**Ethylene Glycol Functionalized Gadolinium Oxide Nanoparticles as a Potential Electrochemical Sensing Platform for Hydrazine and p-Nitrophenol**

**Savita Chaudhary 1,\*, Sandeep Kumar 1, Sushil Kumar 1, Ganga Ram Chaudhary 1, S.K. Mehta <sup>1</sup> and Ahmad Umar 2,3,\***


Received: 20 July 2019; Accepted: 26 September 2019; Published: 1 October 2019

**Abstract:** The current work reports the successful synthesis of ethylene glycol functionalized gadolinium oxide nanoparticles (Gd2O3 Nps) as a proficient electrocatalytic material for the detection of hydrazine and p-nitrophenol. A facile hydrothermal approach was used for the controlled growth of Gd2O3 Nps in the presence of ethylene glycol (EG) as a structure-controlling and hydrophilic coating source. The prepared material was characterized by several techniques in order to examine the structural, morphological, optical, photoluminescence, and sensing properties. The thermal stability, resistance toward corrosion, and decreased tendency toward photobleaching made Gd2O3 nanoparticles a good candidate for the electrochemical sensing of p-nitrophenol and hydrazine by using cyclic voltammetric (CV) and amperometric methods at a neutral pH range. The modified electrode possesses a linear range of 1 to 10 μM with a low detection limit of 1.527 and 0.704 μM for p-nitrophenol and hydrazine, respectively. The sensitivity, selectivity, repeatability, recyclability, linear range, detection limit, and applicability in real water samples made Gd2O3 Nps a favorable nanomaterial for the rapid and effectual scrutiny of harmful environmental pollutants.

**Keywords:** gadolinium oxide; hydrazine; p-nitrophenol; electrochemical sensing; amperometric; selective sensor

#### **1. Introduction**

Aromatic nitro as well as hydrazine are some of the few compounds that are frequently used in the preparation of insecticides, pesticides, pharmaceuticals, and in chemical industries [1–3]. The highly stable nature and lower degradation efficiency of these compounds imparted serious health hazards to human health [4,5]. The utilities of these chemicals in the preparation of explosives are well established in the literature [6]. For instance, according to the U.S Homeland Security Information Bulletin, hydrazine was used in a terrorist attack in 2003 [7]. Hence, from the perspective of safety and security, the development of simple, handy, and competent methodology to monitor these contaminants is the crucial need of the society [8]. To date, there have been a number of analytical instrumental techniques—such as colorimetric, X-ray, fluorescence emission spectroscopy, inductively coupled plasma spectroscopy, atomic absorption spectroscopy, mass spectroscopy, and chromatography—that have been employed for the detection of p-nitrophenol as well as hydrazine [9–12].

All these methods are quite efficient for the detection of these pollutants, but possess delicate functioning, high processing charges, and skilled professionals for data analysis [12]. All these hitches

have restricted the use of these sophisticated techniques in routine applications. Hence, from the viewpoint of human health and environmental security, there is a critical requirement for developing alternative techniques with improved selectivity and sensitivity toward these contaminants [13–15]. Therefore, in this work, we have coupled the sensitivity of electrochemical technique with nanoparticles for developing effective sensors for these harmful pollutants. The developed electrochemical sensor has offered significant benefits such as low processing cost and quick response time. The presence of nanoparticles has further augmented the mass transport during analysis as well as reduced the effect of opposition produced by solution during the measurements. The signal-to-noise ratio is further enhanced in the presence of nanoparticles as compared to conventional macroelectrodes used during the analysis. The higher available surface area of nanoparticles has made them an efficient adsorbent for analytes and provided better-quality responses for contaminants.

In the past, varieties of nanoparticles have been used for preparing electrochemical sensors for hydrazine and aromatic compounds [15–20]. For instance, Mishra et al. [15,16] developed the flexible epidermal tattoo, textile and glove-based electrochemical sensor for the detection of organophosphate molecules. The developed sensor was found to be effective in defense and food security applications. Wei et al. [17] have used the nanohybrids of carbon nanotubes (CNTs) with pyrene-cyclodextrin for the electrochemical sensing of p-nitrophenol with sensitivity of around 18.7 μA/μM. Karthik et al. [18] have used the applications of gold nanoparticles derived from biogenic sources for sensing hydrazine from aqueous media. The developed sensor has shown a linear range 5 nM to 272 μM with a detection limit of around 0.05 μM. Zhang and co-workers have used the application of modified graphene with Pt–Pd nanocubes for detecting aromatic nitro compounds with a detection range of 0.01 to 3 ppm and sensing limit of around 0.8 ppb [19]. Chen et al. [20] have developed the indium tin oxide electrodes, which were further functionalized with β-cyclodextrin and Ag nanoparticles for analyzing the trace amount of nitroaromatic compounds via using electrochemical sensing analysis. Chaudhary et al. [21] have utilized the fluorescence sensing abilities of Gd2O3 Nps nanoparticles in the selective detection of 4-nitrophenol in aqueous media. However, the use of Gd2O3 Nps for the modification of electrodes was less frequent in the literature for the estimation of harmful hydrazine and aromatic compounds.

The current work has utilized the electron transport abilities, high electrical conductivity, and thermal stabilities of Gd2O3 Nps for making effective material in electrochemical sensing for harmful pollutants [21–23]. To date, a diverse range of methodologies has been investigated for the preparation of Gd2O3 Nps [24–28]. The available methods have generally required very high temperature reaction conditions and multistep processing for the synthesis of Gd2O3 Nps. Therefore, it is valuable to systematize a synthetic approach for preparing Gd2O3 Nps at comparatively low temperature in a minimum number of steps, and formed particles will be reliable for developing an effective sensor for harmful pollutants. The present study has emphasized a hydrothermal route for the preparation of Gd2O3 Nps under mild conditions. The hydrothermal process is the best preference due to its superior competence, economical nature, flexible reaction constraints, and prospective ability for the large-scale production of particles. The used methodology has provided better control over the size and shape of the formed particles. Abdullah et al. [29] have reported the fabrication of well crystalline Gd2O3 nanostructures by annealing the hydrothermally prepared nanostructures at 1000 ◦C. The prepared particles were further used for the detection of ethanol. In the current study, the biocompatible coating of ethylene glycol (EG) has provided better control over the agglomeration rate of formed nanoparticles. The presence of EG has a direct influence over the solubility, optical, luminescence, and the morphological characteristics of the prepared nanoparticles. The external template of EG has also modulated the range of non-radiative energy losses in Gd2O particles.

Here in this work, surface-modified gadolinium oxide (Gd2O3) nanoparticles have been used as a proficient electrocatalytic material for the detection of hydrazine and p-nitrophenol by using cyclic voltammetric (CV) and chronoampherometric methods at neutral pH range. The consequences of synthetic parameters such as the concentration of precursors were studied by measuring the optical, photoluminescence, and band-gap variation of the formed particles. The estimation of the sensitivity, selectivity, repeatability, recyclability, linear range, detection limit, and applicability in real water has also been carried out in the current work. The thermal stability, resistance toward corrosion, and decreased tendency toward photobleaching made Gd2O3 nanoparticles a probable contender for preparing a simple, fast, and economical electrochemical sensor for hydrazine and p-nitrophenol.

#### **2. Experimental Details**

#### *2.1. Materials*

GdCl3·6H2O (Gadolinium(III) chloride hexahydrate: Sigma Aldrich, Mumbai, India with purity 99%) was used as a starting material for fabricating Gd2O3 nanoparticles. EG (ethylene glycol: Fluka 98%) and NaOH (sodium hydroxide: Merck 99.9% pure) were used for the synthesis purpose. Hydrazine, p-nitrophenol, benzaldehyde, benzoic acid, benzonitrile, phenol, ethanol, and aniline were procured from Sigma Aldrich, with purity more than 90%. Acetone (BDH, Mumbai, India, 98%) and ethanol (Changshu Yangyuan, Suzhou, China, 99.9%) were used as the washing solvent for obtained nanoparticles. Millipore distilled water was used for the synthesis of nanoparticles.

#### *2.2. Synthesis of Gd2O3 Nanoparticles*

The synthesis of Gd2O3 nanoparticles was done by using the hydrothermal method. In brief, 5 mM of GdCl3·6H2O was added to the 5-mL EG solution under stirring at 50 ◦C solution followed by the addition of 15 mM NaOH. The temperature of the reaction mixture was held constant at 140 ◦C for the first hour, and then raised to 180 ◦C for 4 h. The obtained solution was allowed to cool at room temperature. The obtained yellow precipitates of Gd2O3 nanoparticles were separated out from the reaction media. The obtained precipitates were subjected to calcinations for 3 h at 300 ◦C. The resulting particles were washed with water, acetone, and ethanol to remove the impurities. The corresponding separation was mainly performed by ultracentrifugation at 9000 rpm. The extracted particles were dried in an oven at 50 ◦C and further utilized for different analysis. For the optimization of synthetic parameters for the preparation of Gd2O3 nanoparticles, the respective concentration variations of GdCl3·6H2O have been carried out under respective reaction conditions. In the first instance, the concentration of variations of GdCl3·6H2O were done from 5 to 25 mM in all the reaction mixtures by keeping the concentration of the NaOH fixed at 0.015 M. A UV-visible spectral scan for each sample was taken from 230 to 400-nm wavelength to detect the optical properties of the formed nanoparticles. The optical band gap (Eg) of as-prepared nanoparticles was calculated as a function of the concentration variations of GdCl3·6H2O. The respective fluorescence emission spectra were also studied for the concentration variations of GdCl3·6H2O from 5 to 25 mM.

#### *2.3. Electrode Preparation*

The synthesized nanoparticles were further used to fabricate the electrochemical sensor for hydrazine and p-nitrophenol (Scheme 1). In order to form the modified electrode, the gold electrode with a surface area equivalent to 3.14 mm<sup>2</sup> was first cleaned with alumina slurry and properly washed with distilled water under sonication. After drying the electrode at room temperature, the surface of the gold electrode was coated with the Gd2O3 nanoparticles by using butyl carbitol acetate (BCA) as the binding agent. The as-formed electrode was further dried at 60 ± 5 ◦C for 4–6 h to attain a homogeneous and dried layer of nanoparticles over the surface of electrode. All the electrochemical measurements were carried out on an μAutolab Type-III under neutral pH conditions. In all the analyses, a Gd2O3 customized gold electrode was acting as a working electrode, Ag/AgCl (sat. KCl) was acting as a reference, and Pt wire was acting as a counter electrode.

**Scheme 1.** Schematic illustration of the electrochemical sensor for the detection of hydrazine and p-nitrophenol by using the modified electrode of gold with Gd2O3 nanoparticles.

#### *2.4. Physical Measurements*

The obtained Gd2O3 nanoparticles were characterized with the help of an X-ray diffractometer from Panalytical D/Max-2500 (Malvern, UK), Hitachi (H-7500) Transmission electron microscope (Tokyo, Japan), Thermoscientific UV-vis. Spectrophotometer (Waltham, MA, USA), and FTIR spectrophotometer of Perkin-Elmer (RX1) (Waltham, MA, USA). The fluorescence measurements were carried out with a Hitachi F-7000 photoluminescence spectrophotometer. The photoluminescence analysis was carried out on an Edinburgh Instrument FLS 980 (Bain Square, UK). A JEOL (JSM-6610) scanning electron microscope (SEM) (Tokyo, Japan) with EDX analysis was carried out at 20 kV. The calcinations of the as-prepared Gd2O3 nanoparticles were done in an AICIL muffle furnace at 300 ◦C. Dynamic light scattering measurements were done on a Malvern Zen1690 instrument (Worcestershire, UK). Raman analysis was performed on a Renishaw inVia reflex micro-Raman spectrometer (Wotton-under-Edge, UK). The pH measurements were performed on a Mettler Toledo digital pH meter (Columbus, OH, USA). The surface area of EG-coated Gd2O3 nanoparticles was estimated by using Brunauer–Emmett–Teller (BET) analysis with an N2 adsorption analyzer (NOVA 2000e, Anton Par, Gurugram, India). The separation of the as-prepared Gd2O3 nanoparticles from aqueous media was done on a Remi R-24 centrifuge. Electrochemical measurements were carried out on an μAutolab Type-III cyclic voltammeter (Metrohm, Herisau, Switzerland). The gold electrode with a surface area of 3.14 mm2 was chosen for the analysis.

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

#### *3.1. Characterization and Properties of Synthesized Gd2O3 Nanoparticles*

The crystal structure of formed Gd2O3 nanoparticles has been further scrutinized by investigating the powdered XRD patterns of formed particles (Figure 1a). The absence of any peak related to impurity has confirmed the purity of the prepared nanoparticles [30]. The average crystallite size (D) of 15 nm has been estimated from the diffraction peaks by using the respective values of full width at half maximum (FWHM) via employing Debye–Scherrer's formula [31,32]. The specific surface area and pore diameter of the obtained sample was found to be 15.3 m2·g−<sup>1</sup> and 2.3 nm, which were respectively calculated by using the nitrogen sorption studies at 77 K in accordance with the BET (Brunauer–Emmett–Teller) process. The respective atomic content of Gd2O3 nanoparticles has been confirmed by using EDX analysis (Figure 1b). The obtained spectrum has only displayed the characteristic peaks of the Gd and O atoms in the synthesized sample, which verified the purity of the formed particles.

**Figure 1.** Typical (**a**) XRD pattern, (**b**) EDS spectrum, (**c**) SEM image, (**d**) TEM image, and (**e**,**f**) HRTEM images of Gd2O3 nanoparticles.

The SEM image of the Gd2O3 nanoparticles has been shown in Figure 1c, which has clearly shown the presence of agglomerated nanostructures of Gd2O3 nanoparticles. The presence of contacted particles has been mainly due to the existence of an EG template over the surface of the particles, which has further supported the presence of external electrostatic interactive forces generated from the templates over the surface of the nanoparticles. Detailed information regarding the morphology and structure of Gd2O3 nanoparticles has further been obtained from the HRTEM images presented in Figure 1d. It is clear from the images that the nanoparticles have shown crystallites with an irregular pseudo-spherical shape and a size distribution between 7–15 nm. The crystal spacing of 0.305 nm belongs to the (222) planes of Gd2O3 nanoparticles (Figure 1e,f). The observed result is in good agreement with the reported literature [33]. The presence of the connected nanocrystals has been attained due to the existence of a diverse range of forces (electrostatic, hydrogen bonding, and van der Waals forces) provided by the presence of external templates of EG coating over Gd2O3 nanoparticles. *Coatings* **2019**, *9*, 633

The optical properties of formed Gd2O3 nanoparticles have been investigated by using UV-vis. and fluorescence analysis as a function of variation of the concentration of GdCl3·6H2O salt during the synthesis (Figure 2a,b). The formed particles have shown the characteristic peak between 255–262 nm. The respective peak has been associated with the electronic transition from 8S7/2–6I7/2. [34]. On interpreting the results, it has been found that the absorbance is dependent on the concentration of salt.

**Figure 2.** (**a**) UV-vis., (**b**) fluorescence and photoluminescence (PL) (inset) emission, (**c**) variation of bandgap and agglomeration number, (**d**) Commission Internationale de L'Eclairage (CIE) chromaticity analysis (**e**) Raman and (**f**) FTIR spectra of Gd2O3 nanoparticles.

With increase in the concentration from 5 to 20 mM, there has an increment in the absorbance of around 63%. The change in the peak position was not so prominent with the concentration of the salt. On the other hand, a fluorescence emission peak was observed at 336 nm with λexc = 290 nm (Figure 2b). This peak has been associated with the emission from 6P7/2↔8S7/<sup>2</sup> in Gd(III) ions [35]. The enhancement in the concentration of the starting material has produced an increment of 50.7% in intensity value, whereas the peak position has only shown a change of 10 nm. This variation has displayed the similarity with UV-vis studies. The PL spectra of as-synthesized nanoparticles has shown the characteristic peaks at 390 nm and a broad peak between 400–500 nm with a center at 417 nm (λexc = 350 nm) (inset Figure 2b). The sharp peak at 390 nm has been associated with the radiative recombination of holes with electrons present at the oxygen vacant positions formed due to the photogeneration effect. The other peak was associated with the self-trapped exciton luminescence in formed particles [36]. The optical band-gap values (*E*g) and agglomeration number for Gd2O3 nanoparticles (Figure 2c) have been estimated by using the application of the Brus method [37]. The results have clearly explained the behavioral variation of band-gap value with the concentration of GdCl3·6H2O salt during the synthesis. The decrease in the value of the band gap with the concentration has been associated with the variation of size of the formed particles with the concentration of the

starting material. The particle size was comparatively higher for the nanoparticles prepared with 25 mM GdCl3·6H2O salt. These variations of starting material have directly influenced the size of the particles, and the respective agglomeration number of the particles has also varied in a similar manner. Figure 2d has displayed the respective assignment of their colors in the Commission Internationale de L'Eclairage (CIE) diagram for the Gd2O3 nanoparticles. The respective value of CIE chromaticity coordinates is found to be *x* = 0.3265, *y* = 0.4462, respectively. The outcomes have been associated with the green-shift effect in the formed particles. The formed particles have been further characterized by using the Raman and FTIR spectra of Gd2O3 nanoparticles (Figure 2e,f). On interpretation, Gd2O3 nanoparticles displayed four major Raman peaks at 444.1, 528, 637.8 and 767.4 cm−1, respectively (Figure 2e). These peaks are mainly associated with the Fg and Ag mode for cubic C-type Gd2O3 particles [38]. Moreover, other Raman active modes such as 4Ag, 4Eg, and 14Fg have also contributed toward the peaks in the spectra [39,40]. A sharp IR peak has been detected below 500 cm−1, which has been attributed to Gd–O [40]. The small peak at 1500 cm−<sup>1</sup> has been associated with the δ(O–H) vibrations due to the water bound to the nanoparticles surface in the form of moisture [41–44].

#### *3.2. Electrochemical Behaviour of Hydrazine and p-Nitrophenol on Modified Electrode*

The electrochemical action of formed Gd2O3 nanoparticles has been investigated toward the electrocatalysis of hydrazine and p-nitrophenol (PNP) via using cyclic voltammetric analysis. Figure 3 showed the cyclic voltammograms of a gold electrode in pH 7.0, phosphate buffer (PBS) under various electrode conditions. Interestingly, it has been found that the bare gold electrode does not show any signal in pH 7 PBS buffer. The response has been still negligible for bulk Gd2O3-coated gold electrodes in the presence of hydrazine and p-nitrophenol (Figure 3). On the other hand, well-defined voltammetric signals at 0.68 V have been obtained for the electrocatalysis of hydrazine in the presence of a Gd2O3 nanoparticles-coated gold electrode. In case of p-nitrophenol, one set of reversible redox peaks i.e., (GdR1), oxidation (GdO1) occurred at −0.021 and 0.163 V. Other irreversible reduction peaks (GdR2) at −0.694 V and (GdR3) at 0.4 V were also observed in phosphate buffer solution at pH 7 with a scan rate of 60 mV/s. The obtained results have clearly pointed out that the p-nitrophenol has displayed three types of electrochemical responses with two subsequent types of processes, including reduction and redox couple progression [21]. This might be aroused due to the two-electron oxidation reduction reaction in 4-aminophenol. Whereas, the reduction peaks were associated with the formation of a hydroxylamine group from the nitro group of p-nitrophenol.

**Figure 3.** Cyclic voltammograms for a Gd2O3 nanoparticles-coated gold electrode under various electrode conditions in 0.1 M PBS (pH 7.0). The scan rate was 60 mV/s and the respective concentrations of analyte was kept constant at 1 mM.

From Figure 3, it was found that there has been no signal in the reverse sweep for hydrazine samples. The results have confirmed the irreversible nature of the oxidation process for hydrazine. On other hand, the samples of p-nitrophenol have displayed the redox peaks in both forward and backward directions, which make the analysis of p-nitrophenol reversible in nature [45]. In the presence of Gd2O3 nanoparticles, there has been found a significant increase in the peak current for respective analytes. This has clearly verified the utilities of formed nanoparticles for electro-analytical purposes. The enhancement of peak current has been mainly explained by the enhancement of conductivity of the electrode due to the functionalization of the gold electrode surface with Gd2O3 nanoparticles. These results have further validated that the formed Gd2O3 nanoparticles are capable as efficient electron transporters for the electrocatalysis of harmful pollutants.

The scan rate variations have also been carried out in order to investigate the electron transfer mechanism for a Gd2O3 modified gold electrode in the presence of hydrazine and p-nitrophenol. Figure 4a,b shows the typical voltammograms for a Gd2O3 nanoparticles-coated gold electrode with 1 mM solution of respective analyte (i.e., hydrazine and p-nitrophenol) in 0.1 M PBS solution with pH = 7.0 at different scan rates. The spectrum has revealed a regular enhancement of current response with varying the scan rate from 60 to 900 mV/S for both the analytes. The higher surface area of Gd2O3 nanoparticles has played a critical role for the improved electron transfer process for the catalytic performance toward the understudied analytes. The enhancement of peak current with the scan rate has clearly pointed out the electrocatalytic reaction of hydrazine and p-nitrophenol at the surface of the modified electrode. These current variations have been mainly explained by the surface adsorption to diffusion processes at a façade of modified electrodes [45].

**Figure 4.** Cyclic voltammograms for a Gd2O3 nanoparticles-coated gold electrode for 1 mM (**a**) hydrazine and (**b**) p-nitrophenol at different scan rates ranging from 60 to 900 mV/s.

At lower scan rates, the relative rate of diffusion is quick, and the adsorption of analyte has been found to be slowest during the mass transfer process. On the other hand, at higher scan rates, the diffusion step has been mainly controlling the rate of electrode reactions in the presence of external agents [46]. In the case of p-nitrophenol, both types of associated processes i.e., the reduction and redox couple processes, have shown the significant augmentation of current response as a function of scan rate. The respective calibration curve of peak current versus the square root of the scan rate has displayed a linear relation for hydrazine and p-nitrophenol (Figure 5). These obtained results have clearly pointed out that the oxidation of hydrazine is mainly a diffusion-controlled process at the surface of a Gd2O3 nanoparticles-modified gold electrode [47].

**Figure 5.** The linear dependence of peak current versus square root of scan rate for 1 mM hydrazine and p-nitrophenol.

The number of electrons involved in the overall reaction (*n*) for hydrazine and p-nitrophenol has been calculated from the Randles–Sevcik equation mentioned below [45].

$$ip = \left(2.69 \times 10^5\right) n^{\frac{3}{2}} A D^{\frac{1}{2}} v^{\frac{1}{2}} C \tag{1}$$

where *n* is the number of electron equivalents exchanged during the redox process, *A* (cm2) is the active area of the working electrode, *<sup>D</sup>* (cm2·s<sup>−</sup>1) and *<sup>C</sup>* (mol·cm<sup>−</sup>3) are the diffusion coefficient and the bulk concentration of hydrazine, and *<sup>v</sup>* is the voltage scan rate (V·s<sup>−</sup>1). The obtained CV responses of Gd2O3/Au have clearly explained that the oxidation of N2H4 involves two electron changes, and the irreversible reaction of p-nitrophenol has undergone four electron changes, while the reversible reaction involves two processes—an electron redox process and an irreversible process—and gives a large reduction peak. (Figure 4b) shows the reduction of PNP to 4-(hydroxyamino) phenol. Two coupled redox peaks have indicated the oxidation of 4-(hydroxyamino) phenol to 4-nitrosophenol and the succeeding reversible reduction [48–51]. A schematic representation of the involved mechanism has been illustrated in Scheme 2.

