Plasmonic Rutile TiO₂/Ag Nanocomposites Tailored via Nonthermal-Plasma-Assisted Synthesis: Enhanced Spectroscopic and Optical Properties with Tuned Electrical Behavior

Abstract

In this study, silver nanoparticles (Ag NPs) were synthesized on the surface of rutile-phase titanium dioxide (R-TiO₂) using a plasma-assisted technique. Comprehensive analyses were conducted to investigate the structural, morphological, optical, and electrical properties of the synthesized nanocomposites. Transmission electron microscopy (TEM) images revealed the uniform decoration of Ag NPs (average size: 29.8 nm) on the R-TiO₂ surface. X-ray diffraction (XRD) confirmed the polycrystalline nature of the samples, with decreased diffraction peak intensity indicating reduced crystallinity due to Ag decoration. The Williamson–Hall analysis showed increased crystallite size and reduced tensile strain, suggesting grain growth and stress relief. Raman spectroscopy revealed quenching and broadening of R-TiO₂ vibrational modes, likely due to increased oxygen vacancies. X-ray photoelectron spectroscopy (XPS) confirmed successful plasma-assisted deposition and the coexistence of Ag⁰ and Ag⁺ states, enhancing surface reactivity. UV-Vis spectroscopy demonstrated enhanced light absorption across the spectral range, attributed to localized surface plasmon resonance (LSPR), and a reduced optical bandgap. Dielectric properties, including dielectric constants, loss factor, and AC conductivity, were evaluated across frequencies (4–8 MHz) and temperatures (20–240 °C). The AC conductivity results indicated correlated barrier hopping (CBH) and overlapping large polaron tunneling (OLPT) as the primary conduction mechanisms. Composition-dependent dielectric behavior was interpreted through the Coulomb blockade effect. These findings suggest the potential of plasma assisted Ag NP-decorated R-TiO₂ nanostructures for photocatalysis, sensor and specific electro electrochemical systems applications.

1. Introduction

Titanium dioxide (TiO₂) is a widely studied semiconductor material with extensive applications in photocatalysis, photovoltaics, sensors, and environmental remediation [1,2]. Among its crystalline forms, rutile titanium dioxide (R-TiO₂) holds distinct advantages over anatase TiO₂ due to its superior thermal stability, lower bandgap (~3.0 eV), and higher refractive index [3]. These properties make R-TiO₂ particularly suitable for high-temperature processes, visible-light-driven photocatalysis, and optoelectronic applications [4]. Furthermore, R-TiO₂ exhibits excellent durability under harsh environmental conditions, which enhances its viability for long-term use in industrial and environmental settings [5]

Despite these advantages, R-TiO₂ faces inherent limitations that restrict its optical, spectroscopic, and electrical performance. The lower surface area of R-TiO₂ compared with anatase TiO₂ reduces the number of active sites available for catalytic reactions, while the rapid recombination of photogenerated electron–hole pairs diminishes its photocatalytic efficiency [6]. Additionally, its wide bandgap restricts its activity under visible light, limiting its utilization in solar-driven applications [7]. To address these challenges, the incorporation of noble metal nanoparticles (NPs), such as silver (Ag), onto the surface of R-TiO₂ has emerged as a promising strategy [8]. The presence of Ag NPs on R-TiO₂ surfaces potentially induces localized surface plasmon resonance (LSPR) and alters the optical, spectroscopic, and electrical characteristics of R-TiO₂.

Various methods have been developed for the deposition of Ag NPs onto R-TiO₂, including wet chemical reduction [9], sol-gel synthesis [10], photodeposition [11], thermal evaporation, and hydrothermal techniques [12]. While these methods have demonstrated efficacy, they often involve challenges such as limited control over nanoparticle size and dispersion, use of hazardous chemicals, and scalability constraints [12]. Recent advancements in research have focused on developing innovative synthesis techniques, with plasma-assisted methods emerging as a highly effective alternative [13].

Plasma is a reactive medium composed of free-charged particles and a variety of chemical agents (radicals) that effectively interact with the exposed surface leading to various processes, reduction, functionalization of nanomaterials as well as dissociation of solution molecules [14,15,16]. The plasma-assisted technique enables precise control over nanoparticle size, short time processing and environmentally friendly [17]. The atmospheric-pressure plasma-aided technique provides a simple, cost-effective, compatible with temperature-sensitive materials, and scalable approach to Ag NP deposition [17,18].

