indicated similar structure and chemical bonding. The exceptions were the peaks around 402–412 cm−1, confirming that the change in the oxygen/nitrogen ratio in the Ar:O2:N2 mixture predominantly af-

> **Raman Molecular Vibration Modes (cm−1)**

as-deposited 221 332 373 411 555 589

400 224 338 377 415 555 590 550 234 334 378 409 557 595

fectedthe oxygen bonds in the ZnO lattice.

**Gas Ratio of Ar:O2:N2 (%)**

50:40:10

excite the transitions associated with the optical band gap energy, the first characteristic band was not observed. In Figure 2, typical PL spectra in the spectral range of 400–560 nm are presented for a ZnO:N film deposited in gas mixture of Ar:O2:N2 = 50:10:40 (sample 2.1) and after RTA annealing at 550 ◦C (sample 2.3). sulted from radiative recombination of holes with electrons at the singly ionized intrinsic oxygen vacancies [40,41]. These kinds of defects are deep levels, which contribute to the electron transport properties and were considered further at the data analysis of electrical measurements.

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(sample 2.1) and after RTA annealing at 550 °C (sample 2.3).

as-deposited 237 336 376 412 553 594

as-deposited 236 330 376 402 552 593

In general, in the PL spectra of ZnO films, two emission bands can be observed: one is in the UV region within the 361–369 nm range and the second band is in the visible range of 450–550 nm [32,37]. As the excitation source (Xe bulb lamp) was not intense enough to excite the transitions associated with the optical band gap energy, the first characteristic band was not observed. In Figure 2, typical PL spectra in the spectral range of 400–560 nm are presented for a ZnO:N film deposited in gas mixture of Ar:O2:N2 = 50:10:40

The representative PL spectra of the studied ZnO:N films were characterized by a weak band around 455–460 nm and more strong and broad bands around 412 nm and 500 nm. Due to the broad PL peaks, their locations were difficult to resolve accurately. For this reason, we tried to find the Gaussian components until the best fit was achieved. The dominant 500 nm band was deconvoluted into two Gaussian components as the peaks

The emission peak from the band-to-band recombinations was expected to appear around 382 nm, the wavelength corresponding to the bandgap energy (3.25 ± 0.025 eV) of our ZnO films. However, in the measured PL spectra, the emission peak was centered within the 410–420 nm region, and hence it originated most probably from defect states close to the ZnO bandgap edges. Previous research suggests that zinc interstitial defects, Zni, creating shallow states underneath the conduction band edge, are a possible reason for the appearance of this PL peak [38]. The emission peak at ~490 nm (2.53 eV) was a typical blue green emission of ZnO, which may be attributed to a high density of point defects such as zinc vacancy, VZn, oxygen vacancy, VO, and interstitial oxygen, Oi, in the polycrystalline structure of ZnO [39]. The green emission peak at ~525 nm (2.36 eV) re-

400 233 329 372 402 552 590 550 236 330 376 418 551 593

400 238 335 375 407 554 591 550 239 337 376 411 550 595

50:25:25

50:10:40

*3.2. Photoluminescence Analysis*

were situated around 490 and 524 nm.

**Figure 2.** Representative PL spectra of the ZnO:N films deposited in a gas mixture of Ar:O2:N2 = 50:10:40:(**a**) as-deposited (sample 2.1) and (**b**) after RTA at 550 °C (sample 2.3). **Figure 2.** Representative PL spectra of the ZnO:N films deposited in a gas mixture of Ar:O2 :N2 = 50:10:40: (**a**) as-deposited (sample 2.1) and (**b**) after RTA at 550 ◦C (sample 2.3).

The representative PL spectra of the studied ZnO:N films were characterized by a weak band around 455–460 nm and more strong and broad bands around 412 nm and 500 nm. Due to the broad PL peaks, their locations were difficult to resolve accurately. For·this reason, we tried to find the Gaussian components until the best fit was achieved. The dominant 500 nm band was deconvoluted into two Gaussian components as the peaks were situated around 490 and 524 nm.

The emission peak from the band-to-band recombinations was expected to appear around 382 nm, the wavelength corresponding to the bandgap energy (3.25 ± 0.025 eV) of our ZnO films. However, in the measured PL spectra, the emission peak was centered within the 410–420 nm region, and hence it originated most probably from defect states close to the ZnO bandgap edges. Previous research suggests that zinc interstitial defects, Zni , creating shallow states underneath the conduction band edge, are a possible reason for the appearance of this PL peak [38]. The emission peak at ~490 nm (2.53 eV) was a typical blue green emission of ZnO, which may be attributed to a high density of point defects such as zinc vacancy, VZn, oxygen vacancy, VO, and interstitial oxygen, Oi , in the polycrystalline structure of ZnO [39]. The green emission peak at ~525 nm (2.36 eV) resulted from radiative recombination of holes with electrons at the singly ionized intrinsic oxygen vacancies [40,41]. These kinds of defects are deep levels, which contribute to the electron transport properties and were considered further at the data analysis of electrical measurements.

The positions of emission peaks obtained by Gaussian deconvolution of the PL spectra of ZnO:N films are summarized in Table 3. It can be observed that for the as-deposited films, the increase of the nitrogen content from 10 to 40% resulted in a shift of the Gaussian peaks' position from 493 nm and 524 nm to 498 nm and 530 nm, respectively. This could be attributed to the increased defect generation and disordering of the ZnO lattice by nitrogen doping, as detected by the Raman measurements above. During heat treatment, these defects were partially annealed and, as a consequence, the Gaussian peaks moved to smaller wavelengths with increasing RTA temperature.