**Scheme 2.** The pictorial representation and cyclic voltammetric sweep curves for the Gd2O3 nanoparticles/butyl carbitol acetate/gold (NPs/BCA/Au) electrode for the sensing of hydrazine and p-nitrophenol.

#### *3.3. Amperometric Responses for Hydrazine and p-Nitrophenol*

The amperometric studies are one of the primary techniques to estimate the low concentration of analytes and carry out the relative studies in the presence of interfering analytes. Since the Gd2O3@Au-modified electrode has displayed the higher current response for hydrazine and p-nitrophenol as a model system in the cyclic voltammetric studies, therefore, it has been employed as the amperometric sensor for the detection of hydrazine and p-nitrophenol at low concentration levels.

Figures 6 and 7 depict the amperometric response of Gd2O3@Au for the successive additions of hydrazine and p-nitrophenol at an applied potential of −0.694 and 0.640 V for p-nitrophenol and hydrazine, respectively, in 0.1 M buffer solution with pH = 7. The obtained values of the current response have been estimated after the consecutive injection of 1 μM concentration of hydrazine and p-nitrophenol at the time interval of 60 s in a continuously stirring condition. The Gd2O3@Au-modified electrode has exhibited a considerable and rapid amperometric reaction toward each addition of analyte.

**Figure 6.** (**a**) Amperometric response of the Gd2O3/Au electrode with an increase in the concentration of hydrazine and (**b**) respective peak current vs. concentration plot of hydrazine in PBS at pH = 7.

**Figure 7.** (**a**) Amperometric response of the Gd2O3/Au electrode with an increase in the concentration of p-nitrophenol and (**b**) respective peak current vs. concentration plot of p-nitrophenol in PBS at pH = 7.

The value of the current has reached its stable position within 3 s, demonstrating the fast electro-oxidation of the understudied analyte at the surface of the Gd2O3@Au-modified electrode. In addition, the response current has shown a linear increment for the subsequent additions of analyte over the wide range of concentrations. The respective regression plot of current response versus concentration of both the analytes has displayed a linear relation with the correlation coefficient values of 0.987 and 0.996 for p-nitrophenol and hydrazine, respectively (Figures 6b and 7b). The limit of detection value has been calculated to be 1.527 and 0.704 μM for p-nitrophenol and hydrazine, respectively, by using the equation limit of detection (LOD) = 3σ/slope, where σ is the standard deviation for the particular system [52]. The sensitivity of the developed sensor has been found to be 0.33722 and 0.25734 mA·mM−<sup>1</sup> from the slope of the linear regression for hydrazine and p-nitrophenol,

respectively. The respective reusability, stability, and reproducibility of the as-prepared sensor has also been tested in the current work. The as-developed electrode was kept in the buffer media for one month, and its electrocatalytic efficiency has also been tested against the p-nitrophenol and hydrazine. The results have clearly verified that the formed sensor has displayed a reproducible performance with a decay rate of 5.7% and 6.3% in the oxidation peak current value for p-nitrophenol and hydrazine, respectively. This substantial constancy in results has been further attributed to the stability of Gd2O3 particles, which maintain the efficiency and performance of the electrode for a long period. Additionally, the reproducibility of the developed sensor has been confirmed by estimating the electrochemical response of the Gd2O3 particles as a working electrode for different electrodes in the solution media containing 1 mM of p-nitrophenol and hydrazine.

The obtained relative standard deviation (RSD) of peak currents has been found to be 5.3% and 4.7%, signifying the satisfactory report for the reproducibility of the modified electrode. In order to investigate the reusability of the developed sensor, the as-modified electrode has been rinsed with the respective sample solution. The obtained signal has been tested after the rinsing. It has been found that the obtained signal has maintained 94% of its original strength. The analytical performance of the as-modified electrode has been evaluated with some recent works in Table 1.


**Table 1.** Comparison of detection limit and response time of different electrode materials.

The data has further confirmed the authenticity of the developed sensor in a different range of concentrations with greater selectivity and sensitivity. In order to investigate the application of the formed sensor in real samples, the water samples from different sources have been taken, and the respective analyses have been made for the detection of hydrazine and p-nitrophenol. The relevant stock solutions of understudied analytes of known concentrations were made by using the water samples taken from different sources. The current response of the sensor has been examined for the different concentrations of hydrazine and p-nitrophenol by using amperometric studies. The outcomes of the measurements have shown the excellent recovery rate for the chosen analytes, as indicated in Table 2.

The sensitivity, selectivity, repeatability, recyclability, wide linear range, detection limit, and applicability in real water samples makes Gd2O3 Nps a favorable nanomaterial in the institution of the rapid and effectual scrutiny of harmful environmental pollutants. Thus, the prepared sensor appears to be a probable contender for preparing a simple, fast, and economical electrochemical sensor.


**Table 2.** Determination of hydrazine and p-nitrophenol real water samples. RSD: relative standard deviation.

#### *3.4. Selectivity Study*

In order to apply the proposed sensor for the determination of p-nitrophenol and hydrazine in an aqueous system, the respective selectivity of the sensor has been investigated in the presence of 1 mM of various interfering compounds (benzaldehyde, benzoic acid, benzonitrile, phenol, ethanol, and aniline) at a fixed potential 0.67 V for hydrazine and −0.69 for PNP. Figure 8 has displayed the amperometric response of a Gd2O3/BCA/Au electrode for the same. From the data, it has been found that a negligible change to response current has been detected in the presence of different interfering compounds. However, significant and quick responses were observed in the presence of hydrazine and PNP. These outcomes have clearly explained the high selectivity of the fabricated sensor toward hydrazine and PNP, and enhanced the scope of the developed sensor.

**Figure 8.** Amperometric response of the Gd2O3/Au electrode in the presence of different analytes with a concentration of 1 mM at pH 7 for (**a**) hydrazine and (**b**) PNP.

#### **4. Conclusions**

In summary, the current work has reported the fabrication of the ethylene glycol-mediated synthesis of Gd2O3 Nps. The formation of the particles has been scrutinized by using sophisticated characterization techniques. The effects of synthetic parameters over the optical, photoluminescence, band-gap variation, and agglomeration number of the formed particles have been studied in detail. The fabricated particles have further been employed as a proficient electrocatalytic material for the enzyme-free detection of hydrazine and p-nitrophenol with great sensitivity and selectivity. The developed sensors hold a wider linear range of 1 to 10 μM with low detection limits of 1.527 μM and 0.704 μM for p-nitrophenol and hydrazine, respectively. The sensitivity, selectivity, repeatability, recyclability, wide linear range, and detection limit make Gd2O3 Nps a favorable nanomaterial in the institution of the rapid and effectual scrutiny of harmful environmental pollutants. The realistic application of the developed sensor has also been investigated by spiking known concentrations of hydrazine and p-nitrophenol in different water samples with good recoveries. Consequently,

the successful synthesis of EG@Gd2O3 nanoparticles has immense potential for the design of highly effective electrochemical sensors, and is a probable way to provide momentum to the advancement of new electrode materials.

**Author Contributions:** S.C. for methodology, validation, writing; S.K. (Sandeep Kumar) and S.K. (Sushil Kumar) for formal analysis; A.U., G.R.C. and S.K.M. for data curation.

**Funding:** This research was funded by DST Inspire Faculty award [IFA-CH-17] and DST purse grant II. Sandeep Kumar is grateful to CSIR India for providing a senior research fellowship. Ahmad Umar would like to acknowledge the Ministry of Education, Saudi Arabia for this research through a grant (PCSED-013-18) under the Promising Centre for Sensors and Electronic Devices (PCSED) at Najran University, Kingdom of Saudi Arabia.

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

#### **References**


© 2019 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Iron-Doped Titanium Dioxide Nanoparticles As Potential Sca**ff**old for Hydrazine Chemical Sensor Applications**

**Ahmad Umar 1,2,\*, Farid A. Harraz 2,3, Ahmed A. Ibrahim 1,2, Tubia Almas 1,2,4, Rajesh Kumar 5, M. S. Al-Assiri 2,6 and Sotirios Baskoutas <sup>4</sup>**


Received: 5 January 2020; Accepted: 13 February 2020; Published: 17 February 2020

**Abstract:** Herein, we report the fabrication of a modified glassy carbon electrode (GCE) with high-performance hydrazine sensor based on Fe-doped TiO2 nanoparticles prepared via a facile and low-cost hydrothermal method. The structural morphology, crystalline, crystallite size, vibrational and scattering properties were examined through different characterization techniques, including FESEM, XRD, FTIR, UV–Vis, Raman and photoluminescence spectroscopy. FESEM analysis revealed the high-density synthesis of Fe-doped TiO2 nanoparticles with the average diameter of 25 ± 5 nm. The average crystallite size of the synthesized nanoparticles was found to be around 14 nm. As-fabricated hydrazine chemical sensors exhibited 1.44 μA μM−<sup>1</sup> cm−<sup>2</sup> and 0.236 μM sensitivity and limit of detection (LOD), respectively. Linear dynamic ranged from 0.2 to 30 μM concentrations. Furthermore, the Fe-doped TiO2 modified GCE showed a negligible inference behavior towards ascorbic acid, uric acid, glucose, SO4 <sup>2</sup>−, NO3 <sup>−</sup>, Pb2<sup>+</sup> and Ca2<sup>+</sup> ions on the hydrazine sensing performance. Thus, Fe-doped TiO2 modified GCE can be efficiently used as an economical, easy to fabricate and selective sensing of hydrazine and its derivatives.

**Keywords:** Fe-doped TiO2; hydrothermal; GCE; hydrazine; chemical sensor; amperometry

#### **1. Introduction**

Hydrazine is extensively used in rocket propellants, pesticides, explosives, photography chemicals, antioxidants and plant growth regulators in various related industries and laboratories [1,2]. Effluents released from these industries and laboratories are the foremost environmental source of hydrazine. It is considered one of the poisonous chemicals even at low concentration and is very hazardous to living organisms. The potential symptoms of hydrazine exposure range from eye and nose irritation, pulmonary edema, skin dermatitis, temporary blindness and serious damage to many human organs such as the kidney and liver [3–5]. Due to its mutagenic and carcinogenic behavior, it has been cited in the Environmental Protection Agency (EPA) list. It is therefore mandatory to trace the presence of minor amounts of hydrazine in the aqueous medium.

Previously reported techniques for the detection of hydrazine include spectrophotometry [6], fluorimetry [7,8], chemiluminescence [9] and potentiometry [10,11]. However, all these methods are

accompanied by several disadvantages like low sensitivity, and they require expensive instrumentation followed by complicated procedures. Electrochemical determination is considered to be the promising alternative for the determination of hydrazine in terms of selectivity, sensitivity and portability, with an economical and simple operating procedure [12,13]. Unfortunately, hydrazine undergoes direct oxidation at the bare electrode surface, thus resulting in sluggish electrode kinetics and high over potentials [14,15]. Therefore, chemically modified electrodes have been used to detect hydrazine, which significantly reduce the overpotential, as well as accelerate redox reactions and hence the oxidation current responses [16,17].

Recently, there is a big interest in the utilization of new materials in scientific societies due to improved processibility, applicability and vast applications in various fields of sciences. In the past few years, tremendous work has been done on metal oxide based nanomaterials or composites, especially ZnO [18], CuO [19,20], TiO2 [21,22] and MnO2 [23], which are considered an ideal platform for material preparation used in pharmaceuticals and cosmetics industries, for the treatment of wastewater and in other fields [24]. TiO2 is one of the most extensively used semiconductor materials owing to its non-toxicity, high performance, great stability and low preparation cost. Due to its good conductivity, high electrocatalytic activity and electron transport properties, TiO2 is the most suitable material for devising the electrodes for various electrochemical and biosensing applications. Khodari et al. [25] used an electrochemical sensor, prepared by casting TiO2 nanoparticles (NPs) onto a carbon electrode surface, for the determination of resorcinol. The TiO2 NPs were demonstrated to be competent in boosting the electron transfer between resorcinol and the electrode surface as well as transport of resorcinol molecules to the surface of the sensor. Guo and coworkers [26,27] synthesized TiO2 nanofibers for the bio-sensing of glucose. The fabricated TiO2 nanofibers exhibited good direct electrochemistry as well as magnificent sensitivity along with fast response time for the detection of glucose. In order to further increase the electron transfer rate between the working modified electrode and the analyte solution and, hence, rapid and sensitive current response, doping of metal oxides has been reported as one of the best ways. Among the various dopants, Fe3<sup>+</sup> ions are most suitable owing to their similar ionic radii to that of Ti4<sup>+</sup> ions. Fe3<sup>+</sup> ions (63 pm) can easily replace the Ti4<sup>+</sup> ions (68 pm) from the TiO2 crystal lattice. Further, Fe3<sup>+</sup> ions prevent electron–hole (e−–h+) recombination, better charge separation and incorporation of oxygen vacancies in the crystal lattice and surface of the TiO2 [28,29]. Additionally, the choice of Fe3<sup>+</sup> ions as dopant has also been confirmed for some other semiconductor metal oxides. Hexamethylenediamine (HMDA) grafted and Fe nanoparticles incorporation into SnO2 nanostructures exhibited strong surface affinity of 8.5 μmol/m<sup>2</sup> for H2, as an important aspect of green energy storage applications at ambient temperature and pressure conditions [30]. ZnO powders doped with Fe nanoparticles through an in situ dispersion method showed improved conductivity and capacitance as compared to undoped ZnO nanoparticles [31]. HMDA grafted and Fe nanoparticles doped ZnO nanoparticles showed excellent conductivity which was attributed to the formation of effective proton-conductivity on the surface of the ZnO as well as proton transfer between Fe nanoparticles [32]. Fe-doped TiO2 nanoparticles displayed superior photocatalytic degradation of methylene blue dye, phenol and toxic organic compounds as compared to undoped TiO2 under UV and visible light illumination [33–35].

Herein, Fe-doped TiO2 nanoparticles were prepared through a facile hydrothermal technique and subsequently analyzed using different characterization techniques for their various characteristics and to affirm the formation of the doped TiO2 nanoparticles. Modified glassy carbon electrode (GCE) was fabricated by coating a thin film of Fe-doped TiO2 nanoparticles onto it. As-fabricated hydrazine chemical sensors exhibited excellent hydrazine chemical sensing parameters.

#### **2. Experimental Details**

#### *2.1. Synthesis of Fe-Doped TiO2 Nanoparticles*

For the preparation of Fe-doped TiO2 nanoparticle, a facile hydrothermal method was adopted (2% doping of Fe ion was performed in the Fe-doped TiO2 nanoparticle). Tetrabutyl titanate (Ti(BuO)4) and ferric nitrate (Fe(NO3)3) were used as TiO2 precursor and Fe dopant, respectively. Twenty milliliters of 0.5 M Ti(BuO)4 solution was poured drop-wise to 5 M aqueous NaOH solution followed by the addition of 20 mL of 0.1 M Fe(NO3)3 solution dropwise. A Teflon lined stainless steel autoclave containing above solution was then heated at 200 ◦C for 10 h. The following reaction mechanisms have been proposed for the synthesis of Fe-doped TiO2 nanoparticles (Equations (1)–(3)).

$$\text{Ti(BuO)}\_{4} + 4\text{ NaOH} \rightarrow \text{Ti(OH)}\_{4} + 4\text{ BuONa} \tag{1}$$

$$\text{Fe(NO}\_3\text{)}\_3 + 3\text{ NaOH} \rightarrow \text{Fe(OH)}\_3 + 3\text{ NaNO}\_3\tag{2}$$

$$\text{Ti(OH)}\_{4} \xrightarrow{\text{Calcination}} \text{TiO}\_{2} + \text{ 2 H}\_{2}\text{O} \tag{3}$$

During the hydrothermal process, tetrabutyl titanate was hydrolyzed to titanium hydroxide (Ti(OH)4) and sodium butoxide. Ferric nitrate used as dopant was also hydrolyzed to its hydroxide, i.e., ferric hydroxide (Fe(OH)3). BuONa and NaNO3 were removed by washing with ethanol and deionized (DI) water. These washings also eliminated any unreacted Fe(NO3)3. Thoroughly washed Ti(OH)4 and Fe(OH)3 were dried for 6 h in an electric oven at 70 ◦C. During calcination, titanium hydroxide was converted into TiO2 with Fe3<sup>+</sup> ions occupying the lattice sites in the TiO2 crystal lattice. The calcination was carried out in oxygen atmosphere at 450 ◦C for 2 h. The calculated synthesis yield of the synthesized Fe-doped TiO2 nanoparticles were found to be 0.8 g.

#### *2.2. Characterizations of Fe-Doped TiO2 Nanoparticles*

X-ray diffraction (XRD; PAN analytical Xpert Pro. Cambridge, UK) with Cu–Kα radiation was performed for the analysis of crystal phase and crystallite size. The optical characteristics were estimated by using UV–Vis Spectrophotometer (Perkin Elmer-Lamda 950, Waltham, MA, USA) within the wavelength range of 200–800 nm by dispersing and sonicating the Fe-doped TiO2 in distilled water for 30 min. Scattering properties of the doped TiO2 material were analyzed by Raman spectrum and examined using Raman spectrometer (Perkin Elmer-FTIR Spectrum-100, Waltham, MA, USA) from 100 to 900 cm<sup>−</sup>1. FTIR spectrum of the Fe-doped TiO2 was collected by FTIR spectrophotometer (Perkin Elmer-FTIR Spectrum-100, Waltham, MA, USA) through KBr pelletization from 500 to 4000 cm−1. The photoelectronic properties of the Fe-doped TiO2 nanoparticles were analyzed through photoluminescence spectral measurement.

#### *2.3. Hydrazine Chemical Sensor Fabrication*

Prior to electrode coating with Fe-doped TiO2 nanoparticles, the GCE with surface area of 0.071 cm<sup>2</sup> (Bio-Logic SAS, Seyssinet-Pariset, France) was polished with a 1 μm polishing diamond. After this, the surface of the GCE was further polished with 0.05 μm alumina slurry and finally washed with distilled water. The modified GCEs with active materials (Fe-doped TiO2) were fabricated as follows: The GCE surface was smoothly coated by the Fe-doped TiO2 using ethyl acetate and conducting binder–butyl carbitol acetate followed by drying at 60 ◦C for 3 h. A three electrode electrochemical cell was connected to electrochemical workstation, Zahner Zennium, Germany, with a Pt wire as a counter electrode, and Fe-doped TiO2 nanoparticle modified GCE as working electrode and an Ag/AgCl (saturated KCl) as a reference electrode were used during the electrochemical measurements. Different concentrations of hydrazine (0.2 μM–30 μM) were electrochemically tested. All the electroanalytical experiments were performed in 0.1 M phosphate buffer solution (PBS) of pH 7.4 at room temperature.

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

#### *3.1. Characterizations and Properties of Fe-Doped TiO2 Nanoparticles*

Figure 1 depicts the XRD diffraction patterns of Fe-doped-TiO2 nanoparticles. The XRD studies clearly demonstrated the presence of both anatase and rutile phases with anatase as the major phase. The presence of main diffraction peaks in the XRD pattern of Fe-doped-TiO2 nanoparticles at 2θ = 25.28, 37.8, 48.07, 54.25 and 62.63, 68.9, 70.28 and 75.13 were consistent with (101), (103), (004), (200), (105), (211), (204), (116), (220) and (107) lattice planes of anatase phase (JCPDS No. 21-1272) [24].

**Figure 1.** XRD pattern of Fe-doped TiO2 nanoparticles.

The XRD peaks corresponding to rutile phase also emerged at 2θ = 27.50 and 41.54 diffraction angles corresponding to (121) and (111) diffraction planes (JCPDS No. 21-1276). The synthesized samples did not exhibit any diffraction peaks for Fe, which suggests that the Fe3<sup>+</sup> content in the Fe-doped TiO2 was below the detection limit, and due to almost similar ionic radii, the Fe3<sup>+</sup> ions could substitute Ti4<sup>+</sup> from some of the lattice sites of TiO2 as discussed earlier. This further indicates that Fe3<sup>+</sup> ions were successfully integrated into TiO2 matrix homogeneously without the development of iron oxide on the TiO2 surface. This homogeneous distribution of Fe3<sup>+</sup> ions in TiO2 matrix and low concentration, responsible for the absence of any Fe3<sup>+</sup> peaks in the XRD patterns, have also been reported earlier in the literature [36,37]. The average crystallite size of the synthesized nanoparticles was estimated using the Scherrer formula (Equation (4)) and was found to be around 14 nm.

$$\mathbf{d} = \frac{0.89\lambda}{\beta \cdot \cos\theta} \tag{4}$$

where λ = 1.542 Å, θ = Bragg angle of diffraction, β = full width at half maximum.

The detailed structural and morphological properties of Fe-doped TiO2 were examined by FESEM analysis. The corresponding low and high magnification FESEM images are depicted in Figure 2a–d. A large number of spherical shaped and highly agglomerated Fe-doped TiO2 particles with an average diameter of about 20 nm can be seen. In addition to spherical shapes, some Fe-doped TiO2 with cubic, pentagonal, oval and other irregular geometries can also be seen from high magnification FESEM images (Figure 2c,d).

**Figure 2.** (**a**) and (**b**) Low magnification and (**c**) and (**d**) high magnification FESEM images for Fe-doped TiO2 nanoparticles.

To observe the optical properties of the Fe-doped TiO2 nanoparticles, a UV–Vis absorption spectroscopic study was performed. It can be examined from Figure 3a that Fe-doped TiO2 nanoparticles exhibited a wide absorption peak below 400 nm, which is the typical absorption of TiO2. This peak can be attributed to the electronic excitation from lower energy level to higher energy level in the anatase phase of the TiO2. Furthermore, the change in color of the sample from white (pure TiO2) to creamish yellow (Fe-doped TiO2) depicted the increase in absorption towards visible light due to the incorporation of dopant metal [38].

Figure 3b illustrates the Raman spectra of Fe-doped TiO2 nanoparticles. Five main bands—144.2, 195, 396.4, 513.2 and 634.8 cm<sup>−</sup>1, corresponding to the six Raman active modes—were illustrated for the anatase phase of TiO2 (3Eg, 2B1g and 1A2g) [39]. The spectra indicated the crystalline nature of the synthesized nanoparticles. Furthermore, no additional peak related to the iron oxide was seen, which corroborates the results of XRD.

Figure 3c shows the FTIR spectra of Fe-doped TiO2 nanoparticles. The spectrum reflects that doping had no effect on the bonding environment of the TiO2 host nanoparticles. The broadband at 3433 cm−<sup>1</sup> was assigned to the symmetric and asymmetric stretching vibrations of O–H bonds of the adsorbed water molecules during sample formation. An additional peak at 1628 cm−<sup>1</sup> was attributed to the bending vibration related to the hydroxyl group of the adsorbed water. The band centered at ~520 cm−<sup>1</sup> was due to metal–oxygen, i.e., Ti–O and Fe–O, bonds [40].