In this study, we developed a novel low-temperature plasma-assisted method for the surface modification of R-TiO₂ with Ag NPs without the need for a reducing agent. The use of nonthermal plasma for the synthesis of Ag NPs offers significant advantages, including a rapid process lasting only a few minutes, while maintaining the structural integrity of R-TiO₂ and enabling the uniform deposition of Ag NPs. Pristine and Ag NP-modified R-TiO₂ samples were characterized using a range of complementary techniques, including scanning and transmission electron microscopy (SEM and TEM), X-ray diffraction (XRD), Raman spectroscopy, X-ray photoelectron spectroscopy (XPS), UV–Vis diffuse reflectance spectroscopy (DRS), photoluminescence, and dielectric properties. This work aims to explore effective strategies for synthesizing plasmonic materials using plasma technology for diverse applications, including photocatalysis, sensing, optoelectronic devices, antibacterial coatings, and photoelectrochemical (PEC) cells.

2. Materials and Methods

2.1. Plasma-Aided Synthesis of Ag/R-TiO₂ NPs

Rutile–titanium oxide R-TiO₂ with 99.9% purity (TI-457110) was procured from Nilaco Corporation, Japan, and used as received. The deposition of Ag NPs on the R-TiO₂ NPs was achieved using a nonthermal atmospheric-pressure argon (Ar) plasma setup, more details in [19,20]. The plasma parameters were an Ar flow rate of 2 L/min and a discharge voltage of 6 kV_pp. In this process, 1 g of R-TiO₂ NPs was ultrasonically dispersed in 100 mL of the Ag NO₃ solution. Deionized water was used as the solvent. The AgNO₃ was prepared in two concentration values, 0.3 wt.% and 1.5 wt.% compared with the R-TiO₂ weight. The R-TiO₂-AgNO₃ solution mixture was then exposed to an Ar plasma plume for 5 min to facilitate AgNO₃ reduction, as shown in Figure 1a.

The emission spectrum of Ar plasma shown in Figure 1b indicates that the plasma plume contained various reactive species, such as hydroxyl radicals (OH) and oxygen and nitrogen species, as well as the UV photons. The presence of these excited species indicates the presence of high-energetic electrons. Similarly, the high-intensity Ar I lines at 696.5 nm and 763 nm indicate the presence of long-lifetime argon metastable (Ar*). When the plasma plume struck the R-TiO₂/AgNO₃ solution mixture, a considerable amount of long-lived argon metastable atoms (Ar*), OH radicals, and high-energetic electrons transferred their energy to the solution molecules, which led to various chemical pathways [17,20] contributing to Ag NP formation.

$${Ar}^{*}+H_{2}O\to e_{eq}^{-}+{H_{2}O}^{+}+Ar$$

(1)

$${H_{2}O}^{+}\to H^{+}+{OH}^{· }$$

(2)

$$e^{-}+H_{2}O \to H^{·}+{OH}^{· }+e^{-}$$

(3)

The Ag NPs could be formatted via reduction of Ag⁺ ions with hydrated electrons ($e_{eq}^{-}$) [18,21].

$$e_{eq}^{-}+{Ag}^{+}\to {Ag}^o$$

(4)

Further, the hydrogen atoms ($H^{·}$) are strong reducing agents that can contribute to the reduction of ${Ag}^{+}$ and Ag NP formation [17]. The participant color changing from white to brown indicated the formation of Ag NPs on the surface of R-TiO₂ NPs. The precipitate was filtered and dried, resulting in a faint brown powder.

2.2. Characterization Techniques

The pure R-TiO₂ and Ag-decorated R-TiO₂ NPs were characterized using various techniques. Scanning electron microscopy (Thermo Fisher Quanta FEG250 SEM, Hillsboro, OR, USA) and high-resolution transmission electron microscopy (HRTEM, Joel-JEM-2100, JEOL Ltd., Tokyo, Japan) were used to investigate the morphology and SAED of the nanocomposites. X-ray diffraction (XRD, Rigaku International Corp., Tokyo, Japan) was utilized to determine the crystal structure,. Micro-Raman spectroscopy (SENTERRA II, Carteret, NJ, USA) with a 532 nm laser was employed to analyze the vibrational structure. X-ray photoelectron emission (XPS-Thermo Scientific, Waltham, MA, USA) with a flood gun was used to examine the chemical composition. The DRS spectra were collected by an Evolution 220 spectrophotometer (Thermo-Fisher, Cambridge, UK). The fluorescence spectra of the samples were recorded using a Cary Eclipse (Agilent, Santa Clara, CA 95051, USA) at an excitation wavelength of 350 nm.

For the electrical and dielectric properties, the powders were pressed into pellets using a uniaxial press of pressure 1.9 × 10⁸ Nm⁻². The samples were well polished to obtain uniform parallel surfaces. Silver paste was applied to the sample surface to create contacts, which were then checked for good conduction. The ac conductivity and dielectric properties were measured using the two-probe method with an LCR meter (Hioki Model IM3536, Nagano, Japan) as a function of frequency (4–8 MHz) and temperature (30–180 °C). The sample temperature was measured using a K-type thermocouple connected to a digit-sense thermometer, with a junction in contact with the sample.