The positions of emission peaks obtained by Gaussian deconvolution of the PL spectra of ZnO:N films are summarized in Table 3. It can be observed that for the as-deposited films, the increase of the nitrogen content from 10 to 40% resulted in a shift of the Gaussian peaks' position from 493 nm and 524 nm to 498 nm and 530 nm, respectively. This could be attributed to the increased defect generation and disordering of the ZnO lattice by nitrogen doping, as detected by the Raman measurements above. During heat treatment, these defects were partially annealed and, as a consequence, the Gaussian peaks moved

The above Raman and PL measurements show that inherent structural defects and

**Table 3.** The positions of the PL Gaussian emission peaks of the ZnO:N films. N-induced defect centers were present in the ZnO films. These can be associated with

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to smaller wavelengths with increasing RTA temperature.

The above Raman and PL measurements show that inherent structural defects and N-induced defect centers were present in the ZnO films. These can be associated with localized states in the energy gap of ZnO, which are traps for electric charge carriers. For example, it has been established that the dominant defect centers of oxygen vacancy (PL peak at ~ 492 nm) develop deep energetic levels of 1.4 eV for electrons and 1.6 eV for holes in the ZnO bandgap [42]. 50:25:25 as-deposited 495 527 400 495 523 550 495 522 50:10:40 as-deposited 498 530 400 496 525 550 490 524

#### *3.3. Electric Charge Transport Properties 3.3. Electric Charge Transport Properties*

In order to prove the role of deep levels in the conduction mechanism and to study the current via these deep levels, we conducted a detailed study of the I–V characteristics of the MIS structures formed with the ZnO:N films. The I–V characteristics of the MIS structures with the ZnO:N films, automatically recorded at room temperature, are summarized in Figure 3. A schematic representation of the studied MIS structures is shown as the lower left inset in Figure 3a. In order to prove the role of deep levels in the conduction mechanism and to study the current via these deep levels, we conducted a detailed study of the I–V characteristics of the MIS structures formed with the ZnO:N films. The I–V characteristics of the MIS structures with the ZnO:N films, automatically recorded at room temperature, are summarized in Figure 3. A schematic representation of the studied MIS structures is shown as the lower left inset in Figure 3a.

**Figure 3.** Automatically recorded I–V characteristics of MIS structures with ZnO:N films, deposited at different gas ratios: Ar:O2:N2 = 50:40:10 (**a**), Ar:O2:N2 = 50:10:40 (**b**), and Ar:O2:N2 = 50:25:25 (**c**). The corresponding sample numbers are given in the insets. The lower left inset in (**a**) shows the schematic structure of the measured samples. **Figure 3.** Automatically recorded I–V characteristics of MIS structures with ZnO:N films, deposited at different gas ratios: Ar:O2 :N2 = 50:40:10 (**a**), Ar:O2 :N2 = 50:10:40 (**b**), and Ar:O2 :N2 = 50:25:25 (**c**). The corresponding sample numbers are given in the insets. The lower left inset in (**a**) shows the schematic structure of the measured samples.

In the case of as-deposited ZnO:N films (samples 1.1, 2.1, and 3.1) and those after RTA treatment at 400 ◦C (samples 1.2, 2.2 and 3.2), the current at positive voltages applied to the top Al-dot electrode was higher than the current in the case of negative voltage applied to the same contact. These MIS structures had diode-like properties. At absolute voltages of 8–10 V, the rectification ratios were ~4600, 7055, and 2500 for the as-deposited films obtained at 10%, 25%, and 40% N2 in the gas mixture, respectively. This ratio decreased after 400 ◦C RTA but still remained in the order of 103 . The strongest rectification effect was observed in ZnO:N films grown at 25% N2 in the gas mixture. The existence of asymmetry between the currents at positive and negative voltage values confirms the formation of n-ZnO–p-Si heterojunction. From these observations, it follows that the nitrogen-doped ZnO films were initially n-type and remained n-type after 400 ◦C annealing. In the case of the films deposited at gas ratios of Ar:O2:N2 = 50:40:10 and Ar:O2:N2 = 50:25:25, the comparison of the measured I–V curves of the as-deposited ZnO:N films with those of the corresponding ones after RTA at 400 ◦C showed close current values, which suggests similar concentrations of the N dopants.

For the samples after RTA at 550 ◦C (Figure 3, samples 1.3, 2.3, and 3.3), the current through the ZnO film was higher when negative voltage was applied to the top Al-dot electrode. The asymmetry between the currents at positive and negative voltage values was less pronounced, which was expressed in considerably smaller rectification ratios. At absolute voltages of 8–10 V, their values were 230, 534, and 24 for the films deposited at 10%, 25%, and 40% N2 in the gas mixture, respectively. These results imply that RTA at 550 ◦C transformed the nitrogen-doped ZnO films to p-type ZnO ones. In accordance with the Anderson heterojunction rule, the ZnO valence band offset at the p-ZnO–p-Si was close to 2.6 eV [17]. Because of this valence band offset, the contribution of holes in the p-Si to the forward current in the p-ZnO–p-Si MIS structure was negligible. Therefore, the forward current in these films was accompanied by electron-hole generation at the p-ZnO–p-Si interface.