The photoelectronic properties of Fe-doped TiO2 nanoparticles were studied from the photoluminescence spectrum (Figure 3d). The UV region peak at around 375 nm is likely related to the near-band-edge (NBE) excitonic emission. Interestingly, the energy corresponding to above NBE peak is close to the energy gap of 3.17 eV of anatase TiO2 as reported earlier [40]. The NBE transition originated from the electrons–holes recombination. The incorporation of dopant Fe did not cause any significant alteration in the PL spectrum.

**Figure 3.** (**a**) UV–Vis absorption spectrum, (**b**) room temperature Raman spectrum, (**c**) FTIR and (**d**) PL spectrum of Fe-doped TiO2 nanoparticles.

#### *3.2. Electrochemical Sensing Properties of Hydrazine Using Fe-Doped TiO2 Nanoparticles*

The electrochemical sensing capability of Fe-doped TiO2 nanostructure was examined using cyclic voltammetry. Figure 4a,b shows typical cyclic voltammogram (CV) of bare GCE, undoped TiO2 modified GCE and Fe-doped TiO2 modified GCE in absence and presence, respectively, of 0.5 mM hydrazine in 0.1 M PBS (pH 7.4) at 50 mVs−<sup>1</sup> scan rate. As can be observed from Figure 4a, in blank PBS solution bare GCE, undoped TiO2 modified GCE, and Fe-doped TiO2 modified GCE did not generate any characteristic peak in the selected voltage range. However, with the addition of 0.5 mM hydrazine, no significant peak was observed by bare electrode, but a significant oxidation peak at 0.45 V vs. Ag/AgCl was detected in the case of both undoped TiO2 modified GCE and Fe-doped TiO2 modified electrode (Figure 4b), which clearly illustrated the efficient electrocatalytic activity of the modified GCE. The current response at Fe-doped TiO2 modified electrode was much larger than the undoped TiO2 modified GCE, ~137% larger, which indicates the enhanced electrocatalytic activity of coated GCE after the addition of Fe.

**Figure 4.** Cyclic voltammograms measured at 50 mVs−<sup>1</sup> in 0.1 M phosphate buffer solution (PBS) (pH 7.4) (**a**) with absence of hydrazine and (**b**) in presence of 0.5 mM hydrazine using bare glassy carbon electrode (GCE), TiO2 modified GCE and Fe-doped TiO2 modified GCE.

#### *3.3. Amperometric Studies*

The amperometric (i–t) response was also carried out with the purpose to detect the hydrazine analyte using Fe-doped TiO2 nanoparticle modified GCE. Amperometric studies were performed, and the constant potential was set at a value of 0.45 V with successive addition of hydrazine (0.2–30 μM) into a continuously stirred 0.1 M PBS (pH 7.4). As revealed in Figure 5a, Fe-doped TiO2 modified GCE fabricated sensor illustrated a significant and steep rise in the current value after each successive addition of hydrazine. The measured current increased rapidly with a response time of ~20 s during the amperometric measurements. The corresponding calibration curve for amperometric hydrazine sensing showed a linear response in the concentration range 0.2 to 30 μM (Figure 5b). The correlation coefficient of the line R<sup>2</sup> was found to be 0.998, and the linear equation was y = 0.1019x + 0.0157. The linearity in the plot also confirmed that the hydrazine oxidation process was diffusion-controlled and indicated the fast electron transfer rate, which led to the sharper and well-defined peaks.

**Figure 5.** (**a**) Current–time (i–t) response of Fe-doped TiO2 modified GCE for 0.2–30 μM hydrazine concentrations at a constant potential +0.45V vs. Ag/AgCl. Inset shows an enlarged part of the early stage addition (0.2–1.2 μM). (**b**) The corresponding current–concentration calibration graph.

It has been proposed that hydrazine in slightly basic medium (pH = 7.4) is oxidized onto the surface of the Fe-doped TiO2 to release electrons (Equation (5)) [41,42].

$$2\text{ N}\_2\text{ H}\_4 + 8\text{ HO}^- \rightarrow 2\text{ NO} + 6\text{ H}\_2\text{ O} + 8\text{ e}^-\tag{5}$$

The schematic proposed hydrazine sensing by Fe-doped TiO2 is shown in Figure 6.

**Figure 6.** Proposed mechanism for the hydrazine sensing by Fe-doped TiO2 nanoparticles.

The redox behavior of the Fe3<sup>+</sup> ions facilitate the transfer of electrons released from the oxidation of the hydrazine to the conduction band of the TiO2 or migrate the electrons to reduce the Ti4<sup>+</sup> to Ti3<sup>+</sup> ions. Li et al. [43] proposed that Fe3<sup>+</sup> can act as a hole as well as an electron trap, which further enhances the electron transfer process. Furthermore, Zhu et al. [44] proposed the following reactions to depict the redox nature of Fe3<sup>+</sup> ions ((Equations (6)–(9)).

$$\text{Fe}^{3+} + \text{e}^- \rightarrow \text{Fe}^{2+} \tag{6}$$

$$\text{Fe}^{3+} + \text{h}^{+} \rightarrow \text{Fe}^{4+} \text{ (Hole trap)}\tag{7}$$

$$\text{Fe}^{3+} + \text{Ti}^{4+} \rightarrow \text{Fe}^{3+} + \text{Ti}^{3+} \text{ (Electron migration)}\tag{8}$$

$$\text{Fe}^{3+} + \text{Ti}^{3+} \rightarrow \text{Fe}^{2+} + \text{Ti}^{3+} \text{ (Recombination)}\tag{9}$$

Therefore, when hydrazine comes in contact with Fe-doped TiO2, the electron density in the conduction band increases, which leads to the increase in electrical conductivity and, thus, increase in the current potential. The higher the concentration of the hydrazine, the greater is the electron density in the conduction band and, hence, the current potential is higher.

The sensitivity of the synthesized sensor was calculated from the slope of the calibration curve divided by the electrode area [45,46]. The limit of detection (LOD) of the fabricated hydrazine sensor was accordingly estimated via the following equation (Equation (10)):

$$\text{LOD} = 3.0^{\circ} \text{ } \sigma\_{\text{B}}/b \tag{10}$$

where σ<sup>B</sup> is the standard deviation of the population of the blank signals (0.008 μA) and *b* is the slope of the regression line. It is possible to replace σ<sup>B</sup> by the residual standard deviation of the regression, sy/x, also known as standard error of the regression [47]. The sensitivity and LOD were found to be 1.44 μA μM−<sup>1</sup> cm−<sup>2</sup> and 0.236 μM, respectively. These obtained results demonstrated the potential of the Fe-doped TiO2 nanostructure modified electrode as a suitable electrochemical sensor for sensitive and selective determination of hydrazine. The synthesized hydrazine electrochemical sensor was reliable and depicted fantastic reproducibility. It was ascertained that no significant decrease in sensitivity was observed when tested for more than three weeks while being stored at room temperature in a closed container.

In order to assess the analytical potential of the fabricated sensor for real sample analysis, the selectivity of Fe-doped TiO2 nanostructures to hydrazine was evaluated. The selectivity test of the sensor was conducted to check the influence of some interfering electro-active chemical species on the sensing property of Fe-doped TiO2 by measuring the i–t response. Figure 7 demonstrates the amperometric responses of Fe-doped TiO2 modified GCE for the successive addition of different concentrations of hydrazine and 100 μM of ascorbic acid (AA), uric acid (UA), glucose, SO4 <sup>2</sup><sup>−</sup>, NO3 −, Pb2<sup>+</sup> and Ca2<sup>+</sup> at a regular interval of 100 s in 0.1 M PBS at an applied potential value of +0.45 V. The negligible change in the observed current, even if the various co-existing interfering species were present, undoubtedly revealed the excellent selectivity of the fabricated sensor.

**Figure 7.** Amperometric (i–t) measurement showing the interference behavior of the Fe-doped TiO2 coated GCE upon the successive injections of 10 or 5 μM hydrazine and each 100 μM of AA, glucose, UA, SO4 <sup>2</sup><sup>−</sup>, NO3 <sup>−</sup>, Pb2<sup>+</sup> and Ca2<sup>+</sup> into a continuously stirred 0.1 M PBS (pH 7.4) solution operating at +0.45 V vs. Ag/AgCl.

The sensor parameters of Fe-doped TiO2 nanoparticle modified GCE and other recently reported sensors are compared in Table 1. Detailed comparison shows that the fabricated sensor had excellent electrocatalytic performance for selective detection and sensing of hydrazine.


**Table 1.** Comparison of the sensor parameters of the Fe-doped TiO2 coated GCE sensor with some recently reported hydrazine electrochemical sensor materials.

#### **4. Conclusions**

In summary, highly crystalline Fe-doped TiO2 nanoparticles were synthesized through hydrothermal synthesis and subsequently characterized by different characterization techniques. Finally, Fe-doped TiO2 nanoparticles were applied as an efficient electron mediator for the fabrication of hydrazine chemical sensor using GCE. As-fabricated modified GCE showed sensitivity, LOD and LDR of 1.44 μA μM−<sup>1</sup> cm<sup>−</sup>2, 0.236 μM and 0.2–30 μM, respectively, through an amperometric sensing approach. It was proposed that redox behavior of the Fe3<sup>+</sup> ions facilitates the transfer of electrons released from the oxidation of the hydrazine to the conduction band of the TiO2, which leads to the increase in electrical conductivity. The negligible change in the observed current in the presence of various co-existing interfering chemical species further confirms the excellent selectivity of the fabricated sensor towards hydrazine.

**Author Contributions:** A.U., F.A.H., A.A.I., T.A., R.K., M.S.A.-A. and S.B. conceived of the presented idea and developed and performed the experiments. All of the authors discussed the results and contributed to the final manuscript. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Ministry of Education, Kingdom of Saudi Arabia through a grant (PCSED-013-18) under the Promising Centre for Sensors and Electronic Devices (PCSED) at Najran University, Kingdom of Saudi Arabia.

**Acknowledgments:** Authors would like to acknowledge the support of the Ministry of Education, Kingdom of Saudi Arabia for this research through a grant (PCSED-013-18) under the Promising Centre for Sensors and Electronic Devices (PCSED) at Najran University, Kingdom of Saudi Arabia.

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

#### **References**


© 2020 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **ZnO Nanocrystal-Based Chloroform Detection: Density Functional Theory (DFT) Study**

#### **H. Y. Ammar 1,2,3,\*, H. M. Badran 1,2,3, Ahmad Umar 1,4,\*, H. Fouad 5,6,\* and Othman Y. Alothman <sup>7</sup>**


Received: 5 October 2019; Accepted: 14 November 2019; Published: 19 November 2019

**Abstract:** We investigated the detection of chloroform (CHCl3) using ZnO nanoclusters via density functional theory calculations. The effects of various concentrations of CHCl3, as well as the deposition of O atoms, on the adsorption over ZnO nanoclusters were analyzed via geometric optimizations. The calculated difference between the highest occupied molecular orbital and the lowest unoccupied molecular orbital for ZnO was 4.02 eV. The most stable adsorption characteristics were investigated with respect to the adsorption energy, frontier orbitals, elemental positions, and charge transfer. The results revealed that ZnO nanoclusters with a specific geometry and composition are promising candidates for chloroform-sensing applications.

**Keywords:** nanocrystal; ZnO; density of states; optical and electrical properties

#### **1. Introduction**

The rapid development of various important industries, such as automobiles, pharmaceuticals, textiles, food, and agriculture, has substantially contributed to environmental pollution [1]. The release of various toxic and harmful gases and chemicals from such industries has significantly disturbed the ecosystem, and poses a great threat not only to humans, but to all living beings [2,3]. Among the various toxic gases, chloroform, which is also known as tri-chloromethane or methyl-tri-chloride, is considered to be one of the most toxic gases, and evaporates quickly when exposed to air [4]. It is widely used by chemical companies and in paper mills. Chloroform lasts for a long time in the environment, and its breakdown products, such as phosgene and hydrogen chloride, are as toxic or even more toxic [5]. The exposure of humans to chloroform severely affects the central nervous system, kidneys, liver, etc. Long-term exposure may result in vomiting, nausea, dizziness, convulsions, depression, respiratory failure, coma, and even sudden death [6,7]. It is important to efficiently detect the release of chloroform because of its serious health hazards. Thus, various methods have been reported for detecting chloroform, which involve optical sensors, colorimetric sensors, fluorescent sensors, electrochemical sensors, resistive gas sensors, luminescent sensors, photo-responsive sensors, etc. [8–13].

Among the various sensing techniques, gas sensors have attracted considerable attention because of their facile manufacturing process, high sensitivity, and low detection limit [14–17]. The literature reveals that metal-oxide materials are the most widely used scaffold to fabricate gas sensors [15–20]. In particular, metal-oxide materials are widely utilized to fabricate sensors for toxic and explosive gases [20–24]. It has been observed that the nanocrystal interfaces can significantly influence the optical and electrical properties and charge-trapping phenomena [22–25]. Zinc oxide (ZnO) is one of the most important and functional materials because of its various significant properties, including its wide bandgap; high exciton binding energy, piezoelectricity, and pyroelectricity; high conductivity and electron mobility; good stability in chemical and thermal environments; and biocompatibility [26–28]. Therefore, to improve the gas-sensing performance of ZnO-based gas sensors, various approaches have been employed, such as doping, surface modification, and the fabrication of composites [26,29]. Although ZnO materials are used for various gas-sensing applications [30–34], there are few reports available on the use of ZnO materials for chloroform sensing. Ghenaatian et al. [35] investigated the Zn12O12 nanocage as a promising adsorbent and detector for CS2. Baie et al. examined the Zn12O12 fullerene-like cage as a potential sensor for SO2 detection [36]. Ammar [37] reported that the Zn12O12 nanocage is a potential sorbent and detector for formaldehyde molecules. Nanocrystalline ZnO thin-film gas sensors were investigated by Mayya et al. [38] for the detection of hydrochloric acid, ethanolamine, and chloroform. Additionally, it is important to examine various geometries and other electronic parameters in order to obtain the optimal sensing material based on ZnO.

In this study, we investigated the detection of chloroform (CHCl3) using ZnO nanoclusters via density functional theory (DFT) calculations implemented in a Gaussian 09 program. The effect of various concentrations of CHCl3, as well as the deposition of O atoms, on the adsorption over ZnO nanoclusters was analyzed via geometric optimizations. To fully exploit the ZnO nanocrystals, various calculations related to the gas-sensing properties were performed.

#### **2. Methods and Computational Details**

A quantum cluster consisting of 24 atoms (Zn12O12) was selected to study the interaction between the ZnO nanocage and the CHCl3 molecule. DFT calculations were performed with the Gaussian 09 suite of programs [39]. The calculations were conducted using Becke's three-parameter B3 with the Lee, Yang, and Parr (LYP) correlation functional [40]. This B3LYP hybrid functional contains the exchange–correlation functional, and is based on the exact form of the Vosko–Wilk–Nusair correlation potential [41]. Originally, the functional B included the Slater exchange along with corrections involving the gradient of density [42]. The correlation functional LYP was that of Lee, Yang, and Parr, which includes both local and nonlocal terms [43,44]. For the ZnO nanocage, the standard LANL2DZ basis set [37,45] was used. For the CHCl3 and the deposited O atoms, a 6-31G (d, p) basis set was used. The adsorption energy (*E*ads) of the CHCl3 molecule on the surface of the Zn12O12 nanocage is defined as follows:

$$E\_{\rm ads} = \left[ E\_{\rm (CHCl\_3)\_n/ZnO} - \left( nE\_{\rm CHCl\_3} + E\_{\rm ZnO} \right) \right] / n\_\prime \tag{1}$$

where *E*CHCl3 , *E*ZnO and *E*(CHCl3)*n*/ZnO represent the energies of a single CHCl3 molecule, the pristine Zn12O12 nanocage, and the (CHCl3)*n*/Zn12O12 complex, respectively.

The adsorption energy (*E*ads) of an O atom on the surface of the Zn12O12 nanocage is defined as follows:

$$E\_{\rm ads} = \left[ E\_{\rm O\_2/ZnO} - (nE\_{\rm O} + E\_{\rm ZnO}) \right] / n \,, \tag{2}$$

where *E*<sup>O</sup> and *E*O*n*/ZnO represent the energies of a single O atom and the O*n*/Zn12O12 complex, respectively.

The adsorption energy (*E*i) of a CHCl3 molecule on the deposited O on the Zn12O12 nanocage is defined as follows:

$$E\_1 = \left[ E\_{\mathrm{(CHCl}\_3\big)\_n/\mathrm{O}\_n/\mathrm{ZnO}} - \left( nE\_{\mathrm{CHCl}\_3} + nE\_{\mathrm{O}} + E\_{\mathrm{ZnO}} \right) \right] / n\_\prime \tag{3}$$

where *<sup>E</sup>*(CHCl3)*n*/O*n*/ZnO represents the energy of the (CHCl <sup>3</sup>)*n*/O*n*/ZnO complex.

The positive and negative values of *Ei* indicate the endothermic and exothermic processes, respectively. The binding energy (*E*b) between the *X* and *Y* fragments of the *XY* complex is defined as follows:

$$E\_\mathbf{b} = E\_\mathbf{XY} - (E\_\mathbf{X} + E\_\mathbf{Y}),\tag{4}$$

where *E*XY represents the total energy of the optimized molecule, and *E*<sup>X</sup> and *E*<sup>Y</sup> represent the energies of the two fragments *X* and *Y*, respectively, having the same geometric structure as in the XY complex. The GaussSum 2.2.5 program was used to calculate the densities of states (DOSs) for the Zn12O12 nanocage, CHCl3, and other complex systems [46]. Full natural bond orbital (NBO; NBO version 3.1) analyses were used to estimate the charge distributions for the Zn12O12 nanocages, CHCl3, and other complex systems [47].

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

#### *3.1. Geometric Optimization*

The geometric optimization of a pristine Zn12O12 nanocage was performed. Zn12O12 is composed of eight (ZnO)3 and six (ZnO)2 rings, forming a cluster in which all of the Zn and O vertices are equivalent [48], as shown in Figure 1. The examined structural properties of the Zn12O12 nanocage agreed well with previous studies [37,45,49]. For example, the bond lengths *R*Zn-O of 1.91 and 1.98 Å were close to the previously reported values of 1.91 and 1.98 Å, respectively [37,45], and 1.89 and 1.97 Å, respectively [47]. The calculated highest occupied molecular orbital (HOMO)–lowest unoccupied molecular orbital (LUMO) energy gap for ZnO was found to be 4.02 eV, which agrees well with a previous work [45].

**Figure 1.** Optimized structures and densities of states (DOSs) of the ZnO nanocage (Zn12O12) and CHCl3 used in the calculations. (**a**,**b**) Optimized structure and DOS of the ZnO nanocage; (**c**,**d**) optimized structure and DOS of CHCl3. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent the occupied and virtual states, respectively.

The DOS of the Zn12O12 nanocage was calculated, as shown in Figure 1. A geometric optimization was performed for the chloroform molecule (CHCl3). It is a tetrahedral molecule, as shown in Figure 1. The calculated structural properties of CHCl3 indicated that bond lengths *R*C-H and *R*C-Cl were 1.09 and 1.79 Å, respectively, and angles *A*H-C-Cl and *A*Cl-C-Cl were 107.5◦ and 11.4◦, respectively. The

energy gap (*E*g) between the HOMO and LUMO was calculated to be 7.27 eV. The DOS for CHCl3 was calculated, and is presented in Figure 1.

#### *3.2. CHCl3 Interaction with the Zn12O12 Nanocage*

The geometric optimizations for four probable orientations of CHCl3 on the surface of the Zn12O12 nanocage were investigated. Figure 2 shows the four orientations where the CHCl3 molecule may interact via its H head or Cl head, and may be absorbed over the O site or Zn site of the Zn12O12 nanocage. The adsorption energy was calculated using Equation (1). The electronic properties of the CHCl3 adsorption modes are presented in Table 1. For the first adsorption mode (a), the CHCl3 molecule was weakly chemically adsorbed, and for the other modes (b, c, and d), the CHCl3 molecule was physically adsorbed. The boundary value between the physical and chemical adsorption was considered to be 0.21 eV [50,51]. In mode (a), owing to the chemical interaction, the Fermi level (*E*FL) for the cluster was reduced by 0.17 eV, and the dipole moment (D) was increased to 3.05 Debye. There was no noticeable change in the HOMO–LUMO energy gap. In all of the adsorption modes, it was found that the HOMO–LUMO energy gaps of the CHCl3/ZnO complexes were in the range of 4.00–4.03 eV. Consequently, the adsorption of CHCl3 on the ZnO nanocage had no significant effect on the HOMO–LUMO energy gap.

**Figure 2.** Optimized structures and DOSs of the CHCl3 molecule adsorption on the Zn12O12 nanocage. (**a**) adsorption mode a, (**b**) adsorption mode b, (**c**) adsorption mode c, and (**d**) adsorption mode d. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent occupied and virtual states, respectively.


**Table 1.** Electronic properties of the isomeric configurations of the CHCl3/Zn12O12 complexes, namely: adsorption energy (*E*ads; eV), HOMO (eV), LUMO (eV), Fermi level (*E*FL; eV), HUMO–LUMO energy gap (*E*g; eV), natural bond orbital (NBO) charge (*Q*; au), and dipole moment (*D*; Debye).

Additionally, to investigate the effect of the CHCl3 concentration on the adsorption over the Zn12O12 nanocage, we performed geometric optimizations for *n* CHCl3 molecules (*n* = 1, 2, 3, and 4) adsorbed simultaneously over the Zn12O12 nanocage to form (CHCl3)*n*/ZnO complexes. All of the CHCl3 molecules had an orientation in which the H head of the CHCl3 molecule was directed toward an O site of the Zn12O12 nanocage, which is the most energetic stable orientation, as presented in Figure 2. The adsorption energies (*E*ads) were calculated using Equation (1), and are presented in Table 2.

**Table 2.** Electronic properties of the (CHCl3)*n*/Zn12O12 complexes, namely: adsorption energy (*E*ads; eV), HOMO (eV), LUMO (eV), Fermi level (*E*FL; eV), HUMO–LUMO energy gap (*E*g; eV), NBO charge (*Q*; au), and dipole moment (*D*; Debye).


The optimized structures of (CHCl3)*n*/ZnO and their DOSs are shown in Figure 3. As indicated by Table 2, after the second molecule was adsorbed, the adsorption energy (*E*ads) increased as *n*—the number of adsorbed CHCl3 molecules—increased. Additionally, as the number of adsorbed CHCl3 molecules increased, the Fermi level decreased. Furthermore, although there were no significant changes in the average acquired charge (*Q*CHCl3 ) on the CHCl3 molecules, the dipole moment was sensitive to the number of adsorbed CHCl3 molecules. The HOMO–LUMO energy gap (*E*g), compared with that of the pristine Zn12O12 nanocage (4.02 eV), was not affected by the number of adsorbed CHCl3 molecules.