3. Results and Discussion

3.1. Morphology and Structure of R-TiO₂ and Ag/R-TiO₂ Nanoparticles

Figure 2a–c presents the FESEM images of pristine R-TiO₂ and R-TiO₂/Ag NP nanocomposites. The images clearly show agglomerated nanoparticles with irregular shape and sizes. The high-resolution TEM images of the R-TiO₂/Ag nanocomposites in Figure 2d,e distinctly reveal the deposition of fine Ag NPs over the surface of the R-TiO₂ NPs, particularly in the sample containing 1.5% wt. of Ag nanocomposites.

The Ag NPs were uniformly distributed, with an average size of 29.8 nm. The HR-TEM images also demonstrate that the R-TiO₂/Ag nanocomposites exhibited nanodisc or nanoplate morphology, with sizes ranging from 128.7 to 181.9 nm. Figure 2f displays the lattice fringe pattern of the sample R-TiO₂ + Ag nanocomposite (1.5 wt.% sample). The observed lattice spacing of 0.19 nm was less than the standard interplanar spacing d = 0.32 nm for the R-TiO₂ (110) plane [22], suggesting it was associated with the (111) plane of R-TiO₂ [22] or the (200) plane of the Ag NPs [23]. Furthermore, the SAED pattern (inset of Figure 2f) exhibited a concentric ring–dot pattern that corresponds to diffraction planes of R-TiO₂, confirming that the plasma-assisted Ag NP-decorated R-TiO₂ NPs preserved the structural integrity of the R-TiO₂.

Figure 3 shows the XRD patterns of pure rutile-phase TiO₂ (R-TiO₂) and R-TiO₂ decorated with Ag nanoparticles (Ag NPs) at various weight percentages. The XRD pattern of R-TiO₂ exhibited sharp peaks at 2θ ≈ 27.6°, 36.3°, 39.4°, 41.4°, 54.5°, 56.8°, and 62.9°, corresponding to the (110), (101), (111), (210), (211), (220), and (002) planes of the “tetragonal” rutile phase of TiO₂ (ICDD pattern no: 00–021–1276). These patterns confirm the polycrystalline nature of the R-TiO₂ [23].

Furthermore, as the weight percentage of Ag NPs increased, the intensity of the R-TiO₂ diffraction peaks decreased, suggesting a reduction in crystallinity. Notably, no distinct XRD peaks corresponding to Ag were observed in the patterns of R-TiO₂ decorated with Ag NPs because of the fine distribution of Ag NPs, which was below the detection limit of the XRD instrument. The inset in Figure 3 shows zooming in for the strongest peak of 110. Clearly, the peak position shifted toward the lower 2θ value as Ag content increased in wt.%. The decrease in XRD peak intensity suggests partial disruption of crystalline order or enhanced scattering effects.

Williamson–Hall Plot

The crystallite size (G) and micro strain (ε) were calculated for the synthesized samples using the Williamson–Hall (W-H) method, which distinguishes peak broadening contributions from crystallite size and strain based on their different dependencies on the Bragg angle (θ). Broadening due to crystallite size follows the Debye–Scherrer equation:

$$G={\displaystyle \frac{0.9\times \lambda }{\beta \times \cos \vartheta }}$$

(5)

where λ is the X-ray wavelength (1.54 Å) and β is the full width at half-maximum (FWHM). The strain-induced broadening, caused by crystal distortion, is expressed as β_s = 4 × ε × tanθ. The total broadening (β_hkl) is the sum of the size and strain contributions [24]:

$${\beta }_{hkl}\times \cos \theta ={\displaystyle \frac{0.9\times \lambda }{G}}+4\times \epsilon \times \sin \theta$$

(6)

This equation assumes isotropic crystal properties. ${\beta }_{hkl}\cos \theta$ is calculated for all diffraction peaks and plotted against $4\epsilon \sin \theta$, as shown in Figure 4. The crystallite size is obtained from the y-intercept, and the strain from the slope, of linearly fitted W-H plots. The positive slope indicates tensile strain in the samples. The calculated values of G and ε are summarized in

Table 1. XRD pattern analysis of R-TiO₂ and R-TiO₂/Ag nanocomposites.

Sample	Crystal Size (nm)	Micro Strain e (×10−3)
R-TiO2	43.36	0.6433
R-TiO2 + Ag 0.3 wt.%	69.29	0.505
R-TiO2 + Ag 1.5 wt.%	85.19	0.499

.