It is known that acceptor dopants introduced into ZnO films are self-compensated by donor localized states generated in the ZnO lattice during the dopants' incorporation in ZnO [20,43]. From the results in Figure 3, we can conclude that during RTA at 550 ◦C, the concentration of donor-like defects decreased in the ZnO:N films in such a way that the concentration of nitrogen-acceptor states exceeded that of the remaining donor-like defects. As a result, the n-type as-deposited ZnO:N film was transformed to a p-type one during 550 ◦C RTA. The I–V characteristics of the same MIS structures, measured after a year, showed only minor changes and still exhibited p-type conductivity in these ZnO:N films. Therefore, the transformation of n-type ZnO film into p-type by 550 ◦C RTA treatment is a stable process.

The transformation of ZnO:N film from n-type to p-type by 550 ◦C RTA is similar to the appearance of p-type regions in N-implanted ZnO single crystals after annealing at 500 ◦C, observed by deep-level transient spectroscopy (DLTS) measurements [44]. Similar transformations of N-implanted n-type ZnO films to p-type ones have been observed after RTA treatment at 900 ◦C [45].

The current densities in the ZnO:N films deposited with nitrogen contents of 10 and 25% in the vacuum chamber were similar (Figure 3). Deposition of ZnO:N films at the highest nitrogen content of 40% resulted in lower current density, and it remained lower in comparison to the other ZnO:N films even after RTA treatment at 550 ◦C. This result is a consequence of already mentioned self-compensation of acceptor dopants by the simultaneously created donor-like deep levels during the acceptor doping of ZnO.

In ZnO films doped with nitrogen, the n-type conduction can be attributed to donors formed by interstitial zinc, Zni , and oxygen vacancies, VO, whereas the p-type conduction can be related to the acceptors formed by zinc vacancies, VZn, oxygen interstitial, Oi , and nitrogen substituting oxygen in the O sublattice [18,46,47]. We have confirmed the presence of these kinds of defect states in ZnO:N films by the analysis of IRSE and FTIR spectra [25] and also by the Raman and PL spectroscopic results presented and discussed above in Sections 3.1 and 3.2.

The specific resistivity, ρ(V), values of the ZnO films were calculated from the I–V characteristics shown in Figure 3, in the forward direction with accumulation conditions at the ZnO–p-Si interface, using the relations Rdif = dV/dI and ρ(V) = RdifS/df , where Rdif is the differential resistance, df is the film thickness, and S is the Al-dot contact area. The obtained ρ(V) values for the studied ZnO:N films are summarized in Figure 4.

in Sections 3.1 and 3.2.

**Figure 4.** Specific resistivity ρ(V) versus applied forward voltage of ZnO:N films (deposited at different gas ratios (**a**) and after RTA at 400 °C (**b**) and 550 °C (**c**). The ρ values were calculated from the forward I–V curves shown in Figure 3. **Figure 4.** Specific resistivity ρ(V) versus applied forward voltage of ZnO:N films (deposited at different gas ratios (**a**) and after RTA at 400 ◦C (**b**) and 550 ◦C (**c**). The ρ values were calculated from the forward I–V curves shown in Figure 3.

[25] and also by the Raman and PL spectroscopic results presented and discussed above

obtained ρ(V) values for the studied ZnO:N films are summarized in Figure 4.

The specific resistivity, ρ(V), values of the ZnO films were calculated from the I–V characteristics shown in Figure 3, in the forward direction with accumulation conditions at the ZnO–p-Si interface, using the relations Rdif = dV/dI and ρ(V) = RdifS/df, where Rdif is the differential resistance, df is the film thickness, and S is the Al-dot contact area. The

For the as-deposited and 400 °C RTA n-ZnO:N films (Figure 4a,b), the specific resistivity decreased sharply by more than four orders of magnitude upon increasing the applied voltage from 1 to 8 V. After 550 °C RTA (Figure 4c), the ρ(V) values of the p-ZnO:N films decreased more gradually, as the reduction was approximately two orders of magnitude in the same voltage interval of 1–8 V, but as a whole, they remained above the range of 106 Ωcm. In comparison with the as-deposited ZnO:N films, the smaller change of the specific resistivity in the 550 °C RTA-treated films as a function of the applied voltage is connected with the smaller increase of the current density in the 550 °C RTA films.

In order to elucidate the role of deep levels in the ZnO:N energy gap, we further measured manually, pointbypoint, the I–V characteristics of the MIS structures with ZnO:N films. For these experiments, we chose the sample series 2.1–2.3 with maximal concentration of nitrogen in the ZnO film, expecting a larger effect of N doping on the electrical characteristics. The room-temperature I–V characteristics of the MIS structure with ZnO:N films deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (a) and after RTA at 550 °C (sample 2.3) (b) are presented in Figure 5. To reveal the participation of deep levels in the ZnO energy gap, For the as-deposited and 400 ◦C RTA n-ZnO:N films (Figure 4a,b), the specific resistivity decreased sharply by more than four orders of magnitude upon increasing the applied voltage from 1 to 8 V. After 550 ◦C RTA (Figure 4c), the ρ(V) values of the p-ZnO:N films decreased more gradually, as the reduction was approximately two orders of magnitude in the same voltage interval of 1–8 V, but as a whole, they remained above the range of 106 Ω cm. In comparison with the as-deposited ZnO:N films, the smaller change of the specific resistivity in the 550 ◦C RTA-treated films as a function of the applied voltage is connected with the smaller increase of the current density in the 550 ◦C RTA films.

two-stage I–V measurements were performed. For the as-deposited n-type ZnO films of sample 2.1 (Figure 5a), the initial stage was measured in 20–25min increments starting from zero up to the maximal applied voltage, Va, and the second, return stage was carried out immediately after the initial one, down from Va to zero also in 20–25min increments. In order to elucidate the role of deep levels in the ZnO:N energy gap, we further measured manually, pointbypoint, the I–V characteristics of the MIS structures with ZnO:N films. For these experiments, we chose the sample series 2.1–2.3 with maximal concentration of nitrogen in the ZnO film, expecting a larger effect of N doping on the electrical characteristics.