**Figure 3.** Optimized structures and DOSs for the (CHCl3)*n*/Zn12O12 nanocage. (**a**) CHCl3/Zn12O12, (**b**) (CHCl3)2/Zn12O12, (**c**) (CHCl3)3/Zn12O12, and (**d**) (CHCl3)4/Zn12O12. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent the occupied and virtual states, respectively.

#### *3.3. O Atom Interaction with the Zn12O12 Nanocage*

To improve the sensitivity of Zn12O12 to the CHCl3 molecules, an O atom was deposited onto the cluster. To investigate the ability of the Zn12O12 nanocage to adsorb an O atom, the O atom was added at three different sites, namely: an O site, a Zn site, and the middle of the ZnO bond. Then, a full geometric optimization was performed for the O/Zn12O12 complexes. With the optimization, there are only two possible O/Zn12O12 complexes, as shown in Figure 4. The *E*ads were calculated using Equation (2). As shown in Table 3, the *E*ads values of the O atom on the Zn12O12 nanocage were −1.98 and −1.62 eV for complexes (a) and (b), respectively. This indicated that a chemical bond was formed between the O atom and the Zn12O12 cluster. Additionally, the NBO analysis indicated that the O atom gained negative charges (*Q*O) of −0.71|e| and −0.61|e| for complexes (a) and (b), respectively.


**Table 3.** Electronic properties of the isomeric configurations of the O/Zn12O12 complexes, namely: adsorption energy (*E*ads; eV), HOMO (eV), LUMO (eV), Fermi level (*E*FL; eV), HUMO–LUMO energy gap (*E*g; eV), NBO charge (*Q*; au), and dipole moment (*D*; Debye).

**Figure 4.** Optimized structures and DOSs of O atom adsorption on the Zn12O12 nanocage. (**a**) complex a, and (**b**) complex b. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent the occupied and virtual states, respectively.

This strong interaction is attributed to the charge transfer from the Zn12O12 nanocage to the adsorbed O atom. As indicated by the DOS in Figure 4, the HOMO–LUMO energy gaps (*E*g) of O/Zn12O12 for complexes (a) and (b) were reduced (to 3.53 and 3.78 eV, respectively) compared with that of the pristine Zn12O12 nanocage (4.02 eV; Table 1). Furthermore, for O/Zn12O12 complexes (a) and (b), increases of 0.24 and 0.12 eV, respectively, were observed for the Fermi level (*E*FL), and the dipole moment increased to 2.03 and 0.72, respectively. This indicated that the deposited O atom significantly affected the electronic properties of the Zn12O12 nanocage, and consequently may have affected its ability to adsorb CHCl3 molecules.

#### *3.4. CHCl3 Interaction with O Atoms Deposited on the Zn12O12 Nanocage*

The CHCl3 molecule could interact via its H head or Cl head, and the O atom could be deposited on the Zn or O sites of the nanocage; thus, there were four possible geometric structures for the CHCl3/O/Zn12O12 complexes. Consequently, we performed geometric optimization for the four aforementioned CHCl3/O/Zn12O12 complexes. During the optimization process, we found only three stable CHCl3/O/Zn12O12 complexes, as shown in Figure 5. The properties of the interaction among the CHCl3 molecule, deposited O atom, and Zn12O12 nanocage are presented in Table 4.

**Figure 5.** Optimized structures and DOSs for the CHCl3/O/Zn12O12 nanocage. (**a**) complex a, (**b**) complex b, and (**c**) complex c. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent the occupied and virtual states, respectively.

**Table 4.** Electronic properties of the isomeric configurations of the CHCl3/O/Zn12O12 complexes, namely: adsorption energy (*E*ads; eV), binding energy (*E*b; eV), HOMO (eV), LUMO (eV), Fermi level (*E*FL; eV), HUMO–LUMO energy gap (*E*g; eV), NBO charge (*Q*; au).


*Coatings* **2019**, *9*, 769

The adsorption energies (*E*ads) for the complexes ranged from −0.92 to −2.44 eV. These values indicate a chemical interaction, which may have been due to a charge transfer. This can be explained by the NBO analysis, which revealed that in complexes (a) and (b), the deposited O atom gained negative charges of −0.63|e| and −0.62|e|, respectively. These charges were mainly transferred from the Zn12O12 nanocage, which gained positive charges of 0.56|e| and 0.59|e|, respectively. Additionally, there was a small charge from the CHCl3 molecule, which gained positive charges of 0.07|e| and 0.03|e|, respectively. However, in complex (c), the charge was transferred from the CHCl3 molecule, which gained a positive charge of 0.86|e|, to both the Zn12O12 nanocage and the deposited O atom, which gained negative charges of −0.10|e| and −0.76|e|, respectively. Clearly, the nature of the interaction in complex (c) was significantly different from those for complexes (a) and (b). This led to different binding energies between the CHCl3 fragment and the O/Zn12O12 fragment of the CHCl3/O/Zn12O12 complexes, which were −0.68, −0.15, and −2.46 eV for complexes (a), (b), and (c), respectively. Such interactions between CHCl3 and the O/Zn12O12 nanocage led to an increase in the Fermi level (*E*FL), from −4.64 to −4.49 eV, as well as a reduction of the HOMO–LUMO energy gaps (*E*g), from 3.27 to 3.64 eV for the CHCl3/O/Zn12O12 complexes, compared with 4.02 eV for the pristine Zn12O12 nanocage.

To examine the effect of the CHCl3 concentration on the interaction with the O/Zn12O12 nanocage, we performed geometric optimizations for *n* CHCl3 molecules (*n* = 1, 2, 3, and 4), adsorbed simultaneously over *n* deposited O atoms on the Zn12O12 nanocage. Each CHCl3 molecule interacted via its Cl head with a deposited O atom on the Zn site of the Zn12O12 nanocage. This orientation yielded the highest binding energy between the CHCl3 molecule and O/Zn12O12. The interaction energies were calculated using Equation (3). The optimized structures of (CHCl3)*n*/O/Zn12O12 and their DOSs are shown in Figure 6.

**Figure 6.** Optimized structures and DOSs for the (CHCl3)*n*/O*n*/Zn12O12 nanocage. (**a**) CHCl3/O/Zn12O12, (**b**) (CHCl3)2/O2/Zn12O12, (**c**) (CHCl3)3/O3/Zn12O12, and (**d**) (CHCl3)4/O4/Zn12O12. The distances are in Å, and the DOS is in arbitrary units. The solid and dashed lines represent the occupied and virtual states, respectively.

The interaction energies are presented in Table 5. The adsorption energy remained relatively constant (approximately −0.96 eV) for the first three CHCl3 interacting molecules, and decreased for the fourth CHCl3 molecule (to −0.86 eV). Furthermore, as the number of adsorbed CHCl3 molecules

increased, the average acquired positive charges on CHCl3 (*Q*CHCl3 ) decreased, and the negativity of the average charges on the deposited O atom (*Q*O) decreased, while the negativity of the charges on the Zn12O12 nanocage increased. Additionally, with the increasing number of adsorbed CHCl3 molecules, the Fermi level (*E*FL) increased and the HOMO–LUMO energy gap (*E*g) decreased, compared with the pristine Zn12O12 nanocages. The dipole moment of (CHCl3)*n*/O/Zn12O12 was sensitive to the number of CHCl3 molecules.

**Table 5.** Electronic properties of the (CHCl3)*n*/(O)*n*/Zn12O12 complexes, namely: adsorption energy (*E*ads; eV), HOMO (eV), LUMO (eV), Fermi level (*E*FL; eV), HUMO–LUMO energy gap (*E*g; eV), NBO charge (*Q*; au), and dipole moment (*D*; Debye).


#### *3.5. Zn12O12 Nanocage as a Sensor for CHCl3*

It has been observed that during the adsorption process, the change in the HOMO–LUMO energy gap (*E*g) is related to the sensitivity of the sorbent for the adsorbate. However, the reduction of *E*g of the cluster significantly affects the electrical conductivity, as indicated by the following equation [52]:

$$
\sigma \propto e^{(-E\_g/2kT)} \tag{5}
$$

where σ represents the electrical conductivity, K represents Boltzmann's constant, and *T* represents the temperature. According to Equation (5) and the *E*g values in Tables 1 and 2, the adsorption of the CHCl3 molecule in the gas phase did not lead to significant changes in the *E*<sup>g</sup> of the Zn12O12 nanocage. According to Tables 4 and 5, the CHCl3 molecule adsorption over the oxygenated ZnO significantly reduced the *E*g values.

The energy difference between the nucleophile HOMO and electrophile LUMO is one of the important factors for HOMO–LUMO interactions. As previously mentioned, the chemical bonding between CHCl3 and the oxygenated ZnO cluster in the CHCl3/O/Zn12O12 complexes is due to the charge-transfer mechanism. It can be explained as the contribution from the HOMO of the O/Zn12O12 cluster to the vacant LUMO of the CHCl3 molecule. Figure 7 shows the surfaces of the frontier molecular orbitals (FMOs; HOMO/LUMO) for CHCl3, Zn12O12, O/Zn12O12, and CHCl3/O/Zn12O12. The HOMO and the LUMO of the Zn12O12 cluster are localized on the Zn and O sites, respectively. Thus, the Zn sites are electrophilic centers, whereas the O sites are nucleophilic centers. This explains why the H atom of CHCl3 is attached to the O site in the most stable structure of the CHCl3/Zn12O12 complex. Additionally, the HOMO of the O/Zn12O12 cluster is localized around the deposited O atom. This explains why CHCl3 is attracted to the deposited atom of the O/Zn12O12 cluster.

**Figure 7.** Frontier molecular orbital (FMO) surfaces (HOMO–LUMO) for CHCl3, Zn12O12, O/Zn12O12, and CHCl3/O/Zn12O12.

Figure 8 shows the energy diagrams of the FMOs (HOMO/LUMO) for CHCl3, Zn12O12, O/Zn12O12, and CHCl3/O/Zn12O12. Our FMO studies revealed that the deposited O atom increased the HOMO of the ZnO cluster from −6.81 to −6.32 eV. Consequently, the energy gap between the HOMO of ZnO and the LUMO of CHCl3 decreased, making the charge transfer from the O/Zn12O12 cluster to the CHCl3 easier than that from the pristine Zn12O12 cluster. Thus, the O/Zn12O12 cluster is more sensitive to the CHCl3 molecule than the pristine Zn12O12 cluster.

**Figure 8.** Energy diagram of the FMOs (HOMO/LUMO) for CHCl3, Zn12O12, O/Zn12O12, and CHCl3/O/Zn12O12.

#### **4. Conclusions**

A DOS study of chloroform sensing, based on ZnO nanocrystals, was performed via calculations implemented using the Gaussian 09 suite of programs. A geometric optimization was performed for the ZnO nanocrystal. The DOS for the ZnO nanocrystal was calculated. The calculated gap between the HOMO and the LUMO was found to be 4.02 eV. Furthermore, the effect of the concentration of CHCl3 on its adsorption over the ZnO nanocrystals was investigated. The results indicated that the electrical properties of ZnO were not affected by the concentration of CHCl3. Additionally, the effect of depositing O atoms on the ZnO adsorption properties was examined. The results indicated that the adsorption of CHCl3 on the oxygenated ZnO reduced its bandgap. The findings of this study confirm that the deposition of O on a ZnO nanocluster increases its sensitivity to CHCl3, and may facilitate CHCl3 removal or detection.

**Author Contributions:** Conception and design of the experiments, H.Y.A., H.M.B. and A.U.; implementation of the experiments, H.Y.A., H.M.B. and A.U.; analysis of the data, contribution of the analysis tools, and writing and revision of the paper, H.Y.A., H.M.B., A.U., H.F. and O.Y.A.

**Funding:** This work is funded by Deanship of Scientific Research at King Saud University under the research groups grant (No. RG-1435-052).

**Acknowledgments:** The authors would like to extend their sincere appreciation to the Deanship of Scientific Research at King Saud University for funding this research group (No. RG-1435-052). The authors thank RSSU at King Saud University for their technical support.

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

#### **References**


© 2019 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Multi-Layered Mesoporous TiO2 Thin Films: Photoelectrodes with Improved Activity and Stability**

#### **Enno Gent, Dereje H. Taffa \* and Michael Wark**

Institute of Chemistry, Chemical Technology 1, Carl von Ossietzky University of Oldenburg, D-26129 Oldenburg, Germany; enno.gent@uol.de (E.G.); michael.wark@uol.de (M.W.)

**\*** Correspondence: dereje.hailu.taffa@uol.de; Tel.: +49-441-798-3279

Received: 25 August 2019; Accepted: 24 September 2019; Published: 28 September 2019

**Abstract:** This work aims at the identification of porous titanium dioxide thin film (photo)electrodes that represent suitable host structures for a subsequent electrodeposition of plasmonic nanoparticles. Sufficient UV absorption and electrical conductivity were assured by adjusting film thickness and TiO2 crystallinity. Films with up to 10 layers were prepared by an evaporation-induced self-assembly (EISA) method and layer-by-layer deposition. Activities were tested towards the photoelectrochemical oxidation of water under UV illumination. Enhanced activities with each additional layer were observed and explained with increased amounts of immobilized TiO2 and access to more active sites as a combined effect of increased surface area, better crystallinity and improved transport properties. Furthermore, films display good electrochemical and mechanical stability, which was related to the controlled intermediate thermal annealing steps, making these materials a promising candidate for future electrochemical depositions of plasmonic noble metal nanoparticles that has been further demonstrated by incorporation of gold.

**Keywords:** titanium dioxide; mesoporous; thin film; multi-layered; photoanode; semiconductor; photoelectrochemical water oxidation

#### **1. Introduction**

Since the seminal discovery of water photolysis by Fujishima and Honda in 1972 [1], TiO2-based materials have experienced a rarely diminished attention in material research. Due to its unique set of properties, namely being inexpensive, environmentally benign, stability in aqueous media, recyclability and having favorable band energies, TiO2 is an appealing material choice [2–4]. Moreover, its very flexible processability and intensively explored synthesis routes to modify its structure lead to a wide range of energy and environmental applications [5–7]. To name a few non-photocatalytic applications, these can be sensors [8,9], photovoltaic devices [10,11] and Li-ion batteries [12,13]. Additionally, TiO2 can be found in various (heterogenous) photocatalytic applications ranging from self-cleaning/sterilizing and anti-fogging surfaces [3,14], over water and air purification [15,16] and anticorrosive coatings [17,18] to CO2 photoreduction [19–21] and H2 evolution from water [5,22].

To achieve a good photocatalytic performance, it is necessary to know the desired application before optimizing the catalyst with respect to both its structure and its optoelectrochemical properties. Structurally, mass transfer (reactant to active sites), charge transfer (catalyst surface to reactant) and charge transport (catalyst bulk to its surface) must be considered [4]. Access to more active sites, an improved collection/harvesting of photogenerated charge carriers and light trapping [23] are desirable approaches towards higher quantum yields. In light of this, extensively documented experimental techniques towards various target structures have emerged over the course of recent decades [4]. Today, TiO2 can be shaped into films [24–26], spheres [27,28], fibers [29,30], rods [31,32], tubes [26,33], etc. with variable pore sizes and/or pore size distributions ranging from micro- [34,35] to macroporous [36,37] and—occasionally—bimodal hierarchical [38,39].

Optoelectrochemical properties highly depend on the nanostructural features of titanium dioxide, such as crystallinity and crystallite size, grain boundaries and the types and amounts of defects. However, there are also two major drawbacks that intrinsically impair the efficiency of pure TiO2-based photocatalysts: the first is a band gap of ≈3.0–3.2 eV which limits photon absorption to the UV spectrum and thus to the small fraction of ≈5% solar photon energy [2,40]; the second is relatively fast recombination rates of photogenerated electron–hole pairs due to their limited mobility and short lifetimes [2,41]. There are numerous approaches to address these two issues including metal and/or nonmetal-doping [40,42,43] (to partially replace either the Ti4<sup>+</sup> and/or O2<sup>−</sup> sites), heterojunction formation [20,44], Z-schemes [45,46], decoration with quantum dots [47] and—with more recent success—hydrogenation towards "black" TiO2 [48] or introduction of surface plasmon resonance (SPR) active noble metal nanoparticles (NPs) [41,49–53]. To develop TiO2 photocatalysts, it is, therefore, necessary to address both origins—structure and intrinsic optoelectrochemical properties.

The scope of the present study is focused on a TiO2-based system to identify suitable mesoporous host structures for further modifications related to both origins. It is important to understand the intended modifications, since they define requirements for the investigated material. At a later stage of this study, the electrochemical deposition (ED) of SPR-active Au-NPs is outlined. The basic appeal of decorating TiO2 with such NPs is that both major drawbacks can be tackled simultaneously: Visible-light response can be attained via the SPR effect plus subsequent injection of hot electrons into the conduction band of TiO2 [54] (where they can be consumed for reduction processes [55,56]) and improved electron/hole separation, since the metal–TiO2 interface forms a Schottky barrier thus decelerating charge carrier recombination [41,51]. SPR features are not only size and shape dependent [57], but the surrounding chemical environment and its interaction with the NPs also plays a crucial role. When considering the latter fact together with the intended ED, the most important requirement for the TiO2 host system becomes apparent: being a (meso)porous TiO2 electrode.

Mesoporous TiO2 thin films (MTTFs) not only display tunable host matrices due to their well-definable pore system [58,59], they are also beneficial for facilitated product recovery due to their immobilized nature. A frequently reported technique towards MTTFs—which is commonly used for dip-, spin-coating [60]—is the balanced combination of sol-gel chemistry and the evaporation-induced self-assembly (EISA) in the presence of surfactants which yield the metal oxide and the porous network, respectively [61–63]. Five steps, namely (i) precursor selection and initial sol preparation, (ii) deposition, (iii) controlled aging, (iv) template removal and (v) wall crystallization, govern the final film properties as Soler-Illia et al. pointed out [62]. Recently, the same authors contributed to state-of-the-art designs of plasmon-enhanced MTTFs with structural control of both the host and SPR-active NPs guests [64–68]. Their groups prepared well-ordered MTTFs derived either from mild oxidative at 350 ◦C air calcination [64–66] or extractive template removal [67,68] and employed an impregnation–reduction method to introduce noble metal NPs. The resulting films were for example tested for surface-enhanced Raman scattering (SERS) activity for sensing applications.

The electrochemical incorporation of metallic NPs into MTTFs necessitates electronic conductivity of the films. Thus, improved crystallinity is required to allow deposition throughout the film rather than on the electrode's back contact. This can be achieved by elevated annealing temperatures which usually are accompanied by significant mesostructural deterioration due to mass migration and crystal growth [58,62]. The trade-off between loss of active surface area and improved crystallinity can be manipulated to some extent with an additional thermal treatment under inert-gas atmosphere [47,69]. However, the degree of mesoporosity—which determines accessibility of active sites and diffusion of molecules through the film—seems to be more important than the order of the porosity [62,70].

Photocatalytic activity not only depends on crystal sizes and structure alone but also on the portion of absorbed photons. The EISA approach mostly results limited amount of material per deposited layer [71–73] due the low viscosity and diluted concentrations of EISA-sols [58]. This necessitates a more time-consuming but worthwhile layer-by-layer deposition to prepare thicker films [47,72,74,75]. Such MTTFs display considerably high active surface sites that are accessible throughout the film as

result of its mesoporosity and maximize the number of exploitable photogenerated electron–hole pairs for photocatalytic reactions.

Briefly, the present study aims at the preparation and identification of MTTF (photo)electrode materials that represent suitable host structures for a subsequent (photo)electrodeposition of plasmonic noble metal nanoparticles. Thus, the layer-by-layer deposition of multi-layered TiO2 onto FTO/glass substrates and its effect on performance, structure and film stability is investigated. At first, single-layered MTTFs were prepared according to an established procedure and were further optimized with respect to a calcination temperature of 550 ◦C. In the second part, the layer-by-layer deposition and characterization of multi-layered films are performed showing that an undiminished activity up to 10 layers is observed. Thirdly, performance and mechanical stability tests are conducted to demonstrate the film's durability. Finally, their suitability as porous hosts for subsequent loadings is briefly demonstrated in terms of pulsed electrochemical incorporation of SPR-active Au-NPs.

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

#### *2.1. Chemicals, Materials*

FTO/glass substrates (Pilkington "NSG TECTM A7", 7 Ω-<sup>−</sup>1, 2.2 mm thickness, manufacturer: NSG Group, Gladbeck, Germany), acetone (Carl Roth, ≥99.9%), Pluronic<sup>c</sup> P123 (Sigma Aldrich, M <sup>≈</sup> 5800 g·mol−1), glycerol (VWR, 99.8%), 1,6-diisocyanatohexane (Sigma Aldrich, ≥98%), abs. ethanol (Fisher Scientific, 99.5%), HCl(aq) (Fisher Scientific, 37 wt.%), Ti(OEt)4 (Alfa Aesar, >99%), de-ionized H2O (18.2 MΩ cm), HAuCl4 × 3 H2O (Alfa Aesar, 48.5–50.25 wt.% Au), HClO4 (Bernd Kraft, 70 wt.%). All chemicals were used without further purification.

#### *2.2. Preparation of Mesoporous TiO*<sup>2</sup> *Thin Films*

The preparation of MTTFs proceeded according to the flowchart depicted in Figure 1a.

**Figure 1.** (**a**) Flowchart of steps involved in the film preparation. Green and blue arrows represent pathways to single- and multi-layered MTTFs, respectively. The pale blue arrow indicates a shortcut by skipping characterization/cleaning of MTTFs. (**b**) Scheme and photograph of a TiO2 thin film. Green and blue areas represent the bare FTO and TiO2 part, respectively. The circle indicates the region where electrodeposition might take place later; but more importantly, where optical spectroscopy and photoelectrochemistry are performed.

Pretreatment: FTO/glass substrates were cut into pieces of 25 mm × 46 mm and cleaned by four consecutive ultrasonication steps (namely in 0.1 M NaOH(aq), 0.1 M HCl(aq), ethanol and acetone; each lasting 10 min at 37 kHz). Cleaned substrates were modified with a so-called "cross-linked P123". This strategy has proven to promote the orthogonal orientation of two-dimensional mesochannels (by providing a surface chemically neutral towards both the hydrophilic (PEO) and hydrophobic (PPO) parts of the triblock-copolymer of P123) and was first demonstrated for MTTFs by the Rankin group [76,77]. For doing so, cleaned FTO/glass slides were dip-coated into an acetone-based solution with equimolar amounts of P123 and 1,6-diisocyanatohexane (both 0.696 mM) where the triol glycerol enabled cross-linkage. This solution was prepared by dissolving 0.404 g P123 and a droplet of glycerol into 100 mL acetone (15 min ultrasonication at 37 kHz), precooling in a refrigerator and addition of 12 mg of the diisocyanate inside a glovebox. It was readily used due to the ongoing isocyanate-hydroxyl polyaddition reaction towards polyurethane and the dip-coating procedure was performed at ambient conditions with a 70 mm·min−<sup>1</sup> immersion rate, 20 s holding time and withdrawal rate of 20 mm·min−<sup>1</sup> with an *IDLAB 4 AC Coater* device.

Deposition (dip-coating): EISA-sols were prepared by dissolving 1.30 g P123 in 15.0 g ethanol (15 min ultrasonication at 37 kHz). After being cooled in the refrigerator for 30 min, 3.06 g of concentrated HCl(aq) were added dropwise under vigorous stirring. The resulting mixture was further ultrasonicated and placed in a refrigerator for 30 min again. The precooled solution was transferred into a glovebox where 4.20 g of Ti(OEt)4 were added dropwise under vigorous stirring. The resulting molar ratio in the EISA-sol was Ti(OEt)4:P123:EtOH:HCl:H2O = 1.00:0.0122:17.7:1.6:5.9. Both the chemicals and their stoichometric ratios were adopted from Choi et al. and are known to promote the formation of hexagonal mesopores [78]. However, little changes regarding the used solvent were made towards thicker films based on the procedure described in [76].

Films were deposited onto cross-linked FTO/glass substrates via dip-coating (*IDLAB 4 AC Coater*) into a teflon-made tank (reservoir dimensions: 38 mm height, 28 mm width, 8 mm depth) with a 70 mm·min−<sup>1</sup> immersion rate, 20 s holding time and a 20 mm·min−<sup>1</sup> withdrawal rate. Air-conditioning allowed operation at T < 15 ◦C and relative humidities >70 %RH. After 2 min of hanging inside the dip-coater, the substrates were taken out for the aging phase. Prior to the deposition, the backsides of the cross-linked substrates were masked with a solvent resistant scotch tape.

Humidity aging, preconsolidation, calcination: As-deposited films were aged at low temperatures (4 ◦C) and high relative humidities (>90%) for 2–3 h. This environment has proven to favor the relatively quick formation of ordered mesostructures [78] and was established by placing substrates in dedicated aging boxes, namely sealable boxes with a reservoir of saturated KNO3(aq) that can be stored inside a refrigerator. To evaporate remaining volatile compounds and complete the polycondensation, the films were stored inside a drying furnace at 100 ◦C overnight. For template removal and network crystallization, the scotch tape was peeled of and preconsolidated films were calcined in air by increasing the temperature with +5 ◦C min−<sup>1</sup> to elevated values of 350–650 ◦C holding for 2 h and natural cooling to room temperature.

Pulsed electrodeposition of gold nanoparticles was carried out in a tree-electrode configuration schematically depicted in Figure 2a. A Pt coil, Ag/AgCl(sat. KCl) and a MTTF were connected to an *AMEL 7050* potentiostat and were used as counter, reference and working electrodes, respectively. A salt bridge was used to separate reference electrode and deposition bath.

**Figure 2.** Pulsed electrodeposition of Au-NPs into MTTFs. (**a**) schematic representation of the used setup and (**b**) adjustable parameters.

An aqueous solution of 10 mM HClO4 and 70 μM HAuCl4 with a pH value of 2.1 was used as deposition bath and prior to the addition of the gold salt, the 10 mM HClO4(aq) was purged with nitrogen for 30 min [50]. To ensure that a defined geometric sample area is exposed to the electrolyte solution, the MTTF was masked with a solvent resistant scotch tape leaving out a circular zone of 2 cm diameter (=3.14 cm2). Electrodeposition was carried out at room temperature on single-layered MTTFs and involved reduction reaction towards elemental gold shown in Equation (1) [79]. Two oxidation reactions at the Pt anode are expectable when considering the used chemicals and applied potentials. They are mentioned in Equations (2) and (3) [79] but are of minor importance here.

The used electrochemical parameters rely on a previous study by Bannat et al. and involved a pulse potential of Upulse = –2.0 V vs. Ag/AgCl, a pause potential of Upause = +0.1 V vs. Ag/AgCl, a pulse duration of tpulse = 3.0 s, a pause duration of tpause = 0.1 s and a rectangular pulse shape [50]. Three different amounts of pulses were applied, namely 1, 8 and 16.

All involved physicochemical conditions are summarized in Figure 14 for better comparison.

$$\text{Cathode:}\quad \text{AuCl}\_4^- + \text{ 3e}^- \rightarrow \text{Au} \downarrow + \text{ 4Cl}^- \tag{1} \qquad \qquad \qquad E\_{\text{red}}^\circ = 1.002 \text{ V}\_{\text{RHE}} \tag{1}$$

$$\text{anode:}\quad \text{2Cl}^-\qquad\rightarrow \text{Cl}\_2\quad+\text{2e}^-\qquad\qquad\qquad E\_{\text{ox}}^\circ=1.358\text{ V}\_{\text{RHE}}\tag{2}$$

$$\text{6H}\_2\text{O} \qquad \rightarrow \text{ O}\_2 \quad + \text{ 4H}\_3\text{O}^+ + \text{ 4e}^- \qquad E^\circ\_{\text{ox}} = 1.229 \text{ V}\_{\text{RHE}} \tag{3}$$

#### *2.3. Characterization of Mesoporous TiO*<sup>2</sup> *Thin Films*

Individual multi-layered MTTFs were characterized by means of UV/Vis and fluorescence spectroscopy as well as photoelectrochemical (PEC) water oxidation prior to a next deposition. Further experiments such as XRD, scanning electron microscopy (SEM) and nitrogen sorption were used for selected final samples only.

UV/Vis transmittance spectra in the range of *λ* = 200–800 nm were recorded with a *Varian Cary 4000* UV/Vis spectrophotometer. Prior to the actual measurement, baseline correction was ensured by measuring a cleaned blank FTO/glass substrate that has been calcined with the same temperature program as the investigated MTTF sample.

Fluorescence emissions were measured in the range of *λ* = 400–600 nm with a *Varian Cary Eclipse* fluorescence spectrophotometer. The excitation wavelength was 380 nm.

Photoelectrochemical (PEC) measurements were performed using an electrochemical workstation with a main potentiostat (*ZAHNER ZENNIUM Pro*) and a secondary potentiostat (*ZAHNER PP211*) to power a UV-LED (*<sup>λ</sup>* = 375 nm, *δλFWHM* <sup>=</sup> ±7 nm, 70 W·m−<sup>2</sup> incident photon flux). The experiments were controlled with a computer and the *THALES XT* software package. A PEC cell (*ZAHNER PECC-2*) was filled with aqueous 1 M NaOH of pH = 13.6 and a three-electrode configuration was established by using a platinum wire, Ag/AgCl(sat. KCl) and the MTTF as counter, reference and working electrodes, respectively. Electrical contact to the films was made with help of copper wires and copper tapes attached to the bare FTO part of each sample. All experiments were performed under front-side illumination, thus the light entered the PEC cell from the electrolyte side through a quartz window. To eliminate interference by external irradiation, the PEC cell was placed into a light exclusion box.

For better comparison towards results from literature, the applied potentials versus Ag/AgCl(sat. KCl) were also translated into potentials relative to the reversible hydrogen electrode (RHE) by using the Nernst equation:

$$E\_{\rm RHE} = E\_{\rm Ag/AgCl(sat. KCl)} + 0.059\,\text{V pH} + E\_{\rm Ag/AgCl(sat. KCl)}^{\circ}\tag{4}$$

$$\text{with } E\_{\rm Ag/AgCl(sat. KCl)}^{\circ} = 0.1976\,\text{V}$$

$$\text{and pH} = 13.6$$

Chronoamperometry (*It*) under UV irradiation was performed with six illumination-dark-periods of 10 s each at a fixed potential of 1.20 V versus RHE. Linear voltammetric scans, or briefly (*IV*) scans, were carried out under UV irradiation or in the dark to ensure that no impurity-related adulterations of the *It* data was present. Accordingly, voltage scans were performed in the range of +0.4 V to +2.25 V versus RHE with a scan rate of 15 mV·s<sup>−</sup>1. Prior to photoelectrochemical measurements of gold-loaded MTTFs, the electrolyte had to be purged for 30 min with nitrogen to replace dissolved air oxygen. This step is necessary to eliminate the reductive dark current related to the oxygen reduction reaction.

Cleaning was necessary to remove characterization-related impurities prior to the deposition of an additional MTTF layer. Conductive adhesive and remaining NaOH from the PEC characterization were removed by acetone and rinsing with de-ionized water, respectively. The cleaning was completed by a short period of ultrasonication in de-ionized water and an overnight drying at 100 ◦C.

Scanning electron microscopy (SEM) was carried out using a *FEI Helios NanoLab 600i*. Prior to the experiment, substrates were cut and affixed onto aluminum sample holders using conductive carbon tape and copper wires before applying conductive silver adhesive. SEM images of both surfaces and breaking edges were recorded with a secondary electron detector in a working distance of 4 mm and with an acceleration voltage of 10 kV and an electron beam current of 0.17 nA.

Combined energy dispersive X-ray spectroscopy and scanning electron microscopy (EDX-SEM) was used for gold-loaded MTTFs. Acceleration voltage and electron beam current were raised to 20 kV and 5.5 nA, respectively. F, Si, Ti, Sn and Au were selected as elements to assign and overlay pictures of the collected EDX mappings were generated with the *EDX Genesis* software.

Grazing incident X-ray diffractograms (GI-XRD) were obtained with a *PANalytical Empyrean* diffractometer with CuK*α* = 1.54 Å radiation at an incident angle of 0.5◦. The Scherrer equation was applied to determine the average crystallite sizes.

Nitrogen adsorption-desorption was measured with a *Micromeritics ASAP 2020* device. Prior to the adsorption measurements, the samples were degassed at 150 ◦C for 2 h at a base pressure of <sup>2</sup> × <sup>10</sup>−<sup>6</sup> mm Hg. For each measurement, two identically prepared MTTF substrates were placed inside a special substrate holder. Surface areas calculated according to the formalism established by Brunauer, Emmett and Teller (BET) from two selected partial pressures (0.005 p/p0 and 0.038 p/p0) and were related to the MTTFs' geometric surfaces due to the very low and in general hard to determine mass of individual films. The two-point calculation was selected in all cases since it was only possibly to collect full sorption isotherms for samples with >4 layers. Pore size distributions were calculated from the desorption branches of the isotherms using the Barrett–Joyner–Halenda (BJH) model.

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

#### *3.1. Single-Layered Films*

Figure 3 shows top-view SEM images of four single-layered MTTFs after different calcination temperatures between 400–550 ◦C. After the 400 ◦C treatment, the desired well-ordered mesoporous structure was obtained which is in good agreement with the protocol by Nagpure et al. [77]. With higher temperatures, the mesostructure becomes increasingly deteriorated, involving wall collapses and particle agglomeration. This observation that a formerly pristine network consisting of closely packed micelles within an amorphous TiO2 matrix experiences structural deterioration (and formation of worm-like structures) upon air calcination has been commonly observed for MTTFs and is explained in terms of mass migration that accompanies the proceeding crystallization [62,80]. At the macroscopic scale, higher temperature treatment affected the structure of the film leading to cracks encompassing larger regions of FTO which appear much brighter than the TiO2 in the SEM images. The measured thickness was around 450 nm for a single-layered film calcined at 550 ◦C (shown later in Figure 7). This is almost twice as much as the ≈250 nm mentioned by Rankin et al. [77] who reported the use of an identical EISA-sol composition as well as similar aging/thermal treatment. Unfortunately, these authors did not mention their dip-coating conditions (most importantly the used withdrawal rate). Other thicknesses were not determined within this study, but it was shown elsewhere that the thermally induced structural collapse is accompanied with a shrinkage of films [81].

**Figure 3.** Influence of calcination temperature on mesofilm morphology as observed by top-view SEM. Higher calcination temperatures cause mesostructure deteriorations and the formation of cracks at T > 500 ◦C.

To analyze phase composition and crystallinity as a function of calcination temperature, grazing incidence X-ray diffraction was used. The obtained diffractograms for single-layered MTTFs treated at different temperatures between 350–550 ◦C are shown in Figure 4. Reference patterns of anatase TiO2 [82], rutile TiO2 [83] and SnO2 [84] (as a representative for FTO) are included for reflex assignment and a diffractogram of a blank FTO/glass substrate was measured to identify the contribution of the FTO to the recorded diffractograms.

**Figure 4.** Normalized GI-XRD patterns of single-layered MTTFs calcined at different temperatures. Reference patterns [82–84] and selected hkl values are displayed for better interpretation.

The observed reflexes in the blank FTO diffractogram can be explained with the used reference pattern of SnO2 in its rutile modification [84]. It is worth mentioning that the (200) orientation at 37.5◦ of 2*θ* is preferential in the used FTO/glass substrates. After a 350 ◦C thermal treatment, a weak anatase (101) reflex at 25.3◦ of 2*θ* is already visible but FTO-related reflexes dominate (e.g., those related to the (110), (101), (200) and (211) planes at 26.3◦, 33.7◦, 37.5◦ and 51.3◦ of 2*θ*, respectively). Those reflexes are part of all diffractograms of single-layered MTTFs as result of the small TiO2 thickness. However, they start to diminish with additionally deposited layers (shown later in Figure 8). With higher calcination temperatures, the anatase phase becomes more crystalline as more intense (still broad) signals are observed (see insets). Average crystallite sizes for the most pronounced anatase reflex, namely (101), were calculated using the Scherrer equation and the results are summarized in Table 1. The observed average grain size increases for elevated calcination temperatures from 16 nm to 28 nm for samples treated at 400 and 550 ◦C, respectively are caused by mass migration and particle agglomeration during crystallization [62,80]. This agrees with the structural deterioration observed from the SEM imaging. No formation of rutile was observed in the investigated temperature range, indicating anatase is the exclusive phase in the MTTF. For instance, the most prominent rutile reflex, namely the (110) plane expected at 27.4◦ of 2*θ*, was not formed at all (see upper inset in Figure 4). However, also less pronounced rutile reflexes such as its (101), (111), and (301) plane (expected around round 36.0◦, 41.2◦ and 69.0◦ of 2*θ*, respectively) were absent. This is in accordance with previous reports, where highly thermally stable MTTFs were investigated and the formation of rutile was not observed for temperatures below 700 ◦C [60,85,86].

**Table 1.** Analysis of anatase (101) GI-XRD peak broadening of single-layered MTTFs calcined at different temperatures.


The photoactivity of single-layered films was measured by means of photoelectrochemical water oxidation under UV illumination. Linear voltammetric scans of thermally treated MTTFs at different temperatures are shown in Figure 5.

**Figure 5.** Influence of calcination temperature on photoelectrochemical water oxidation activity. (**a**) Voltammetric scans and (**b**) selected photocurrent densities for each corresponding calcination program.

The observed photocurrent densities strongly depend on the chosen calcination program and reach a maximum for the 550 ◦C treated samples. Higher calcination temperatures result in decreased photocurrent densities, which is likely related to a larger degree of mesostructural collapse that can no longer be overcompensated by simultaneously improving crystallinity of the MTTF. Such trade-offs between improved optoelectrochemical properties on the one hand and loss of accessible active sites on the other are commonly observed for MTTFs [58,62]. The optimal temperature for a given photocatalytic device not only depends on the involved chemicals and their procession but also on its later purpose. In the context of this study, the calcination program with a holding temperature of 550 ◦C was chosen for the preparation of multi-layered films. This thermal treatment gives highest photoactivities due to improved crystallinity and a still accessible (albeit deteriorated) pore structure.

#### *3.2. Multi-Layered Films*

Multi-layered films with up to 10 layers have been prepared in terms of an "interrupted" layer-by-layer deposition which is schematically depicted in Figure 1 on page 3. Accordingly, a typical deposition cycle consisted of the dip-coating deposition, followed by a high humidity aging treatment at low temperatures, an overnight preconsolidation at 100 ◦C and finally the calcination at the optimized temperature of 550 ◦C. It is emphasized that the term "interrupted" derives from the fact that individual films were characterized prior to the next layer deposition. Such a characterization usually involved UV/Vis and fluorescence spectroscopy and—more importantly—photoelectrochemical characterization in terms of water oxidation in a highly alkaline media. This interim characterization/cleaning is meant to ensure (i) the sole effect of additional layers by avoiding sample variations and (ii) the robustness of the films without losing activities. However, its potential effect on the observed photoelectrocatalytical and structural properties must be considered in context of our experimental findings.

Figure 6 illustrates the optical properties of the obtained MTTFs. Part (a) shows a side-by-side photograph of selected multi-layered substrates. Increasing opacities in the visible range already indicate additional amounts of TiO2 with each new layer.

**Figure 6.** Photograph and optical spectra of multi-layered MTTFs calcined at 550 ◦C. (**a**) Photographs of selected samples, (**b**) UV/Vis transmittance spectra, (**c**) Tauc plots, and (**d**) fluorescence spectra.

This observation is confirmed in terms of the UV/Vis transmittance spectra shown in Figure 6b where the increasing opacity is apparent in both the UV and visible range. While reduced UV transmittance is entirely intentional and the decisive motivation for multi-layer formation in the first place, the optical opacity might be beneficial for intended future incorporation of plasmonic nanoparticles. The intense back-scattering of visible photons within the film implies that more visible photons could be consumed for surface plasmon resonance.

Figure 6c shows the Tauc plots and optical band gaps. While single- and double-layered films show slightly higher band gaps of ≈3.4 eV and ≈3.3 eV respectively, films consisting of four and more layers display values close to the expected 3.2 eV of TiO2 in its anatase modification [2]. This observation of slightly larger optical band gaps for 1L and 2L could hint at quantum size effects that disappear with particle growth due to the additional thermal processing of multi-layered films. Analysis of peak broadening from the Gi-XRD pattern (Figure 8) however, does not reveal increased grain sizes. In fact, the average crystallite sizes for the most dominant (101) reflexes remain relatively unaltered at 27 nm. However, these grains might consist of coalesced (aggregated) small nanoparticles, which only sinter completely during additional thermal treatment.

The deposition of additional TiO2 is furthermore confirmed by fluorescence spectroscopy as increasingly intense fluorescence emissions indicate in Figure 6d. These results are only considered to be qualitative indicators for the presence of higher TiO2 amounts and any further interpretation is not considered here. Briefly, the experience that one and the same sample occasionally shows noticeable fluctuations of fluorescence intensities gives reason for this perception.

A more reliable evidence for the additional deposition of TiO2 is the development of film thickness which was determined by SEM imaging of breaking edges and is shown in Figure 7.

**Figure 7.** Development of film thickness as function of deposited MTTF layers. (**a**) Breaking edge SEM images of multi-layered MTTFs calcined at 550 ◦C and (**b**) their corresponding thickness-versus-layer plot.

The development of 0.45, 0.75, 1.05, 1.4, 1.7 and 2.15 μm for 1-, 2-, 4-, 6-, 8- and 10-layered films, respectively is in fact not as clearly indicative for a linear correlation between film thickness and number of deposited layers as claimed by other authors [47,73–75,87,88]; however, a subsequent gain of TiO2 is evident. Considering the development from 2–10 layers, a linear increase (of ≈150–200 nm per layer) can be determined. Thus, only the first two layers with their above-the-average thickness exhibit an exception.

One possible explanation for the exceptionally high thickness of 0.45 μm of the first layer could be that the surface initially deposited onto (FTO) displays a significantly different environment compared to that of all subsequent layers (mesoporous TiO2). Another explanation regarding the second layer could be related to the 550 ◦C air calcination and total oxidative removal of surfactants that fully exposes the porous structure of the first layer. This in turn might allow another EISA-sol infiltration compared to higher-layered MTTFs later. In fact, the Grosso group reported a layer-by-layer synthesis route involving an interposed dip-coating into a diluted ethanolic surfactant solution (just these two chemicals) to allow pore refilling of thermally exposed mesopores prior to subsequent dip-coating into EISA-sol [73]. Such a treatment allowed a linear increase of film thickness. For more than two layers

indeed, a more or less linear dependence is observed indicating an "equilibrated sub-surface" for the subsequent deposition cycle.

Figure 8 shows GI-XRD patterns of the obtained multi-layered MTTFs together with reference patterns.

**Figure 8.** Normalized GI-XRD patterns of multi-layered MTTFs calcined at 550 ◦C. Reference patterns [82–84] and selected hkl values are displayed for better interpretation.

In case of single- and double-layered films, FTO-related reflexes are dominant, but they quickly decline with further layers. Still noticeable FTO reflexes such as (101) and (211) (at 33.7◦ and 51.3◦ of 2*θ*, respectively) for 10-layered films could possibly be related to the porous and non-compact TiO2 structure. All multi-layered MTTFs appear to consist entirely of the anatase phase. An improved crystallinity in terms of sharper anatase reflexes as a result of the multiple thermal treatments was expected, since improved crystallinity of anatase within the bottom layers (resulting from repeated thermal processing with the additional layer) had already been demonstrated for a relatively lower interim temperature treatment at 350 ◦C [74]. The authors performed incident angle dependent GI-XRD scans between 0.1–0.4◦ on a five-layered MTTF, thus manipulating the X-ray's penetration depth and obtain insights on the crystallinity changes at the bottom layers. Similarly, the MTTFs investigated in our work may have improved bottom-layer crystallinity as the intermediate thermal treatment was higher (550 ◦C) than in the previous report [74]. The average crystallite sizes obtained from the most dominant (101) reflexes remain relatively unaltered at 27 nm (see Table 2).

**Table 2.** Analysis of anatase (101) GI-XRD peak broadening of single-layered MTTFs calcined at different temperatures.


The specific surface area of the films was investigated with nitrogen sorption experiments and BET analysis. The BET surface areas were derived from two selected partial pressures (0.005 p/p0 and 0.038 p/p0) and were related to the MTTFs' geometric surfaces due to the very low and in general difficult to determine mass of individual films. Figure 9a shows the calculated BET surfaces as a function of the number of deposited MTTF layers.

**Figure 9.** Nitrogen adsorption-desorption isotherms of multi-layered MTTFs calcined at 550 ◦C. (**a**) The calculated BET surface areas as a function of the number of deposited MTTF layers and (**b**) nitrogen sorption isotherms and pore size distribution of selected substrates.

Surface areas increased almost linearly up to four layers where it reached 188 m2·m−2. Beyond that almost no changes were observed. However, the trend of quickly saturated surface areas is in very good agreement with similar observation by Procházka et al. [74]. They studied P123-templated MTTFs with 1–5 layers which were obtained via layer-by-layer dip-coating onto FTO/glass and with 2 h lasting thermal treatment at 350 ◦C prior to the next deposition cycle (but without interruption by characterization/cleaning). Their finding of an unchanged surface area for more than 3 layers was explained by the compensation of two opposing processes: added surface area per deposition and lost surface area due sintering of bottom layers upon thermal treatment [74]. This very explanation is likely to apply here as well. Moreover, they documented two additional features: the stepwise improvement of the crystallinity of anatase nanoparticles within bottom layers during repeated thermal processing and the formation of denser crust films surfaces as a result of one-dimensional constrained sintering. Although these effects might apply here as well, they are not clearly pronounced in the collected XRD and SEM data, respectively as stated before. Figure 9b presents full nitrogen sorption isotherms of selected multi-layer films (4 L, 8 L, 10 L) showing mesopore condensation. An inset shows pore size distribution from desorption isotherms. With additional layers (and additional thermal treatments), the pore size distributions shift to smaller values indicating pore narrowing due to "wall thickening" during the subsequent sol infiltration process.