The corresponding I–V curves in the reverse direction are given as an inset in Figure 5a. The considerable asymmetry between the currents at positive and negative voltage values confirmed the formation of the n-ZnO:N–p-Si heterojunction. The forward current in the return stage was higher than the forward current in the initial stage, resulting in a counterclockwise hysteresis in the measured I–V curves. Some of the electrons injected into the ZnO film during the initial stage were captured at deep levels with time constants higher than the time needed for measuring the initial stage, and their release during the return stage led to the observed excess, higher forward current in the return stage. Thus, the hysteresis in the I–V characteristics indicates that deep levels in the ZnO energy gap take part in the charge transport mechanism through the n-ZnO:N films. The room-temperature I–V characteristics of the MIS structure with ZnO:N films deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (a) and after RTA at 550 ◦C (sample 2.3) (b) are presented in Figure 5. To reveal the participation of deep levels in the ZnO energy gap, two-stage I–V measurements were performed. For the as-deposited n-type ZnO films of sample 2.1 (Figure 5a), the initial stage was measured in 20–25min increments starting from zero up to the maximal applied voltage, Va, and the second, return stage was carried out immediately after the initial one, down from Va to zero also in 20–25min increments. The corresponding I–V curves in the reverse direction are given as an inset in Figure 5a. The considerable asymmetry between the currents at positive and negative voltage values confirmed the formation of the n-ZnO:N–p-Si heterojunction. The forward current in the return stage was higher than the forward current in the initial stage, resulting in a counterclockwise hysteresis in the measured I–V curves. Some of the electrons injected into the ZnO film during the initial stage were captured at deep levels with time constants higher than the time needed for measuring the initial stage, and their release during the return stage led to the observed excess, higher forward current in the return stage. Thus, the hysteresis in the I–V characteristics indicates that deep levels in the ZnO energy gap take part in the charge transport mechanism through the n-ZnO:N films.

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**Figure 5.** Room temperature I–V characteristics of MIS structures with ZnO:N films, deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 °C (sample 2.3) (**b**). The initial and return stages of the I–V measurements are denoted with empty and full triangles, respectively. The corresponding I–V curves in the reverse direction are given as insets. **Figure 5.** Room temperature I–V characteristics of MIS structures with ZnO:N films, deposited at Ar:O2 :N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 ◦C (sample 2.3) (**b**). The initial and return stages of the I–V measurements are denoted with empty and full triangles, respectively. The corresponding I–V curves in the reverse direction are given as insets. **Figure 5.** Room temperature I–V characteristics of MIS structures with ZnO:N films, deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 °C (sample 2.3) (**b**). The initial and return stages of the I–V measurements are denoted with empty and full triangles, respectively. The corresponding I–V curves in the reverse direction are given as insets.

sample 2.3

The forward I–V characteristics of the MIS structures with ZnO:N films annealed at 550 °C (sample 2.3) are shown in Figure 5b, withthe corresponding I–V curves in reverse direction are given as inset. As expected, the forward current was with negative voltages applied to the Al-dot contact on the ZnO:N surface. As shown in Figure 5b, the forward current in the return stage was also higher than the forward current in the initial stage, leading to the appearance of counterclockwise hysteresis in the forward I–V characteris-The forward I–V characteristics of the MIS structures with ZnO:N films annealed at 550 ◦C (sample 2.3) are shown in Figure 5b, withthe corresponding I–V curves in reverse direction are given as inset. As expected, the forward current was with negative voltages applied to the Al-dot contact on the ZnO:N surface. As shown in Figure 5b, the forward current in the return stage was also higher than the forward current in the initial stage, leading to the appearance of counterclockwise hysteresis in the forward I–V characteristics. This counterclockwise hysteresis shown in Figure 5b, as the one in Figure 5a, also confirms that the deep levels in the ZnO energy gap take part in the charge transport mechanism in the p-ZnO:N films. The forward I–V characteristics of the MIS structures with ZnO:N films annealed at 550 °C (sample 2.3) are shown in Figure 5b, withthe corresponding I–V curves in reverse direction are given as inset. As expected, the forward current was with negative voltages applied to the Al-dot contact on the ZnO:N surface. As shown in Figure 5b, the forward current in the return stage was also higher than the forward current in the initial stage, leading to the appearance of counterclockwise hysteresis in the forward I–V characteristics. This counterclockwise hysteresis shown in Figure 5b, as the one in Figure 5a, also confirms that the deep levels in the ZnO energy gap take part in the charge transport mechanism in the p-ZnO:N films.