To correlate the structural properties of MTTFs involved in the present study with their activity, PEC water oxidation in a highly alkaline media was used as a test reaction. More precisely, chronoamperometric experiments were carried out under controlled chopped-light UV illumination conditions at 1.2 V versus RHE. For a better reliability, two identically prepared samples were tested. The corresponding experimental results are depicted in Figure 10.

**Figure 10.** Photocurrent densities (during PEC water oxidation) of multi-layered MTTFs calcined at 550 ◦C as a function of the number of deposited MTTF layers. (**a**) The corresponding chronoamperometric measurements. (**b**) Average values of the two sets of samples used during the chronoamperometric experiment. A green line indicates the average activity increase for the first eight layers. (**c**) derivative representation of (**b**) to emphasize the specialty of both 10L films.

From the raw data depicted in Figure 10a it is clearly visible that the photocurrent increases with the number of layers. The similarity of results for a pair of samples with the same number of layers (denoted as dashed and solid lines of same color) underlines their good reproducibility. Figure 10b shows the average photocurrents of each pair as a function of the number of layers. It becomes clearly visible that the activity increases almost linearly in the range of 1–8 layers, leading to an average increase of ≈15.5 <sup>μ</sup>A·cm−<sup>2</sup> per layer. To our surprise, the 10-layered MTTFs exceed this trend by exhibiting a significantly higher photocurrent per layer value of ≈19.4 <sup>μ</sup>A·cm−<sup>2</sup> layer−1. Figure 10c further emphasizes this circumstance in terms of a derivate plot.

Without focusing too much on the exceptionally active 10-layered films, the observation of an undiminished activity increases up to 10 L (if even beyond was not tested) has to be explained in the context of the already discussed findings: One is that the thickness increases almost linearly (at least if single- and double-layered films are not considered). Another is that the surface area is almost constant (again if single- and double-layered MTTFs are not considered). Therefore, the enhanced activity needs to be explained by access to more active sites derived from improved optoelectrochemical properties and decreased recombinational losses in bottom layers near the FTO. Additionally, the higher rates of charge carrier formation due to stronger absorption of thicker materials contribute as well. It should be pointed out that the same thickness-versus-surface-area behavior was observed and successfully explained in terms of increased bottom-layer crystallinity before [74].

A further possible explanation could be related to the "interrupted" layer-by-layer deposition. In fact, it may induce structural altering as result of characterization/cleaning steps between each deposition. Figure 11 shows top-view SEM images of single- versus ten-layered MTTFs at four different magnifications.

**Figure 11.** SEM top-view image of (**a**) single- and (**b**) ten-layered MTTFs at different magnifications showing that an extended network of microscopic cracks is formed in case of the multi-layered sample.

Single-layered films show some isolated cracks which expose the FTO back contact (white lines in the images). The topology of the ten-layered films, however, was not expected. SEM imaging reveals the formation of an extended network of small cracks which are much smaller in size but much more prominent than those for single-layered films. Topologically, these structures resemble a hierarchically macro-mesoporous systems that were prepared by controlled phase separation and surfactant templating [89]. Such a structure provides improved transport properties as the electrolyte's infiltration/diffusion into the MTTF becomes facilitated and therefore is likely to contribute to the enhanced activity as well. At this point it must be determined in terms of uninterrupted control experiments whether this extensive crack formation is related to the characterization/cleaning step between deposition or if it is caused by the multiple calcination steps.

#### *3.3. Multi-Layered Films Stability*

To gain insights on the chemical and mechanical stability of the films, we have performed different sets of stability measurements. Electrochemical stability tests were performed on 10-layered MTTFs and involved the steps of disassembling, cleaning, drying and assembling between each chronoamperometry measurement. These measurements were carried out for ten consecutive days and—for better reproducibility—two sets of samples treated at 450 ◦C and 550 ◦C (two films each) were used (see Figure 12a). At the end of the tenth measurement, the sample "550 ◦C\_A" was mechanically rubbed with a cotton bud (Q-Tip<sup>c</sup> ) (see Figure 13a).

**Figure 12.** (**a**) Photocurrent stability tests of four 10-layered MTTFs (2× 450 ◦C, 2× 550 ◦C) recorded at ten different days (disassembling, cleaning, drying, assembling between each measurement). Photocurrent densities were derived from chronoamperometric experiments and taken after 32.5 s (right in the middle of an illumination phase). (**b**) Three-dimensional representation of the chronoamperometric results that correspond to a the "550 ◦C\_A" sample from (**a**).

As shown in Figure 12a,b, all films deliver stable photocurrents during the first 10 experiments. Films treated at 450 ◦C show values around ≈<sup>140</sup> <sup>μ</sup>A·cm−<sup>2</sup> and films treated at 550 ◦<sup>C</sup> ≈ <sup>200</sup> <sup>μ</sup>A·cm<sup>−</sup>2. To set this into perspective, stable photocatalytic activities of MTTFs (loaded with Pt NPs) were reported previously by Ismail for both long term and repeated measurements of the photocatalytic gas-phase oxidation of acetaldehyde [75].

After the rubbing test of "550 ◦C\_A" (red diamond), still comparable photocurrents are observed with very slight variations, suggesting the films are quite robust and intact. The slight increase of photocurrents at measurement 11 might be related to the use of a freshly prepared new electrolyte solution. When considering the photographs taken before and after rubbing, the applied mechanical stress did not result in macroscopic changes of the films (see Figure 13a). Corresponding SEM micrographs, however, show that the structural integrity of the films is affected to an extent as scratch marks and microstructural changes reveal. Overall, there seems to be no significant material abrasion since the activity stays nearly unaffected which confirms the good adhesion. However, MTTFs cannot withstand a scotch tape adhesion test (not shown).

In hope for an accelerated film preparation process, two sets of comparative experiments were carried out: In the first set, we have prepared multi-layered films without an intermediate temperature step and tested the mechanical stability in a similar way. Unfortunately, these films are not mechanically stable at all and can be peeled off easily (see Figure 13b). It is worth mentioning that a very comparable

synthesis route (also without intermediate calcination) was recently reported by Rankin et al. They used MTTFs with up to five layers as negative electrodes in Li-ion batteries and did not report stability problems [88]. In another attempt towards faster production of thicker films, a more viscous EISA-sol was used by having a lower amount of the solvent EtOH in the initial EISA-sol. The obtained MTTFs exhibited the same problem of bad adhesion (photograph not shown).

**Figure 13.** Two different samples before and after the "rubbing test". (**a**) Photograph and SEM top-view images of the 10-layered MTTF "550 ◦C\_A" from Figure 12. Although its microscopic structure is affected by the rubbing, this sample has delivered stable photocurrents prior and after this procedure. (**b**) Photograph of a 3-layered MTTF that had been calcined only once to save energy and time, namely after the drying phase of its third and last deposition cycle. This film is not mechanically stable.

#### *3.4. Proof of Concept: SPR-Active Gold Nanoparticles Inside MTTFs*

In this brief subsection, the suitability of our MTTFs as porous host systems for the incorporation of SPR-active noble metal particles is investigated. As a test reaction, the cathodic pulsed electrodeposition (pulsed-ED) of gold nanoparticles was used.

The pulsed-ED technique provides the benefit of being simple to operate but still versatile, since it allows introduction of Au nanoparticles of different size/shape by use of different deposition potentials, pulse/pause sequences, pulse numbers etc. [50,90]. Therefore, pulsed-ED enables both fine-tuning of plasmonic properties and intense contact areas between TiO2 and Au the same time. Individual single-layered MTTFs were treated with three different amounts of pulses, namely 1, 8 and 16.

Figure 14 shows the results of the experiments. Figure 14a summarizes the deposition bath composition and electrochemical pulse parameters. Optical analysis is shown in Figure 14b, where both a photograph and the corresponding UV/Vis transmittance spectra show the surface plasmon resonance features related to the introduced gold particles. Specifically, an SPR band is present for all three Au-modified samples and is in the green region of the visible electromagnetic spectrum around wavelengths of 530 nm. The tendency towards slightly increased wavelengths after more applied pulses goes along with the expected particle growth and the size (and shape) dependence of SPR features [57]. Furthermore, it should be pointed out that the porous TiO2 host structure confines the growth of incorporated gold inside the film—with the limit being dendritic Au structures as partial replica of the original pore system [50].

**Figure 14.** Properties of SPR-active gold-modified MTTFs obtained after three different amounts of applied cathodic pulses. (**a**) physicochemical conditions and pulse parameters during pulsed electrodeposition, (**b**) UV/Vis transmittance spectra and photograph, (**c**) chronoamperometric results (dashed and solid lines before and after pulsed-ED, respectively) with inset showing relative increase of photocurrents, (**d**) fluorescence spectra with inset showing relative decrease of emission intensity, (**e**) cross-sectional SEM images and corresponding EDX mapping of two selected samples, (**f**) top-view SEM image of most active sample near a cracked part showing Au-NPs.

The activity before and after the pulsed-ED was determined in terms PEC water oxidation under UV illumination and the corresponding chronoamperometric measurements are shown Figure 14c. The data show that exposure to 1, 8 and 16 pulses leads to highest, second highest and lowest photocurrents, respectively with corresponding relative photocurrent increases of 87%, 62% and 1%

(inset of Figure 14c). These increased activities are explained by the formation of a Schottky barrier at the Au–TiO2 interface which causes spatial separation of photogenerated charge carriers involving transfer of electrons towards the noble metal [41,49]. The consequence is a reduced rate of charge carrier recombination which extends the amount of harvested charge carriers and their participation in photoelectrochemical reaction—ultimately leading to the observed improved photocurrents. This is astonishingly well confirmed in terms fluorescence emissions whose relative decreases of 86%, 68% and –2% (shown in Figure 14d) go almost perfectly hand in hand with the relative photocurrent increases. Both experiments prove that the formation of intense contact areas between the noble metal and the semiconductor is possible with the chosen pulsed-ED approach. The observed decreased activity for additional pulses might be a result of pore blocking or screening due to an excessive amount of gold. In fact, MTTFs that experienced an extended electrochemical treatment involving 16 pulses have an extensive deposition of Au on the film surface (not shown).

To further demonstrate the successful incorporation and to possibly obtain first insights regarding spatial distributions of Au-NPs, SEM-EDX mapping was performed on breaking edges of the samples exposed to 1 and 16 pulses. The corresponding SEM images and SEM-EDX overlays are shown in Figure 14e. Regions related to glass, FTO and TiO2 can be easily identified and distinguished. More importantly, the incorporation of Au-NPs is indicated but information related to their spatial distribution (and their structural properties including size distribution) cannot be extracted from those images. It appears from the presented overlays that Au is enriched at the film surface. SEM-EDX mapping confirms the incorporation of Au but for more detailed information of spatial and size distribution, selected thin Au/MTTF filaments should be analyzed by transmission electron microscopy (TEM) which is intended for future investigation. Finally, a top-view SEM image of the most active sample is shown in Figure 14f. It was measured nearby a crack in the MTTF and the incorporation of Au-NPs is confirmed once more as they appear as small dots with much brighter contrast compared to the TiO2 or the FTO.

#### **4. Conclusions**

Multi-layered mesoporous TiO2 thin film (photo)electrodes with up to 10 layers were prepared by an EISA method and repeated dip-coating/aging/drying/calcination cycles ("interrupted" layer-by-layer deposition). The experimental conditions were derived from an existing route towards single-layered films and were optimized with respect to the calcination temperature.

Each additional layer was accompanied with an almost linear increase of activity due to increased amounts of immobilized TiO2 and access to more active sites as a combined effect of increased surface area and better crystallinity. While additional TiO2 causes stronger absorption of UV photons and thus more photogenerated e−–h<sup>+</sup> pairs, accessible porous structure plus improved crystallinity allows their collection and participation in photochemical reactions. Another contribution is related to improved transport properties due to an extended network of cracks which likely derived from the interim characterization/cleaning steps. These combined effects explain the enhanced activity. Since 10-layered films show the highest photocurrents in this study, there is no indication of an already saturated activity which implies that the it can be further maximized with additional layers.

Moreover, the 10-layered films possess both a good photoelectrocatalytical and mechanical stability making them suitable candidates for future modifications. It was further shown that two faster routes towards thicker films lead to poor adhesion (an easy peeling off) rendering them useless for a further usage.

Initial results show the suitability of MTTFs as porous host substrates. Plasmonic gold nanoparticles were incorporated via pulsed electrochemical deposition and the presence of Au-NPs was confirmed (SEM-EDX, UV/Vis spectroscopy). All nanocomposites showed improved activities under UV illumination compared to pure MTTFs which is in excellent agreement with decreased fluorescence emissions. These observations are explained by reduced h+–e<sup>−</sup> rates the Au–TiO2 interface (Schottky barrier). Further study is in progress to get in depth insights at structural properties of noble metal deposits, visible-light responses and optimum number of layers of the MTTF host.

**Author Contributions:** Conceptualization, E.G., D.H.T. and M.W.; Methodology, E.G. and D.H.T.; Software, E.G.; Validation, E.G. and D.H.T.; Formal Analysis, E.G. and D.H.T.; Investigation, E.G.; Resources, M.W.; Data Curation, E.G. and D.H.T.; Writing—Original Draft Preparation, E.G.; Writing—Review and Editing, E.G. and D.H.T.; Visualization, E.G.; Supervision, M.W.; Project Administration, M.W.; Funding Acquisition, M.W.

**Funding:** This research was funded by the Deutsche Forschungsgemeinschaft (DFG) within the SPP 1613 (WA 1116/28-1) and INST 184/154-1 FUGG and the Bundesministerium für Bildung und Forschung (BMBF) within the project *PROPHECY* (*PRO*zesskonzepte für die *PH*otokatalytische CO2-Reduktion verbunden mit Lif*E*-*CY*cle-Analysis).

**Acknowledgments:** The authors further thank Lea Mohrmann who was significantly involved in the preparation of multi-layered MTTFs.

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

#### **References**


c 2019 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Fully Reversible Electrically Induced Photochromic-Like Behaviour of Ag:TiO2 Thin Films**

#### **Stavros Katsiaounis 1, Julianna Panidi 1,2, Ioannis Koutselas 1,\* and Emmanuel Topoglidis 1,\***


Received: 23 December 2019; Accepted: 28 January 2020; Published: 3 February 2020

**Abstract:**A TiO2 thin film, prepared on fluorine-dopedindium tin oxide (FTO)-coated glass substrate, from commercial off-the-shelf terpinol-based paste, was used to directly adsorb Ag plasmonic nanoparticles capped with polyvinylpyrollidone (PVP) coating. The TiO2 film was sintered before the surface entrapment of Ag nanoparticles. The composite was evaluated in terms of spectroelectrochemical measurements, cyclic voltammetry as well as structural methods such as scanning electron microscopy (SEM), transmission electron microscopy (TEM) and X-ray diffraction (XRD). It was found that the Ag nanoparticles are effectively adsorbed on the TiO2 film, while application of controlled voltages leads to a fully reversible shift of the plasmon peak from 413 nm at oxidation inducing voltages to 440 nm at reducing voltages. This phenomenon allows for the fabrication of a simple photonic switch at either or both wavelengths. The phenomenon of the plasmon shift is due to a combination of plasmon shift related to the form and dielectric environment of the nanoparticles.

**Keywords:** TiO2 films; Ag nanoparticles; optical properties; spectroelectrochemistry; cyclic voltammetry; surface plasmon

#### **1. Introduction**

In recent years there has been significant interest in optically transparent electrodes, due to their range of applications, including solar cells, light-emitting diodes and printable electronics [1,2]. Mesoporous (mp) nanocrystalline titanium dioxide (TiO2) films are optically transparent for wavelengths greater than 390 nm due to their wide energy band gap, at ca. 3.2 eV [3]. The TiO2 films comprise a rigid, porous network which is built with 10–40 nm nanocrystalline TiO2 nanoparticles. These films usually exhibit pore sizes between 5 and 20 nm, sufficiently large for dye molecules [4], metal nanoparticles [5–8], biomolecules [9], gases [10], quantum dots [11,12] and perovskites [13] to diffuse throughout their porous structure. Their surface area is typically much greater (by up to 1000 times) than their geometric area [9]. In addition to their optical transparency and high surface area, these films exhibit good chemical stability, excellent optoelectronic properties and electrochemical activity at potentials above their conduction band edge. Therefore, TiO2 films have been utilized in many applications such as photovoltaics [4], electrochromic windows and displays [14], antireflective coatings [13,15], batteries [14,16], touch screens [17], light-emitting diodes [18], supercapacitors [15,19], photocatalysis and photoelectrochemistry [20–23] or spectroelectrochemistry [9] applications, sensors [24–26] and biosensors [9].

The TiO2 films are highly photosensitive and exhibit optically induced properties. They are also applicable in non-linear optics and optical devices [27,28]. However, one of their main disadvantages is that their large energy band gap (*E*<sup>g</sup> = 3–3.4 eV), that lies in the ultraviolet (UV) region, limits their optical response in the visible region of the electromagnetic spectrum. In order to increase their absorption in the visible spectral range, the introduction of active absorbing units is required for optoelectronic devices. In fact, a number of studies over the last few years show that this can be achieved by incorporating noble metal nanoparticles with plasmonic effect in the TiO2 matrix [27,29,30]. Plasmons are collective oscillations of electrons residing in unfilled energy bands and generally appear as pronounced resonances in the optical absorption spectra of metallic nanoparticles. This photochromic effect adds an interesting new aspect to the rich optical behaviour exhibited by silver nanoparticles (AgNPs). This effect can be studied among others in light scattering and absorption, non-linear signal enhancement and electroluminescence. Plasmonic nanostructures with increased and tunable optical absorption are used in various electronic devices, such as in thin solar cells through efficient scattering of the incident light in semiconducting absorber. Applications of plasmonic materials can also be considered for photochromic materials, which can reversibly change their colour under illumination or applied bias [31]. Lastly, plasmonic nanomaterials have also been proposed for a wide range of applications such as information storage and large-scale displays [32].

Multicolour photochromism was reported in nanocomposite Ag-TiO2 films when these were prepared photocatalytically using a sol-gel route and consisted of AgNPs embedded in anatase TiO2 [33]. The photochromic effect of the composites relied on burning a reversible spectral hole in the plasmon band.

Noble metal nanoparticles, such as AgNPs, exhibit an absorption band in the visible region of the spectrum caused by the surface plasmon resonance (SPR), which occurs at a different frequency from that of the bulk plasmon. The resonance wavelength of the SPR in AgNPs and its intensity are extremely dependent on the particle's environment (dielectric constant and interparticle distance) as well as on their geometry, size and shape. The incorporation of these nanoparticles into the TiO2 matrix will extend their utility and device applications. Ag-TiO2 films exhibit photochromic properties and, therefore, could be used for information storage, displays, smart windows and switches. Additional to these applications, the optically transparent semiconductor TiO2 is used to carry out direct spectroelectrochemical experiments of molecules, such as redox proteins which were seen to be electrochemically active and changes in optical spectra were correlated with changes in applied potential [9].

Herein, we present the use of AgNPs as a simple electrically induced photonic switch when they are deposited on mesoporous (mp) TiO2 films. The electrochromic behaviour of the Ag-TiO2 nanocomposite film was characterized by cyclic voltammetry (CV) and spectroelectrochemistry (SEC). The crystallinity of the film was characterized by X-ray diffraction (XRD) and the surface morphology was examined by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The optical properties of the composite films were investigated via ultraviolet–visible (UV–Vis) spectroscopy. The optical properties and their morphology revealed a hybrid material whose plasmon can be tuned via the application of external bias. Furthermore, the fabrication of a simple photonic switch (on a rigid support) was attempted by assessing the electrochromic behaviour of the Ag-TiO2 films by the application of controlled voltages using SEC. The novelty of the present work is based on the simple and straightforward preparation conditions of both the mp TiO2 layer and AgNPs layer, which provide an optically interesting material. The SPR effect of AgNPs on mesoporous TiO2 films could be influenced by charge transfer and local electric field enhancement. In the charge transfer mechanism, the SPR excites the electrons in the Ag nanoparticles, which are transferred to the conduction band of TiO2, leaving a "plasmonic hole" in the metal nanoparticle [34].

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

#### *2.1. Materials*

Commercial 18NRT TiO2 paste with average final nanoparticle size of 20 nm was purchased from Dyesol (Elanora, Australia) and used without any further purification. Fluorine doped tin oxide-coated (FTO) glass with resistance of 15 Ω/sq was purchased from Hartford Glass (IN, US). Sodium dihydrogen orthophosphate (0.01 M) was used to prepare the supporting electrolyte, and its pH was adjusted to 7 using NaOH. All other reagents were of chemical grade. All aqueous solutions were prepared in distilled, deionised water of resistance R = 18 MΩ cm. Silver nitrate (AgNO3, MW:169.87) and polyvinylpyrrolidone (PVP, MW:10000) were supplied by Sigma Aldrich Chemie GmbH (Taufkirchen, Germany) as well as Na2S in form of solid platelets. P25 nanotitania powder was commercially acquired from Degussa.

#### *2.2. Mesoporous TiO2 Film Electrodes Preparation*

Dyesol TiO2 nano-product (18NRT) was used to prepare thin TiO2 films on FTO glass slides. The slides were first cleaned in a detergent solution using an ultrasonic bath for 15 min, followed by rinsing with de-ionised (DI) water and ethanol. TiO2 was deposited on the substrate via the doctor-blade technique [9], by masking the glass substrates with tape which enabled the control of the thickness and the width of the area spread. The films were then allowed to dry for 20 min (evaporation of the solvent) before being sintered for 20 min at 450 ◦C. The resulting TiO2 films were then cut in 1 cm2 pieces.

#### *2.3. P25 Film Preparation*

P25-TiO2 suspension was prepared from the P25 nanotitania powder, consisting of 80% anatase and 20% rutile which is manufactured by flame synthesis. An aqueous suspension of P25-TiO2 was prepared by mixing 6 g of P25 TiO2 powder, 60 μL of acetylacetone, 4 mL of water, 15 mL of ethanol, 1 mL of acetic acid and 60 μL of Triton X-100. The preparation of films on FTO glass is the same as with the Dyesol paste. For completeness, P25 have exhibited same properties to those prepared with the Dyesol product, with the exception that the Degussa powder leads to film with large scattering, thus, rendering it harder to measure its optical properties.

#### *2.4. Preparation of Silver Nanoparticles (AgNPs) and TiO2-AgNPs Films*

We dissolved 0.01 g/mL of PVP at room temperature in triply distilled H2O (10 mL), to which 100 μL of 0.5 M AgNO3 solution were slowly added under stirring for one hour. The solution was kept for 24 hours after which its colour darkened. Similar reactions were achieved by an extra addition of equal volumes of 0.5 M Na2S aqueous solution to the AgNO3:PVP mixture. TiO2 films were immersed in AgNO3:PVP solution and kept for 24 hours. This procedure enables the reduction of Ag<sup>+</sup> [35], its agglomeration to nanoparticles as well as the prevention of further agglomeration to very large nanoparticles [36].

#### *2.5. Film Characterization*

The adsorption onto the TiO2 films was monitored by recording the UV–Vis absorption spectra of the immobilized films at room temperature using a Shimadzu UV-1800 spectrophotometer. Contributions to the spectra from scatter and absorption by the TiO2 film alone were subtracted by the use of AgNPs-free reference films. Prior to all spectroscopic measurements, the films were removed from the immobilization solution and rinsed in a buffer (methanol) solution several times to remove non-immobilized nanoparticles or excessive AgNO3 solution. The photocatalytic process can be associated with the Ag-TiO2 films, where ultra bandgap irradiation of titania generates an electron-hole pair, with possibility to reduce silver ions at the surface of the titania to silver metal. The conditions under experiments were free from UV irradiation. Typically, the Ag-TiO2 prepared films appeared dark brown black in transmitted light. Chemical and thermal reduction have also been used to prepare Ag-TiO2 films which, however, lacked the switching mechanism reported here.

XRD analyses of the TiO2 films on FTO glass with or without AgNPs were performed using a Bruker D8 advance X-Ray Diffractometer (Bruker AXS GmbH, Karlsruhe, Germany) using a Cu Kα-radiation source set with an anode current of 40 mA and accelerating voltage of 40 kV with a scanning speed 0.015 degrees/second. The diffraction patterns were indexed by comparison with the Joint Committee on Powder Diffraction Standards (JCPDS) files number 21-1276 and 21-1272 for rutile and anatase respectively. The morphology and thickness of the TiO2 film was analysed by a ZEISS EVO MA 10 SEM equipped with an energy-dispersive spectrometer (EDS, Oxford Instruments, 129 eV resolution). The thin films were in some cases sputtered with gold, of 5 nm thickness, in order increase the conductivity of the samples prior the SEM imaging. TEM studies of the TiO2 nanoparticles, crystallites and Ag nanocrystals were carried out with a Philips CM20 electron microscope equipped with a Gatan GIF200 energy filter.

#### *2.6. Electrochemical Measurements*

Electrochemical and spectroelectrochemical experiments were performed using an Autolab PGStat 101 potentiostat. The spectroelectrochemical cell was a 6 mL, three-electrode teflon cell with quartz windows, employing a platinum mesh flag as the counter electrode, an Ag/AgCl/KClsat reference electrode, and the Ag-TiO2 film on FTO conducting glass as the working electrode. The electrochemical cell had an inlet and an outlet for passing gas into it. The electrolyte, an aqueous solution of 10 mM sodium phosphate (pH 7), was thoroughly de-aerated by bubbling with Argon prior to any electrochemical measurements and an Argon atmosphere was maintained throughout the measurements. For spectroelectrochemistry, the above cell was incorporated in the sample compartment of the Shimadzu UV-1800 spectrophotometer, and the absorption changes were monitored as a function of the applied potential. All potentials are reported against Ag/AgCl and all experiments were carried out at room temperature.

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

The surface morphology, structure and thickness of the TiO2 films prior and after deposition of AgNPs were analysed by SEM. The SEM images (Figure 1a) of TiO2 film showed disordered porosity and comprise a rigid, porous network of TiO2 nanoparticles of average size 30–40 nm. The film exhibits great homogeneity and even size distribution, while all nanoparticles are bonded together through the sintering process, creating a rich mesoporous surface. The thickness of the coated titania films (Figure 1b) was found to be around 6 μm by analysing the cross-section SEM images. These results confirm that the mesoporous film structure of TiO2 could provide an excellent surface for the AgNPs to diffuse throughout the porous structure. In addition, the porous film could provide many active sites for electrocatalytic reactions. Mesoporous layers are most suitable for immobilizing electroactive compounds, as the surface area available for sensitization and hence electrochemistry can be increased by over two orders of magnitude with respect to a flat electrode, while ensuring satisfactory access to the pores [37]. Agglomerates (showed with a red arrow in Figure 1c) are formed when AgNPs are deposited on top of the TiO2 films. The AgNPs agglomerates are shown as white dots on the SEM images at the sample surface and exhibit a large size distribution from 200 to 800 nm, with an average size of 600 nm. Back-scattered electrons (BSE) images have also provided evidence towards the different chemical content of the agglomerates with respect to that of the surface. The EDS spectrum of Ag-TiO2 (Figure 1d) shows the main peaks of Ti, O, Ag confirming the large amount of silver present. The extra addition of equal volumes of 0.5 M Na2S in the AgNO3:PVP mixture, which most probably would have passivated the surface of the AgNPs with a thin layer of Ag2S, was not detected in the EDS spectra, yet it is considered that a thin layer of Ag2S may have formed on the AgNPs surface. Finally, EDS elemental analysis led to the conclusion that the TiO2:Ag molar ratio at the surface was 8, while in the Supplementary Information (SI Figure S1) a surface EDS mapping can be observed that indicates the successful Ag coverage.

**Figure 1.** Scanning electron microscope (SEM) images for a bare TiO2 film on fluorine-doped indium tin oxide (FTO) substrate (**a**) top view and (**b**) cross section, (**c**) SEM image of Ag-TiO2 film and (**d**) energy-dispersive spectrometry (EDS) of Ag-TiO2 on FTO substrate.

TEM was employed on Ag-TiO2 films in order to further investigate the surface morphology and the size of the AgNPs. TiO2 nanoparticles (Figure 2a,b) exhibit the shape of platelets with external large dimension of 20–45 nm, in agreement with the specifications of the Dyesol and/or P25 products. AgNPs (Figure 2c,d) form different structures; in some cases plate-like, when they are deposited on TiO2 films with average diameter of 15 nm. High-resolution TEM (HR-TEM) imaging, Figure 2c, permitted easy differentiation of Ag nanocrystals (small dark areas) and TiO2 crystallites (large bright areas). Ag nanocrystals can be observed on the surface of the TiO2 particles as dark lines. It is presumed that even smaller Ag nanoparticles are in existence scattered throughout the porous TiO2 surface.

**Figure 2.** Transmission electron microscope (TEM) images (**a**,**b**) of TiO2 at low TEM resolution and high-resolution TEM (HR-TEM) images (**c**,**d**) of Ag-TiO2 composites. Images, a and c, on the left have been magnified at the selected dotted squares and placed to the right as b and d, respectively.

The color change of the Ag-TiO2 films is significant upon bias application. Figure 3 presents the thin films of TiO2 as prepared from Dyesol (Figure 3a) and from Degussa (Figure 3d). After immersing the TiO2 films in the AgNPs solution the film color changes (Figure 3b), since AgNPs were adsorbed on the TiO2 film. In order to perform the electrochemical measurements, negative bias was applied in the film and the color of the film became darker (Figure 3c,e).

**Figure 3.** Digital photos (**a**) mesoporous TiO2 film as fabricated from Dyesol and (**d**) from Degussa precursors. Ag-TiO2 films (**b**) before and (**c**,**e**) after bias application.

In order to investigate any changes in the crystal structure of the TiO2 films affected by the AgNPs, the XRD patterns (Figure 4) of FTO-conducting substrates, TiO2 and Ag-TiO2 film electrodes were measured. The XRD of TiO2 and AgNPs-TiO2 electrodes revealed similar characteristic peaks at 2θ: 25.28◦, 37.8◦ and 48.05◦ which correspond to the indices of anatase TiO2 (101), (004) and (200) hkl planes and are consistent with the reported values of the JCPDS file (21-1272). As expected, no characteristic peaks that correspond to the rutile TiO2 indices were observed since the Dyesol TiO2 paste is 100% anatase. The only difference was observed on the relative intensities of the peaks, which may be due to the fact that doping alters the crystallinity but not the crystal structure of the Ag-TiO2 films. In addition, two slightly intense peaks at 26.54◦ and 38◦ which correspond to the plane indices 110 and 200 of the FTO glass could be observed. The XRD pattern of the AgNPs-TiO2 film showed a new peak at 44.4◦ which can be assigned to the (200) plane of Ag. Furthermore, a closer look of the peak at 38◦ of the Ag-TiO2, which corresponds to the (200) FTO index, also showed a small shoulder at 38.1◦ which can be assigned to the Ag (111) plane. These two peaks agree with the JCPDs card of Ag as presented in Figure 4. No rutile phase or any other modification is observed for the Ag-TiO2 film depending on AgNPs incorporation.

**Figure 4.** X-ray diffraction (XRD) patterns of a TiO2 film with or without silver nanoparticles (AgNPs) deposited on its surface, FTO glass and Joint Committee on Powder Diffraction Standards (JCPDS) cards of anatase and rutile TiO2. Inset: magnification of the peak at 38o corresponding to the fcc Ag (111) plane.

The TiO2 films deposited on conducting FTO glass slides had been soaked in solutions of AgNPs and rinsed with NaH2¬PO4 buffer to remove any loosely bound nanoparticles. The mp TiO2 films combine transparency in the visible region of the electromagnetic spectrum with a high surface area accessible to molecules from a surrounding solution (e.g., AgNPs). In many cases in the past, this allowed the adsorbed molecules (dyes, proteins, electrochromic species and surfactants) to achieve the densities required for informative electronic absorption spectroscopy, whether that was for the development of solar cells, electrochemical biosensors, electrochromic devices or catalytic applications [3,4,9,27,38,39]. Therefore, the transparency of the electrodes can be useful as it could allow the AgNPs adsorption process and their electrochromic properties to be monitored by UV–Vis absorption spectroscopy, which is a technique also used by other authors for the structural characterization of the AgNPs in a dielectric matrix [27,38–41].

Figure 5 shows the optical absorption spectra of the AgNPs solution used as dopants for the surface of the TiO2 films. A relatively broad absorption band is located at 452 nm properly corresponding to the size and form of the AgNPs; this peak is due to the SPR of the AgNPs and is within the spectrum range (400–450 nm) reported for them depending on their shape and size [27,30,38,39,41]. Adsorption of AgNPs on TiO2 films results in light coloration (see Figure 3) of the films, indicating that a large amount of AgNPs has been immobilized into the mesoporous TiO2 film. Also shown is the absorption spectrum of a bare TiO2 film which is transparent and colourless in the visible region, showing a characteristic absorption increase below 400 nm due to the onset of TiO2 band gap excitation. Therefore, the optical transparency of the TiO2 allows the adsorption process of the AgNPs to be monitored by UV–Vis absorption [30,42]. The electronic absorption spectra (Figure 5) for the nanocomposite Ag-TiO2 film showed features typical of AgNPs superimposed on a background arising from scattering by the TiO2 layer. The resulting spectrum of AgNPs on the TiO2 film showed a characteristic absorption band at 413 nm (a clear blue shift in comparison with the absorption spectra of the AgNPs in solution). The Ag-TiO2 film exhibits a much narrower, blue shift and more defined optical spectrum which means that the active metallic cores of the deposited AgNPs are smaller than those in the solution and probably spherical [27,38,41,43]. The shift is related to the size of the AgNPs and to the interaction

between TiO2 and Ag, as well as to the fact that the majority of the AgNPs that have filled the pores are smaller than the pores themselves. No band is observed of organic residuals, which were used to bind the TiO2 nanoparticles, remaining in the film due to the titania film sintering at 450 ◦C. If the observed plasmonic bands were broad between 510 and 590 nm, that would have implied that the nanoparticles are large and/or of non spherical nature [44]. Increasing the AgNO3 solution concentration, where the AgNPs were created, had as a result the increase of the absorbance of the Ag-TiO2 film, as was also reported by other researchers in the past [30]. There is also the possibility that some of the non-reduced silver ions (Ag<sup>+</sup>) could be adsorbed on the hydroxylated TiO2 surface according to the following reaction, as suggested by other authors also in the past: Ag<sup>+</sup> + TiO2 <sup>−</sup> <sup>→</sup> Ti–O–Ag <sup>+</sup> <sup>H</sup><sup>+</sup> [45].

The plasmon resonance in the Ag-TiO2 thin films strongly depends on the crystalline phase and dielectric constant of the TiO2 matrix [46]. In theory, the SPR peak wavelength increases with increasing dielectric constant of the matrix and depends on the refractive index.

**Figure 5.** Optical absorption spectra in the visible range (380–700 nm) of (a) AgNPs solution, (b) TiO2 film and (c) Ag-TiO2 film.

It should be mentioned that the form of the absorption profile does not only signify that the 415 nm absorption peak of the Ag-TiO2 films is characteristic of spherical AgNPs. It also carries information due to the band absence between 520 and 540 nm which is usually related to bigger AgNPs or AgNPs dimmers. Similarly, no band is observed around 620 nm which is usually due to non-spherical Ag nanoparticles or to larger particles or to manifestation of higher-order plasmon modes called quadrupolar modes. No band is also observed around 670 nm which suggests a longitudinal plasmon band of Ag nanorods. Any broadness of the observed absorption bands could refer to different morphologies of the deposited nanoparticles, to broad particle size distribution, and to agglomeration processes, while all the above discussion is in agreement with the TEM analysis.

The optical band gap of the Ag-TiO2 films is expected to be smaller to that of pure (blank) TiO2 films due to the effect of the AgNPs, as decrease of Eg with silver addition has also been reported by other authors [27,38,39,41]. Also, the Ag<sup>+</sup> ions probably exist on the surface of the anatase TiO2 films by forming Ag–O–Ti bonds [47], which may introduce trap states affecting the energy band gap. Finally, it is possible that the potential applied can oxidize or reduce the Ag nanoparticles to silver oxide and reverse this as suggested by Kuzma et al [48].

Figure 6 presents the spectroelectrochemical spectrum for a Ag-TiO2 film electrode after remaining for 2 min at each negative applied potential (−0.1 to −1.1 V) applied. The plasmon (or Soret) band at 413 nm up until the application of −0.4 V remains constant, but upon the application of −0.5 V, it starts to increase in intensity without though changing shape. The application of −0.6 V to the film causes a big increase of the peak which at the same time becomes thinner and sharper. Afterwards, and by gradually applying more negative biases up to −1.1 V, the peak at 413 nm continues to rise and at the same time becomes sharper. However, several minutes after the end of the spectroelectrochemical measurements and the application of no voltage, the absorbance of the film shows that it has not returned to its initial state, but rather a shift of the main peak from 413 nm to 440 nm has been observed.

**Figure 6.** Ultraviolet–visible (UV–Vis) spectral changes of a Ag-TiO2 film electrode upon the application of increasingly negative potentials (0 to −1.1 V vs. Ag/AgCl).

Following this route, increasing negative biases (−0.2 to −0.9) were applied again to the doped film and its absorption spectra were recorded. Figure 7 shows that upon the application of up to −0.4 V, a small increase in the size of the peak at 440 nm was monitored and a slight blue shift. However, upon the application of higher negative biases (−0.8 or −0.9 V) the peak shifts back to 413 nm. Figure 7 also shows that upon the application of a positive bias (0 to 0.6 V) the peak shifts again to 440 nm. The switch from 413 nm to 440 nm, depending on the bias that was applied to the film, was repeated several times and always with success.

**Figure 7.** UV–Vis spectral changes of a Ag-TiO2 film electrode under the application of negative or positive potentials showing the shift of the Soret peak from 413 nm to 440 nm and vice versa.

In order to gain an insight into the kinetic mechanism, the optical absorbance (OA) or optical density (OD) were measured for a range of negative and positive voltages. Figure 8a,b show the curves regarding the kinetics for the absorption change at 413 (a) and 440 nm (b) respectively. The absorbance was initially monitored during the application of −0.8 V for 600 s and immediately afterwards upon the application of 0.15 V for another 600 s. Figure 8a demonstrates that the continuous application of a sufficient cathodic current (−0.8 V) causes the fast increase of the absorbance at 413 nm of the Ag-TiO2 film in 4.5 s (τ1/<sup>2</sup> = 2.6 s). By stepping back the potential from −0.8 V to +0.15 V, the absorbance at 413 nm starts dropping quite fast at the first 20 s and slower afterwards until it reaches the OD value it exhibited before the application of the negative bias. Similar results were obtained in Figure 8b for the kinetics for the absorption change at 440 nm. The continuous application of the −0.8 V caused the fast increase of OD at 440 nm of the Ag-TiO2 in 6 sec (τ1/<sup>2</sup> = 3.8 s). The application of the +0.15 V caused the gradual drop of the OD at 440 nm and reached completion after 50 s.

**Figure 8.** Kinetics for the absorption change at 413 nm (**a**) and 440 nm (**b**) of a 6 μm thick Ag-TiO2 film by the application of 0.15 V and −0.8 V.

The spectroelectrochemical studies reported in Figures 6 and 7 were further supported by cyclic voltammetry (CV). The CVs of mp TiO2 films with or without the adsorption of AgNPs on their surface, in aqueous 10 mM NaH2PO4 electrolyte of pH 7, at a scan rate of 0.1 V/s, are presented in Figure 9. Upon scanning the potential of a bare semiconductive TiO2 film cathodically in a neutral pH aqueous medium (NaH2PO4, pH 7), a transition from insulating to conductive behaviour is observed. The characteristic charging/discharging currents were assigned to electron injection and storage into sub-band gap/conduction band states of the TiO2 film, until the metal oxide becomes fully degenerated once the applied potential reaches the conduction band potential [9,33,34,49]. The current shows a plateau at potentials where the film behaves as an insulator (positive biases and up to around −0.3 V). At that range the electrical response is dominated by the Helmholtz capacity of the uncovered FTO glass/electrolyte solution interface at the bottom of the TiO2 film [50,51]. At more negative potentials (−0.3 V and higher), the cathodic current displays an exponentially rising behaviour (at −0.8 V) that is considered by many authors as the reduction of superficial Ti ions [52,53]. This rising behaviour is then transformed to a peak (at −0.75 V), when the direction of voltammetry is reversed to anodic. This is considered as re-oxidation of reduced Ti ions and is a slow and irreversible process [53]. These potentials are more negative than the conduction band potential and the TiO2 film behaves as a conductive metallic electrode. As the scan rate becomes slower (Figure 10) the height of the anodic peak progressively diminishes until, at very slow scan rates (0.01 V/s), it disappears. According to many reports [1,7,53,54] the origin of this behaviour is due to the charging/discharging of electrons in the film and a charge transfer mechanism. However, no cathodic or anodic peaks due to a redox reaction are observed. The CV integrates to approximately 0, indicating negligible Faradaic currents.

In addition, Figure 9 shows the CV of a TiO2 film electrode after the adsorption of AgNPs on its surface. The Ag-TiO2 film exhibited two pairs of redox peaks that were absent from the CV of the bare TiO2 film. For the first pair, a small cathodic peak is observed at +0.07 V and a small anodic peak at +0.27 V. For the second pair the cathodic peak is much bigger and appears at −0.65 V and the anodic peak which is much smaller and broader appears around −0.53 V. All these peaks are attributed to two different reduction/oxidation states of Ag on the TiO2 films. The redox reactions are associated with Ag → Ag2O and Ag2O → Ag (oxidation and reduction peaks, respectively) [55,56]. The pair of peaks at the lower biases correspond to the reduction and oxidation of the AgNPs that have adsorbed inside the mesoporous network of the TiO2 film and the other pair of smaller peaks most probably correspond either to small deposited AgNPs or to AgNPs adsorbed only on the outer surface of the TiO2 film [46]. Also both re-oxidation peaks are broader and smaller to the reduction peaks due to the large band gap of the TiO2 film. This observation most probably requires further experiments in order to draw any final conclusions. However, this redox behaviour (two sets of redox peaks, a two phase reduction) of the AgNPs on the TiO2 films has also been monitored on immobilization studies of a heme based redox dyes, such as iron(III) 5,10,15,20-tetrakis(1-methyl-4-pyridyl) porphyrin, FeTMPyP, on the same films [33].

**Figure 9.** Cyclic voltammograms (CVs) in 10 mM NaH2PO4 buffer, pH 7 for Ag-TiO2 and blank TiO2 films. The scan rate was 0.1 V/s.

The effect of the scan rate on the voltammetric behaviour of the Ag-TiO2 film was also investigated. Slower scan rates were applied to the Ag-TiO2 electrode in order to try to obtain a reversible peak-shaped CV. Figure 10 illustrates that even at slow scan rates no simple reversible behaviour for the Ag-TiO2 film was observed, consistent with the currents being limited by the low TiO2 conductivity at moderate potentials. However, the cathodic peak potential at −0.65 V shifted negatively with the increase of the scan rate. In addition, there is a good linear relationship between the cathodic peak current and the scan rate in the range (0.01 to 0.1 V/s), indicative of a surface-confined electrochemical process. This has also been observed in other systems where instead of AgNPs, surfactants or proteins were adsorbed on the TiO2 films [39].

Also, it was found that the cathodic and anodic peak currents corresponding to the adsorbed AgNPs vary linearly with scan rates in the range of 10 to 100 mV·s–1. The CVs exhibit a current increase as the scan rates are steadily increased, significantly larger when the 0.1 V/s scan rate results are compared to those of the 0.01 V/s. This is clear indication that the adsorbed AgNPs underwent a surface-controlled process.

**Figure 10.** Cyclic voltammograms (CVs) in 10 mM NaH2PO4 buffer pH 7, for a Ag-TiO2 film at different scan rates.

#### **4. Conclusions**

All the above experimental data clearly suggest that the electrophotochromism of the Ag-TiO2 composite is due to oxidation/reduction of the Ag nanoparticles, which forms a thin layer of Ag2O on the metallic core, forming core/shell nanoparticles; the core is metallic while the outer shell is semiconducting, where the peaks agree well with those found by Kuzma et al. [46]. In fact, it is suggested that the oxide thickness is about 1.5 nm when the composite absorbs at the 440 nm peak and at most 0.5 nm while the composite exhibits a peak at 413 nm. Furthermore, it is possible that the particles may be charged upon the application of negative potential on the Ag-TiO2 electrode, which would lead their separation due to their charging; such an effect would not be possible at positive electrode potentials since silver has a lower conduction band than that of the FTO electrode and than that of TiO2.

#### **Supplementary Materials:** The following are available online at http://www.mdpi.com/2079-6412/10/2/130/s1.

**Author Contributions:** E.T. and I.K. conceived and designed the experiments; S.K., J.P., E.T. and I.K. performed the experiments; E.T., I.K., S.K. and J.P. analyzed the data; E.T. and I.K. contributed reagents/materials/analysis tools; I.K., S.K., E.T. and J.P. wrote the paper. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding from funding agencies in the public, commercial or not-for-profit sectors.