tics. This counterclockwise hysteresis shown in Figure 5b, as the one in Figure 5a, also confirms that the deep levels in the ZnO energy gap take part in the charge transport mechanism in the p-ZnO:N films. The room-temperature dependences of ln(J) versus ln|V| in the forward direction and in the initial stage are presented in Figure 6. For the MIS structures with as-deposited ZnO:N films, the slope of the plots is about 8 in the dominant parts of the plots and, around 2 for the other parts. This means that the charge transport in these ZnO:N films is carried out via deep levels in the ZnO energy gap. This transport mechanism is described as trap charge limited current (TCLC) (see for example [48]) or trap-assisted space charge The room-temperature dependences of ln(J) versus ln|V| in the forward direction and in the initial stage are presented in Figure 6. For the MIS structures with as-deposited ZnO:N films, the slope of the plots is about 8 in the dominant parts of the plots and, around 2 for the other parts. This means that the charge transport in these ZnO:N films is carried out via deep levels in the ZnO energy gap. This transport mechanism is described as trap charge limited current (TCLC) (see for example [48]) or trap-assisted space charge limited current (see for example [49]). For the MIS structures with 550 ◦CRT-treated ZnO:N films, the slope of the lnJ versus ln|V| plot (Figure 6) is around 2.5 in the range of 2–18 V. Since this slope is also higher than 2, this means that the charge transport in the p-type ZnO:N is also carried out via deep levels in the ZnO energy gap. The room-temperature dependences of ln(J) versus ln|V| in the forward direction and in the initial stage are presented in Figure 6. For the MIS structures with as-deposited ZnO:N films, the slope of the plots is about 8 in the dominant parts of the plots and, around 2 for the other parts. This means that the charge transport in these ZnO:N films is carried out via deep levels in the ZnO energy gap. This transport mechanism is described as trap charge limited current (TCLC) (see for example [48]) or trap-assisted space charge limited current (see for example [49]). For the MIS structures with 550 °CRT-treated ZnO:N films, the slope of the lnJ versus ln|V| plot (Figure 6) is around 2.5 in the range of 2–18 V. Since this slope is also higher than 2, this means that the charge transport in the p-type ZnO:N is also carried out via deep levels in the ZnO energy gap.

and after RTA at 550 °C (sample 2.3).

**Figure 6.** Logarithm of room-temperature forward current density as a function of logarithm absolute voltage for the MIS structures with ZnO:N films, deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) and after RTA at 550 °C (sample 2.3). **Figure 6.** Logarithm of room-temperature forward current density as a function of logarithm absolute voltage for the MIS structures with ZnO:N films, deposited at Ar:O2 :N2 = 50:10:40 (sample 2.1) and after RTA at 550 ◦C (sample 2.3).

The AC conductance, Gm, of these ZnO:N films, measured in the 0.5–500 kHz test frequency range, is plotted in Figure 7 as lnGm against lnω, where ω is the angular

**Figure 6.** Logarithm of room-temperature forward current density as a function of logarithm absolute voltage for the MIS structures with ZnO:N films, deposited at Ar:O2:N2 = 50:10:40 (sample 2.1)

The AC conductance, Gm, of these ZnO:N films, measured in the 0.5–500 kHz test frequency range, is plotted in Figure 7 as lnGm against lnω, where ω is the angular

The AC conductance, Gm, of these ZnO:N films, measured in the 0.5–500 kHz test frequency range, is plotted in Figure 7 as lnGm against lnω, where ω is the angular frequency. The slope of the plot for the as-deposited n-ZnO:N film (sample 2.1) is equal to 0.66. This value is close to the most widespread value 0.7 of the power exponent in the Jonscher universal power law for AC conductance in the case of hopping or tunneling of charge carriers via deep levels near to the Fermi level in the semiconductor energy gap [50]. For the annealed p-ZnO:N films (sample 2.3), the slope of this plot is equal to 0.24. frequency. The slope of the plot for the as-deposited n-ZnO:N film (sample 2.1) is equal to 0.66. This value is close to the most widespread value 0.7 of the power exponent in the Jonscher universal power law for AC conductance in the case of hopping or tunneling of charge carriers via deep levels near to the Fermi level in the semiconductor energy gap [50]. For the annealed p-ZnO:N films (sample 2.3), the slope of this plot is equal to 0.24. *Nanomaterials* **2022**, *12*, x FOR PEER REVIEW 12 of 18 frequency. The slope of the plot for the as-deposited n-ZnO:N film (sample 2.1) is equal to 0.66. This value is close to the most widespread value 0.7 of the power exponent in the Jonscher universal power law for AC conductance in the case of hopping or tunneling of

charge carriers via deep levels near to the Fermi level in the semiconductor energy gap

**Figure 7.** Logarithm of AC conductance, Gm, versus logarithm of angular frequency, ω, of MIS structures with ZnO:N films deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) and after RTA at 550 °C (sample 2.3). **Figure 7.** Logarithm of AC conductance, Gm, versus logarithm of angular frequency, ω, of MIS structures with ZnO:N films deposited at Ar:O2 :N2 = 50:10:40 (sample 2.1) and after RTA at 550 ◦C (sample 2.3). **Figure 7.** Logarithm of AC conductance, Gm, versus logarithm of angular frequency, ω, of MIS structures with ZnO:N films deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) and after RTA at 550 °C (sample 2.3).

In order to reveal the character of t**he** charge transport mechanism through the ZnO:N film, current–voltage measurements must be done at different temperatures. In Figure 8a, the I–V characteristics of the MIS structure with as-deposited n-ZnO:N, measured at 298 and 77 K, are given as the logarithm of the forward current density versus applied voltage. The averaged current density measured at 298 K under high applied voltages was 2.53 times higher than the current density measured at 77 K. Using the expression ln(J298/J77) = (qφa/k) (1/77–1/298), the effective activation energy, qφa, of these current densities was estimated. For the as-deposited n-ZnO:N film, it was equal to qφa = 8.3 meV. This small qφa value indicates that the current in this film was carried out predominantly by electron tunneling from the occupied deep level to the nearest unoccupied one in the In order to reveal the character of the charge transport mechanism through the ZnO:N film, current–voltage measurements must be done at different temperatures. In Figure 8a, the I–V characteristics of the MIS structure with as-deposited n-ZnO:N, measured at 298 and 77 K, are given as the logarithm of the forward current density versus applied voltage. The averaged current density measured at 298 K under high applied voltages was 2.53 times higher than the current density measured at 77 K. Using the expression ln(J298/J77) = (qϕa/k) (1/77–1/298), the effective activation energy, qϕa, of these current densities was estimated. For the as-deposited n-ZnO:N film, it was equal to qϕa = 8.3 meV. This small qϕa value indicates that the current in this film was carried out predominantly by electron tunneling from the occupied deep level to the nearest unoccupied one in the ZnO energy gap. The current at 77 K was carried out by intertrap tunneling mechanism [43]. The measured excess current at 298 K is a consequence of the appearance of thermally activated carrier hopping in these ZnO:N films at temperatures higher than 77 K [51]. In order to reveal the character of t**he** charge transport mechanism through the ZnO:N film, current–voltage measurements must be done at different temperatures. In Figure 8a, the I–V characteristics of the MIS structure with as-deposited n-ZnO:N, measured at 298 and 77 K, are given as the logarithm of the forward current density versus applied voltage. The averaged current density measured at 298 K under high applied voltages was 2.53 times higher than the current density measured at 77 K. Using the expression ln(J298/J77) = (qφa/k) (1/77–1/298), the effective activation energy, qφa, of these current densities was estimated. For the as-deposited n-ZnO:N film, it was equal to qφa = 8.3 meV. This small qφa value indicates that the current in this film was carried out predominantly by electron tunneling from the occupied deep level to the nearest unoccupied one in the ZnO energy gap. The current at 77 K was carried out by intertrap tunneling mechanism [43]. The measured excess current at 298 K is a consequence of the appearance of thermally activated carrier hopping in these ZnO:N films at temperatures higher than 77 K [51].

In the case of intertrap tunneling, the current density J is given by

**Figure 8.** lnJ versus V characteristics, measured at 298 and 77 K, of MIS structures with ZnO:N films deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 °C (sample 2.3) (**b**). **Figure 8.** lnJ versus V characteristics, measured at 298 and 77 K, of MIS structures with ZnO:N films deposited at Ar:O2 :N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 ◦C (sample 2.3) (**b**).

**Figure 8.** lnJ versus V characteristics, measured at 298 and 77 K, of MIS structures with ZnO:N films

J = J0 sinh[B(V − Vfb)] (1)

J = J0 sinh[B(V − Vfb)] (1)

deposited at Ar:O2:N2 = 50:10:40 (sample 2.1) (**a**) and after RTA at 550 °C (sample 2.3) (**b**).

In the case of intertrap tunneling, the current density J is given by

In the case of intertrap tunneling, the current density J is given by

$$\mathbf{J} = \mathbf{J}\_0 \sinh[\mathbf{B}(\mathbf{V} - \mathbf{V}\_{\text{fb}})] \tag{1}$$

where J0 = 2qν{exp[−2(2m\*q)1/2ϕt 1/2w/*h*¯]}/w2 and B = (2m\*q)1/2w2/*h*¯ϕt 1/2df [52]; w is the mean distance between deep levels situated around the quasi-Fermi level; ν is the electron attempt to escape frequency from deep levels; and qϕt is the energy position of deep levels in the ZnO energy gap. The electron effective mass, mn\* for n-type ZnO is taken as equal to 0.23 me [53]. The electron attempts to escape frequency, ν = 8.54 × 1012 s −1 , is estimated from the relation hν = kTD [54], where the Debye temperature, TD, for wurtzite ZnO is 410 K [55], and all other symbols have their common meaning. From Equation (1) it follows that the values of w and qϕt can be determined from the slope of the plot lnJ versus V and the intersection of the extension of this plot toward the lnJ axis at V = 0. These values of w and qϕt determined from the lnJ-V plot in Figure 8a at 77 K were 4.71 × 10−7 cm and 1.48 eV, respectively. The density of deep levels Nt in the ZnO energy gap was estimated by the expression Nt = 1/w3 and it was equal to Nt = 9.62 × 1018 cm−3 . Because of the energy position of deep levels is qϕt = 1.48 eV, the position of electron quasi-Fermi level in these n-ZnO MIS structures was 0.12 eV above the middle of the ZnO energy gap.

When the concentration of deep levels is known, one may calculate the mobility µt for this intertrap tunneling charge transport by the relation µt = J/(qNtE), where J is the current density at the applied voltage V and E is the electric field across the ZnO film. In the accumulation conditions at the ZnO/Si interface, E is given by V/df , where df is the ZnO film thickness. The calculated low value of µt = 1.04 × 10−6 cm2V −1 s −1 is characteristic for charge transport via deep levels in the semiconductor band gap.