**Acknowledgments:** We would like to thank Nikolaos Boukos from NCSR "Demokritos" for the TEM images.

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

#### **References**


© 2020 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **A Rapid Synthesis of Mesoporous Mn2O3 Nanoparticles for Supercapacitor Applications**

**You-Hyun Son 1, Phuong T. M. Bui 1, Ha-Ryeon Lee 1, Mohammad Shaheer Akhtar 1,2,\*, Deb Kumar Shah <sup>1</sup> and O-Bong Yang 1,2,\***


Received: 20 August 2019; Accepted: 24 September 2019; Published: 30 September 2019

**Abstract:** Mn2O3 nanomaterials have been recently composing a variety of electrochemical systems like fuel cells, supercapacitors, etc., due to their high specific capacitance, low cost, abundance and environmentally benign nature. In this work, mesoporous Mn2O3 nanoparticles (NPs) were synthesized by manganese acetate, citric acid and sodium hydroxide through a hydrothermal process at 150 ◦C for 3 h. The synthesized mesoporous Mn2O3 NPs were thoroughly characterized in terms of their morphology, surfaces, as well as their crystalline, electrochemical and electrochemical properties. For supercapacitor applications, the synthesized mesoporous Mn2O3 NP-based electrode accomplished an excellent specific capacitance (*Csp*) of 460 F·g−<sup>1</sup> at 10 mV·s−<sup>1</sup> with a good electrocatalytic activity by observing good electrochemical properties in a 6 M KOH electrolyte. The excellent *Csp* might be explained by the improvement of the surface area, porous surface and uniformity, which might favor the generation of large active sites and a fast ionic transport over the good electrocatalytic surface of the Mn2O3 electrode. The fabricated supercapacitors exhibited a good cycling stability after 5000 cycles by maintaining ~83% of *Csp*.

**Keywords:** Mn2O3; mesoporous materials; electrochemical characterizations; electrode; supercapacitors

#### **1. Introduction**

A popular electrochemical energy-storage system, the supercapacitor has been a well-explored device as a heartening energy storage because of its high-power density, marvelous cycling and fast charge-discharge mechanisms [1–4]. Supercapacitors, on the basis of their charge-storage process, are classified into two types: i) double layer electrochemical capacitors (EDLCs), built from the electrode and electrolyte interface for the ions adsorption-desorption, and ii) the electrode materials-governed faradaic reaction-based pseudocapacitors [5–7]. In comparison with the EDLCs, the pseudocapacitors have exhibited a fast and desirable reversible redox reaction that promotes an excellent charging and discharging process, resulting in the enhancement of the charge storage capacities. In addition, pseudocapacitors display a highly competent storage device for rechargeable batteries owing to their fast energy harvesting, high energy and high-power delivery [8–10]. In supercapacitors, the active electrodes are normally prepared with carbon-based materials, conducting polymers, and a variety of transition metal oxide materials [11–14]. In recent years, transition metal oxides, such as RuO2, MnO2, Mn3O4, NiO, Nb2O5, V2O5, CoO*x*, MoO3, and TiO2, have been frequently applied in preparing an effective electrode for supercapacitors due to their abundance in nature, inexpensive nature and extraordinary redox activity [13–21].

Apart from other transition metal oxides, the manganese-based oxides (MnO*x*) and their derivatives, like MnO2, Mn2O3, and Mn3O4, are popularly used as electrode materials in supercapacitors because they have a non-toxic nature, good structural flexibility and excellent chemical and physical stability in various electrolytes [22]. In particular, Mn2O3 materials as anode materials in lithium ion batteries have shown a high capacity and demonstrate a theoretically high specific capacity of ~1018 mAh/g, while also exhibiting a high specific capacitance [23]. Until now, the Mn2O3 materials-based electrodes in supercapacitors have been less explored and could be expected to have a high capacity and storage properties as they show an excellent environment compatibility and a resistance in acidic/alkaline electrolytes. Numerous efforts have been made to improve the performance of Mn2O3-based electrodes by adopting various modifications, such as chemical modifications [24], the incorporation of high surface-area conductive materials [25,26] and nanostructure fabrication [24,27]. Therefore, the synthesis of Mn2O3 materials with a controlled size, morphology and density through a cost-effective, simple and environment friendly method could be potentially sound [28,29]. It is expected that mesoporous Mn2O3 nanomaterials might show a surface that is fruitful for fast ion transportation in electrochemical supercapacitors [30].

In this work, a rapid and low temperature hydrothermal process was used to synthesize well-crystalline mesoporous Mn2O3 materials that were successfully applied as electro-active materials for a pseudosupercapacitor. The prepared mesoporous Mn2O3 materials-based electrode exhibits a high specific capacitance of 460 F·g−<sup>1</sup> at a scan rate of 10 mV·s−1, with a good cycling stability after 5000 cycles.

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

In a typical synthesis, a mixture of 1 g of manganese acetate (Mn(CH3CO2)2, Sigma-Aldrich, Saint Louis, MO, USA) and 0.5 g citric acid (Samchun Chemicals, Seoul, Korea) was dissolved in 100 mL of deionized (DI) water. Using an aqueous 5 M NaOH solution, the pH of the reaction mixture was adjusted to pH ~10 as the solution color changed from bright brown to dark brown. The hydrothermal process was carried out by transferring the reaction mixture into a Teflon beaker that was tightly sealed. Finally, a temperature of 150 ◦C was maintained for 3 h. After cooling down the autoclave, the obtained precipitates were collected by filtration, washed several times with DI water and with ethanol. The collected precipitates were dried in an oven overnight at 80 ◦C, and finally the obtained black powder was calcined at 500 ◦C in 1 h to remove other impurities.

For the electrochemical supercapacitor application, the synthesized mesoporous Mn2O3 electrode was prepared by mixing 85 wt.% Mn2O3 powders, 10 wt.% carbon black (Super P, Venatech, Seoul, South Korea), 3 wt.% carboxyl methyl cellulose (CMC, Sigma-Aldrich) and 2 wt.% polytetrafluoroethylene (PTFE, TCI chemical, Tokyo, Japan) in DI water to obtain a paste that was spread over the Ni foam using a glass rod via the rolling method. Afterward, the mesoporous Mn2O3-coated Ni foam electrodes were dried in the oven at 80 ◦C for 20 min to remove the solvent. For the electrochemical measurement, a three-electrode system, comprised of Mn2O3-coated Ni foam as the working electrode, Pt wire as the counter electrode and Ag/AgCl as the reference electrode, was used, and a cyclicvoltametric measurement (VersaSTAT4, AMETEK, Inc., Berwyn, PA, USA) was performed in an aqueous 6 M KOH electrolyte. All cyclicvoltametry (CV) measurements were observed at different scan rates ranging from 10 to 500 mV·s−<sup>1</sup> in the voltage range of 0 to 1.0 V. A potentiostat/galvanostat (VersaSTAT4, AMETEK, Inc.) was used to analyze the electrochemical impedance spectroscopy (EIS) of the fabricated supercapacitor, based on the mesoporous Mn2O3 electrode, with a frequency ranging from 0.1 Hz to 1 MHz. For the calculation of *Csp*, the mass loading of the mesoporous Mn2O3 on the electrode was ~0.001 g.

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

The morphological features of the synthesized mesoporous Mn2O3 materials were analyzed by field emission scanning electron microscopy (FESEM, Hitachi S-4800, Tokyo, Japan) and transmission

electron microscopy (TEM, JEM-ARM200F, JEOL, Peabody, MA, USA) observations. Figure 1a,b shows the FESEM images of the synthesized mesoporous Mn2O3 materials at low and high magnifications. The spherical small particles, which are highly uniform, are visible in Figure 1a. From FESEM observations, it is difficult to identify the porosity of the synthesized materials. At a high magnification mode (Figure 1b), the obtained particles possess highly porous structures with average sizes of 10–30 nm. A similar observation for synthesized mesoporous Mn2O3 materials was detected in the TEM and high-resolution transmission electron microscopy (HRTEM) analyses, as shown in Figure 1c,d. As seen in Figure 1c, the synthesized Mn2O3 materials show a similar spherical shape, with a few semi-spherical particles having average particle sizes of ~10–30 nm. A close look at Figure 1c shows that the porosity of the synthesized Mn2O3 materials might be defined by the presence of visible voids over the particle surfaces. As reported earlier [31], these voids in materials may be a detrimental factor of the porosity of the materials. The HRTEM image is shown in Figure 1d; it expresses clear lattice fringes from when the measurement was focused on one particle from Figure 1c. From the HRTEM image, the interplanar distance between two lattice fringes is estimated to be ~0.27 nm, which is well-indexed to the (222) plane of Mn2O3 [32].

**Figure 1.** The (**a**,**b**) FESEM, (**c**) TEM and (**d**) HRTEM images of the synthesized mesoporous Mn2O3 materials.

The crystalline behavior and crystal planes of the synthesized mesoporous Mn2O3 materials were determined via a wide-angle X-ray diffraction (XRD, PANalytical, Malvern, United Kingdom) measurement, as shown in Figure 2. The well-defined diffraction peaks at 18.9◦, 23.1◦, 33.1◦, 38.2◦, 45.2◦, 47.3◦, 49.5◦, 55.1◦ and 65.8◦ are associated to (200), (211), (222), (400), (332), (422), (431), (440) and (622) planes, respectively. All of the obtained diffraction peaks are assigned perfectly to the Bixbyite crystal phase α-Mn2O3, with JCPDS no. 41-1442 and space group Ia3, lattice constants *a* = *b* = *c* = 9.4091 Å, α = β = γ = 90◦. To estimate the crystallite sizes (*CS*) of the synthesized mesoporous Mn2O3 materials, the Debey–Scherrer equation was used [33]:

$$\text{CS} = \frac{0.95 \times \lambda}{\beta \cos \left(\theta\right)}\tag{1}$$

where β is the breadth of the observed diffraction line at its half-intensity maximum, *K* is the so-called shape factor, which usually takes a value of about 0.9, and λ is the wavelength of the X-ray source used in the XRD. By taking the maximum diffraction peak of (222), the crystallite size of the mesoporous Mn2O3 materials is found to be 28 nm, which is very close to the FESEM and TEM results.

**Figure 2.** The XRD patterns of the synthesized mesoporous Mn2O3 materials.

Figure 3 shows the infrared (IR, Nicolet, IR300, Thermo Fisher Scientific, Waltham, MA, USA) and Raman spectroscopic studies (Raman microscope, Renishaw, UK) that define the structural properties of the synthesized mesoporous Mn2O3 materials. As seen in Figure 3a, two sharp IR bands are observed at 610 and at 520 cm<sup>−</sup>1, assigning the stretching vibrations of Mn–O units and the asymmetric Mn–O–Mn stretching vibration, respectively [34]. Other IR bands at 1640 and 3343 cm−<sup>1</sup> are detected, related to –OH and the water species from atmospheric moisture. It is believed that the observation of the IR bands at 520 and 610 cm−<sup>1</sup> clearly reveals the formation of Mn2O3 without other impurities. Figure 3b depicts the Raman scattering spectroscopy of the synthesized mesoporous Mn2O3 materials. The mesoporous Mn2O3 NPs present a strong Raman band at 651 cm<sup>−</sup>1, including with two weak Raman bands at 268 and at 175.0 cm−1. The strong Raman band at 651 cm−<sup>1</sup> represents the characteristic of the Mn2O3 along with the space group Ia3 structure [35], suggesting the typical symmetric stretching Mn–O–Mn bridge in Mn2O3. The main Raman band is well-matched with the reported literature of Mn2O3 [35]. Additionally, two weak Raman bands at ~268 and ~175 are assigned to the out-of-plane bending modes of Mn2O3 and the asymmetric stretching of the bridge oxygen species (Mn–O–Mn) [36], respectively.

To explain the thermal and structural properties, a thermal gravimetric analysis (TGA, Thermal analyzer, TA Instrument Ltd., New Castle, DE, USA) was performed for the synthesized mesoporous Mn2O3 materials, as displayed in Figure 4a,b. As seen in Figure 4b, four decomposition temperatures were visibly identified in the range of 25 to 800 ◦C. After 500 ◦C, the subsequent weight loss of ~1.5% that started from 500 to 600 ◦C is ascribed to the thermal decomposition of Mn2O3 to MnO, a decomposition that is usual in metal oxides. In the beginning, the synthesized mesoporous Mn2O3 materials were recorded as having a very small weight loss of ~0.5% up to 500 ◦C, which usually occurs via the removal of water/moisture from the sample. This suggests that the synthesized mesoporous Mn2O3 NPs exhibit a remarkably good stability with a high crystalline nature of Mn2O3.

**Figure 3.** The (**a**) Fourier-transform infrared spectroscopy (FTIR) and (**b**) Raman spectrum of the synthesized mesoporous Mn2O3 materials.

**Figure 4.** (**a**,**b**) TGA plot of the synthesized mesoporous Mn2O3 materials.

The synthesized mesoporous Mn2O3 materials were further characterized in terms of their composition and the oxidation states of the elements using an X-ray photoelectron spectroscopy (XPS, AXISNOVA CJ109, Kratos Inc., Manchester, UK) analysis. Figure 5a shows the survey XPS spectrum of the synthesized mesoporous Mn2O3 materials, revealing Mn 2*p*, Mn 3*s* and O 1*s* with weak C 1*s* peaks. The high-resolution Mn 2*p* XPS spectra of the synthesized mesoporous Mn2O3 materials is shown in Figure 5b and demonstrates doublet binding energies at 641.0 (Mn 2*p*3/2) and 652.8 eV (Mn 2*p*1/2). It is noted that the doublet Mn 2*p*, with an estimated spin–orbit splitting value of 11.8 eV, is nearly the same as the values reported for Mn2O3 [37]. Additionally, Figure 5c displays the high-resolution Mn 3*s* with characteristic doublet binding energies at 88.8 and 83.5 eV for Mn2O3. From the Mn 3*s* XPS, the peak separation for the doublet binding energies are ~5.3 eV, which is very close to the peak separation of Mn 3*s* in Mn2O3 [38], which again implies the formation of Mn2O3. Figure 5d depicts the high-resolution O 1*s* XPS spectrum. The two binding energies at 528.9 and 530.7 eV are normally related to oxygen O2<sup>−</sup> in the lattice of Mn–O–Mn, indicating the formation of Mn2O3. Therefore, the XPS analysis implied the formation of a pure Mn2O3 form without any other oxide impurities [39].

The surface and porous behavior of the synthesized mesoporous Mn2O3 materials have been elucidated by analyzing the nitrogen (N2) adsorption-desorption isotherms and BET surface analyzer, as shown Figure 6. The N2 adsorption-desorption isotherms (Figure 6a) present a regular type IV isotherm with a small hysteresis in the relative pressure (*p*/*p*0) range of 0.75–0.9, which clearly imitated the mesoporous characteristic of materials. Figure 6b shows how the pore size distribution (*d*) of the mesoporous Mn2O3 materials was recorded in the range of 10–30 nm, which is the case for mesoporous materials. Using a Brunauer-Emmett-Teller (BET) surface analysis, the specific surface area (*s*) for the synthesized mesoporous Mn2O3 materials was determined to be 76.9 m2·g<sup>−</sup>1. Thus, the synthesized Mn2O3 materials are mesoporous in nature which could provide a larger surface area for the high ion diffusion of ions in the electrolyte for high-performance supercapacitors.

**Figure 5.** (**a**) Survey, (**b**) high resolution Mn 2*p*, (**c**) Mn 3*s*, and (**d**)O1*s* XPS of the synthesized mesoporous Mn2O3 materials.

**Figure 6.** The (**a**) N2 adsorption-desorption isotherm and (**b**) pore size distribution plot of the synthesized mesoporous Mn2O3 materials.

The synthesized mesoporous Mn2O3 materials were utilized as electroactive materials to evaluate the supercapacitor properties. The parameters for the supercapacitor with the synthesized mesoporous Mn2O3 electrode were determined by measuring the cyclicvoltametry (CV) at different scan rates in a 6 M KOH electrolyte. Figure 7a shows a series of CV curves of the synthesized mesoporous Mn2O3 electrode at different scan rates in a 6 M KOH electrolyte. Generally, in an electrochemical reaction, the conversion from Mn3<sup>+</sup> to Mn4<sup>+</sup> in an Mn2O3 electrode occurs via an oxidation reaction. This oxidation reaction normally accelerates the reaction kinetics of OH<sup>−</sup> over the Mn2O3 cubic lattice through the chemisorbed and/or intercalated. The charge/discharge process can be explained by the following electrochemical reaction, which involves the chemisorption/intercalation of HO− over the Mn2O3 surfaces:

$$\text{Mn}\_2\text{O}\_3 + \text{HO}^- \xrightarrow[\text{Discharge}]{\text{C}^{\text{Charge}}} \text{Mn}^{4+} \text{ [Mn}^{3+}] \text{OH}^- \text{O}\_3 + \text{ }\overline{\text{e}}$$

The prepared Mn2O3 electrode depicts the oxidation-reduction pair peaks with different curve shapes of the charge–discharge curves in relation to the variation in the scan rates. The observations principally suggest the origin of the faradaic pseudo-capacitance in an alkaline electrolyte over the surface of an Mn2O3 electrode. Moreover, the well-defined oxidation reduction pair within 0−1 V can also be ascribed to a faradaic redox reaction [40,41]. The plot of the capacitance versus the scan rate for a fabricated pseudosupercapacitor based on a mesoporous Mn2O3 electrode is displayed in Figure 7b. In general, the specific capacitance (*Csp*) is estimated from the CV curves using the following equation [42]:

$$\mathbb{C}\_{\rm sp} = \frac{1}{m\nu(V\_2 - V\_1)} \int\_{V\_1}^{V\_2} I(V)dV \tag{2}$$

where *Csp* is the specific capacitance (F·g−1), *I* is the current (A), *v* is the scan rate, *m* is the mass of the active material (g) and Δ*V* is the potential range (V). As seen in Figure 7b, the fabricated pseudosupercapacitor based on the mesoporous Mn2O3 electrode exhibits a high *Csp* of ~460 F·g−<sup>1</sup> at a scan rate of 10 mV·s<sup>−</sup>1. This might be explained by the fact that there is an improvement of the surface area, the porous surface and uniformity, which might favor the generation of large active sites [43,44] and fast ionic transport over the surface of the Mn2O3 electrode. Specifically, the porous structure could provide a large and accessible surface area to ion adsorption, improve the accessibility of cations and shorten the ion diffusion path. The stability of the mesoporous Mn2O3 electrode is explored by observing the multicycle CV measurements in a 6 M KOH electrolyte. From Figure 7c, the slight shift in the oxidation peaks expresses the good stability of the mesoporous Mn2O3 electrode in the alkaline electrolyte. It is also seen that the anodic current in the Mn2O3 electrode is positively shifted with the increase, while the cathodic current is negatively shifted due to the increment in the electrical polarization and the fast and irreversible reactions as the scan rates increase. Here, the fast, irreversible reactions after 25 cycles might result from the accumulation of Mn4<sup>+</sup> ions in the electrochemical system. The stability of the fabricated pseudosupercapacitor based on a mesoporous Mn2O3 electrode is shown in Figure 7d by measuring the capacitance after 5000 cycles. From Figure 7d, after 5000 cycles, the electrochemical system shows a reasonably high stability by retaining 83% of the initial capacity. Furthermore, the inset of Figure 7d presents the FESEM image of the mesoporous Mn2O3 electrode after 5000 cycles. The morphology of the Mn2O3 materials is not altered, except for an aggregation of small particles after cycling, suggesting the stability of the Mn2O3 electrode in the alkaline electrolyte. Hence, the reproducibility of a pseudosupercapacitor based on a mesoporous Mn2O3 electrode in a KOH electrolyte is remarkable and implies the low dissolution of the electro-active materials in a strong alkaline electrolyte after an electrochemical process.

The electrochemical Impedance Spectroscopy (EIS) for the fabricated pseudosupercapacitor based on the synthesized mesoporous Mn2O3 electrode was conducted to understand the electrochemical and electrical properties. Figure 8 shows the EIS plot, which was measured in the frequency range of 0.1–10<sup>5</sup> Hz. In the illustration of the equivalent circuit, the starting point in the Nyquist plot at a high-frequency region and the starting point at a low-frequency region represent the series resistance (*Rs*) and the charge transfer resistance (*Rct*) of the electrode/electrolyte interface, respectively. *Cdl*, *W* and *Cpseudo* explain the double layer capacitance that arose by the parallel connection to *Rct*, the Warburg diffusion element and the Faradaic capacitance generated by the contribution of the Mn2O3 electrode. As shown in Figure 8, the pseudosupercapacitor based on the synthesized mesoporous Mn2O3 electrode features a large phase angle near the low-frequency region, indicating the faradic capacitance behavior of the electrode. Importantly, in the EIS plot, the straight line in the low-frequency region, and the absence of any small semicircle in the high-frequency region, are indicative of the good capacitive nature of the present electrochemical system based on a mesoporous Mn2O3 electrode. Likewise, the observed straight line at a low frequency might reduce the diffusion length and accelerate the ions transportation on the mesoporous surface of the Mn2O3 electrode, as evidenced by the CV results.

**Figure 7.** The (**a**) CV curves and (**b**) specific capacitance of the synthesized mesoporous Mn2O3 electrode at different scan rates ranging from 10 to 500 mV·s<sup>−</sup>1, (**c**) multicycles CV curves and (**d**) variation in the specific capacitance of the synthesized mesoporous Mn2O3 electrode after 5000 cycles. The inset shows the FESEM image of the synthesized mesoporous Mn2O3 electrode after 5000 cycles.

**Figure 8.** The EIS plot of the synthesized mesoporous Mn2O3 electrode at the Zim versus the Zre mode. The inset shows its corresponding equivalent circuit.

#### **4. Conclusions**

In summary, a facial hydrothermal process was adopted to synthesize well-crystalline mesoporous Mn2O3 materials for the fabrication of a pseudocapacitor. The crystalline and structural characterizations confirmed the formation of Mn2O3 materials without displaying any other oxide forms. The surface properties were analyzed, showing the mesoporous nature of Mn2O3 materials and an estimated high surface area of 76.9 m2·g−<sup>1</sup> with a good pore distribution. The fabricated pseudo-capacitor based on a Mn2O3 mesoporous particle electrode shows a reasonably high specific capacitance of ~460 F·g−<sup>1</sup> at 10 mV·s−<sup>1</sup> in a 6 M KOH aqueous solution. The enhancement in the capacitive properties might be attributed to the high surface area, porous surface and uniformity of the unique mesoporous particles morphology, resulting in the large generation of active sites and a fast ionic transport over the surface of the Mn2O3 electrode.

**Author Contributions:** Conceptualization, Y.-H.S., P.T.M.B., M.S.A., H.-R.L., D.K.S. and O-B.Y.; methodology, Y.-H.S., P.T.M.B., M.S.A., H.-R.L. and D.K.S.; software, Y.-H.S. and P.T.M.B.; validation, P.T.M.B., M.S.A., H.-R.L. and D.K.S.; formal analysis, Y.-H.S. and P.T.M.B.; investigation, Y.-H.S. and P.T.M.B.; resources, Y.-H.S. and P.T.M.B.; data curation, Y.-H.S., P.T.M.B., H.-R.L. and D.K.S.; writing—original draft preparation, Y.-H.S., P.T.M.B., and M.S.A.; writing—review and editing, Y.-H.S., P.T.M.B., M.S.A., and O-B.Y.; supervision, M.S.A., and O-B.Y.

**Funding:** This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MSIT) (NRF-2019R1F1A1063999).

**Acknowledgments:** This paper was also acknowledged the support of research funds from Chonbuk National University in 2018.

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

#### **References**


© 2019 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 (http://creativecommons.org/licenses/by/4.0/).

*Article*