The excess of the current at 298 K over that at 77 K (Figure 8a) is evidence for thermally activated carrier hopping in these n-ZnO MIS structures. N.F. Mott [51] has proposed that in the case of variable range hopping (VRH), the temperature-dependent conductivity can be expressed by σ(T) = σ0exp{−(T0/T)1/4}, where σ is the specific conductivity and T0 is the characteristic temperature. This dependence is appropriate for the Mott VRH in ZnO [56–58]. One may calculate the value of characteristic temperature T0 by using the expression ln(J298/J77) = T0 1/4[1/(77)1/4 − 1/(298)1/4], where the current density is averaged over high applied voltages. For the as-deposited n-ZnO MIS structures (Figure 8a) this value was T0 = 8.45 × 103 K.

The localization length, Bohr radius, aB,ZnO, of the localized states in the ZnO energy gap is given by the expression aB,ZnO = (4πεZnO*h*¯ 2 )/(m\*q2 ), where εZnO is the dielectric constant of ZnO, m\* is the effective electron mass in ZnO, and other symbols have their usual meanings. In the corresponding 1 MHz C-V curve of sample 2.1 (not shown in this paper), the measured maximal capacitance value in the accumulation regime was equal to Cmax = 98 pF. As Cmax is expressed by Cmax = εε0S/df , the dielectric constant εZnO of the ZnO:N film was determined as εZnO = εε0 = 1.03 × 10−12 F/cm. Knowing εZnO and mn\* = 0.23 me [53] quantities, the calculated value of aB,ZnO was equal to aB,ZnO = 2.674 × 10−7 cm.

The density of localized states, N(ε), in the case of hopping mechanism is given by the expression N(ε) = 16α 3/kT0, where the decay constant α is the inverse of the Bohr radius (α = 1/aB,ZnO). The numerical factor 16 is taken from Ambegaokar et al. [59]. With α = 3.74 × 106 cm−1 and T0 = 8.449 × 103 K the calculated value of N(ε) was equal to 1.15 × 1021 cm−3 .eV−1 . When N(ε) is known, the most probable hopping distance Rm can be estimated by the expression Rm = (9/8παN(ε)kT)1/4 [60,61]. At the temperature of 187.5 K, taken as the middle value of temperature range of 77–298 K, Rm was calculated, and it was equal to Rm = 2.677 × 10−7 cm. This Rm value is practically equal to the localization radius aB,ZnO = 2.67 × 10−7 cm, and it is 14 % higher than the half-distance w = 2.35 × 10−7 cm between adjacent deep levels obtained from the I–V measurement at 77 K. The most probable energy difference of deep levels, ∆W, taking part in the electron hopping is given by ∆W = (3/4πRm 3N(ε)) [60,61]. The calculated value of ∆W = 10.78 meV for this hopping in the n-ZnO–p-Si MIS structures was close to the effective thermally

activation energy qϕa = 8.3 meV in the as-deposited n-ZnO films. Therefore, in addition to temperature-independent intertrap tunneling, thermally activated hopping also took place in these n-type ZnO films in the temperature range of 77–298 K.

In Figure 8b, the forward I–V characteristics, measured at 298 and 77 K, of the MIS structures with ZnO:N films after RTA at 550 ◦C (sample 2.3) are given as the logarithm of the current density versus applied voltage. The averaged current density measured at 298 K under high applied voltage was 17.21 times higher than the current density measured at 77 K. The effective activation energy of the current densityof this p-ZnO:N film, obtained from the expression ln(J298/J77) = (qϕa/k)(1/77 − 1/298) was equal to qϕa = 25.46 meV. This value shows that the current in these p-ZnO:N films was also carried out predominantly by tunneling of holes from the occupied deep level to the nearest unoccupied one in the ZnO:N energy gap. Because of that, at 77 K the current density J is also given by the equation (1), where the electron effective mass, mn\* = 0.23 me, is replaced with the hole effective mass, mh\* = 0.59 me [62].

The values of qϕt and w, calculated from the slope of the plot lnJ versus V in Figure 8b and the intersection of the extension of this plot toward lnJ axis at V = 0, were qϕt = 1.27 eV and w = 3.724 × 10−7 cm, respectively. Under high forward bias the concentration of injected holes in the annealed at 550 ◦C p-ZnO film was equal to Nt = 1/w3 = 1.94 × 1019 cm−3 . Taking into account that our previous EDX studies have shown a nitrogen/oxygen ratio of 1/8 in both RF-sputtered and annealed ZnO:N films [25], the total concentration of nitrogen in the 550 ◦C RTA p-ZnO film (sample 2.3) was 1.82 × 1021 cm−3 . The comparison between the concentration of acceptor deep levels (Nt = 1.94 × 1019 cm−3 ) with the total concentration of N atoms incorporated into ZnO:N films (1.82 × 1021 cm−3 ) shows that less than 1% of acceptor levels were not compensated. Because of this strong charge compensation, it is difficult to achieve an increase in the conductivity of ZnO by introducing a larger amount of nitrogen into a ZnO:N film.

Assuming that temperature-dependent current in these p-type ZnO:N films is also governed by the Mott VRH [51], the characteristic temperature T0 obtained from the relation ln(J298/J77) = T0 1/4[1/(77)1/4 − 1/(298)1/4] was equal to T0 = 7.44 × 105 K. The values of the Bohr radius aB,ZnO, the decay constant α, the density of localized states, N(ε), the most probable hopping distance, Rm, and the most probable hopping energy difference, ∆W, were calculated in the same way as the corresponding parameters for the n-type ZnO:N films. The obtained values were, respectively, aB,ZnO = 1.042 × 10−7 cm, α = 9.59 × 106 cm−1 , and N(ε) = 2.2 × 1020 cm−3 .eV−1 , and at 187.5 K the Rm and ∆W values were 3.2 × 10−7 cm and 33.07 meV, respectively. These results reveal that in the p-type ZnO films, the most probable hopping distance Rm was about three times larger than the Bohr radius aB,ZnO and this relation confirms that Mott variable range hopping of holes occurred in these p-type ZnO:N films. Therefore, the additional current at 298 K compared to that at 77 K in these p-type ZnO:N films was carried out by charge carrier hopping, close to the Fermi or the quasi-Fermi level in the ZnO:N energy gap.

For better insight, the above calculated charge transport parameters of the intertrap tunneling and variable range hopping mechanism in ZnO:N films deposited at a gas ratio of Ar:O2:N2 = 50:10:40 (sample 2.1) and treated by RTA at 550 ◦C (sample 2.1) are summarized in Table 4.

**Table 4.** Intertrap tunneling and variable range hopping parameters of ZnO:N films deposited at gas ratio of Ar:O2 :N2 = 50:10:40 (sample 2.1) and treated by RTA at 550 ◦C (sample 2.1): qϕa—effective thermally activation energy; w—mean distance between deep levels; qϕt and Nt—energy position and density of deep levels, respectively, in the ZnO energy gap; T0—characteristic temperature; aB,ZnO—Bohr radius of the localized states in the ZnO energy gap; N(ε)—density of localized states; Rm and ∆W—most probable hopping distance and energy difference of deep levels, respectively.

The decrease of the ZnO band gap density of localized states from N(ε) = 1.15 × 1021 cm−3 eV−1 for the as-deposited ZnO:N (sample 2.1) to N(ε) = 2.2 × 1020 cm−3 eV−1 for the p-ZnO:N after 550 ◦C RTA (sample 2.3) correlates well with the observed decrease in the slope values of lnGm versus lnω plots (Figure 7) from 0.66 to 0.24, respectively. The decrease of the density of localized states in the ZnO bandgap by the 550 ◦C RTA is evidence that the concentration of donor-like defects decreased during this treatment. The densities of localized states N(ε) given in Table 4 are within the range of 9.93 × 1019–2.08 × 1022 cm−3 eV−1 reported for VRH conduction in n-type polycrystalline ZnO films [63]. Values of N(ε) = 2×1020 cm−3 eV−1 and ∆*W* = 29 meV, close to ours, for the densities of localized states and corresponding characteristic energies, respectively, have been obtained from the current–voltage characteristics of amorphous ZnON thin film transistors [64].

The established transformation of n-type conduction to p-type in the studied ZnO:N films viaannealing the samples at 550 ◦C suggests that if the deposition could be accomplished at such elevated temperatures as 550 ◦C, the density of localized states could be reduced below N(ε) = 2.2 × 1020 cm−3 .eV−1 , as was obtained in the 550 ◦C RTA-treated ZnO:N films. In this way, it would be possible to obtain nitrogen-doped ZnO films with weaker N-acceptor self-compensation and a smaller density of localized states in the p-type ZnO band gap than that we observed herein.

#### **4. Conclusions**

Nitrogen-doped ZnO thin films were deposited onto Si substrates by RF magnetron sputtering in nitrogen-containing ambient gas, followed by rapid thermal annealing at 400 ◦C and 550 ◦C in nitrogen. Raman and PL spectroscopic analyses revealed defects due to O and Zn vacancies and interstitials, which are related to carrier traps in the ZnO:N films.

Detailed analyses of the I–V characteristics of the MIS structures with the ZnO:N films revealedn-type conduction in both as-deposited and annealed at 400 ◦C ZnO:N films. The reason for n-type conductivity in these ZnO:N films is a result of the self-compensation of the N-acceptor levels by donor-like defects generated during N doping. RTA at 550 ◦C decrease the concentration of donor-like defects below the concentration of N acceptors and, as a consequence, the n-type ZnO:N films transformed into p-type ZnO:N ones. Repeated measurements after a year proved that the resulting conduction-type transformation at the annealing stage at 550 ◦C is a stable process.

The observed counterclockwise hysteresis in the measured point-by-point I–V characteristics of both n- and p type ZnO:N films confirms the presence of deep levels in the ZnO energy gap which participate in the charge transport through these ZnO:N films. The slope of the lnJ versus lnV plots reveals that the current in both n- and p-type ZnO:N films is trap charge limited current (TCLC), while the slope of the lnGm versus lnω plots confirms the tunneling or hopping of charge carriers via deep levels in the ZnO energy gap.

At 77 K, the current in both n- and p-type ZnO:N films is carried out by intertrap tunneling via deep levels in the ZnO energy gap. The higher forward averaged current density at 298 K than that measured at 77 K reveals additional thermally activated variable range carrier hopping in these ZnO:N films.

**Author Contributions:** Formal analysis, investigation, writing-original draft preparation, S.S. and M.N.; conceptualization, investigation, writing—original draft preparation, writing-review and editing, A.S.; formal analysis, investigation, D.S. and I.S., formal analysis, writing—review and editing M.A.; visualization, E.A. and M.M.; conceptualization, writing-review and editing, data curation, project administration, funding acquisition, M.G. All authors have read and agreed to the published version of the manuscript.