As inferred from Table 1, the deposition of Ag NPs on R-TiO₂ led to an increase in crystal size and a decrease in tensile strain. The Ag NPs reduce grain boundary energy, which supports grain growth. Furthermore, the reduction in strain observed in the R-TiO₂/Ag NP system can indeed be attributed to strain relaxation mechanisms. Specifically, Ag NPs may help relieve structural stress in the lattice by promoting dislocation motion or grain boundary relaxation, or by filling lattice vacancies. These mechanisms could contribute to reduction in strain and stabilization of the TiO₂ crystal structure. Although the specific strain-relaxation mechanisms were not directly explored in this study, the observed changes in lattice strain and grain growth support the idea that Ag NPs play a role in these processes, which can enhance the material’s properties for applications such as photocatalysis or fluorescence enhancement.

3.2. Vibrational and Chemical Composition Analysis

The Raman spectra of pristine R-TiO₂ and R-TiO₂/Ag nanostructures with varying Ag wt.% are presented in Figure 5. For pristine R-TiO₂, as depicted in Figure 5a, prominent peaks appeared at 447 cm⁻¹ and 610 cm⁻¹, corresponding to the E_g and A_1g Raman modes characteristic of the rutile TiO₂ phase [25]. Additionally, a subtle peak at 143 cm⁻¹ and a broad hump at 230 cm⁻¹ were observed, attributed to the B_1g mode and multiphonon scattering, respectively [26]. These findings aligned well with the XRD results. The E_g mode corresponded to the symmetric stretching vibration of O–Ti–O in TiO₂, while the B_1g and A_1g modes were associated with the symmetric and antisymmetric bending vibrations of O–Ti–O bonds, respectively. The multiphonon scattering observed at 230 cm⁻¹ arose from lattice disorder, indicating the nanometer-scale nature of the TiO₂, consistently with the observations from the SEM/TEM images. The Raman spectra of the R-TiO₂/Ag nanostructures, shown in Figure 5b,c, displayed the characteristic Raman modes of pristine R-TiO₂. However, these modes exhibited a notable reduction in intensity and slight broadening, suggesting alterations in the vibrational properties of the R-TiO₂.

No Raman modes associated with Ag nanoparticles (NPs) were observed because of the low loading of Ag NPs on the R-TiO₂ surface. The quenching of Raman modes in the R-TiO₂/Ag nanostructures was attributed to the absorption of 532 nm laser light by the Ag NPs decorating the R-TiO₂ surface [14]. Additionally, the broadening of the Raman modes may be linked to an increase in oxygen vacancy content or changes of the R-TiO₂ NPs size [27].

Figure 6A(a) presents the wide-scan XPS survey spectra of pristine R-TiO₂ and R-TiO₂ decorated with Ag nanoparticles (NPs). The survey spectrum of pristine R-TiO₂ showed prominent peaks at 458 eV and 529.2 eV, corresponding to Ti 2p and O 1s, respectively. In contrast, the XPS survey spectra of R-TiO₂ decorated with Ag NPs (Figure 6A(b,c)) exhibited an additional peak at 368.9 eV, attributed to Ag 3d, confirming the successful plasma-assisted deposition of Ag NPs onto the R-TiO₂ surface [28]. As expected, the intensity of the Ag 3d peak increased with the Ag wt.% ratio, signifying a higher Ag loading.

Figure 6B(a) displays the high-resolution Ti 2p spectrum of R-TiO₂, which showed two primary peaks at approximately 457.7 eV and 463.3 eV, corresponding to Ti 2p₃/₂ and Ti 2p₁/₂, respectively. Deconvolution of the Ti 2p₃/₂ peak revealed two components: one at 456.3 eV, attributed to Ti³⁺, and another at 457.68 eV, assigned to Ti⁴⁺. Similarly, the Ti 2p₁/₂ peak was resolved into two subpeaks at 461.2 eV and 463.4 eV, corresponding to Ti³⁺ and Ti⁴⁺, respectively [29]. For the R-TiO²/Ag NP samples (Figure 6B(b,c)), the Ti 2p₃/₂ peak deconvolves into three components, including a new peak at a lower binding energy of 454.2 eV, attributed to Ti²⁺. Similarly, the Ti 2p₁/₂ peak resolved into three subpeaks, corresponding to Ti⁴⁺, Ti³⁺, and Ti²⁺. Notably, the relative intensity of the Ti³⁺ component increased significantly upon Ag NP deposition, indicating enhanced titanium reduction. Additionally, the binding energy positions of the Ti 2p₃/₂ and Ti 2p₁/₂ peaks shifted to lower values in the R-TiO₂/Ag NP samples compared with pristine R-TiO₂. This shift suggests that Ag NPs facilitate electron transfer to the TiO₂ surface, promoting the stepwise reduction of Ti⁴⁺ to Ti³⁺ and Ti²⁺. The increased local electron density around titanium atoms, driven by this electron transfer, accounts for the observed decrease in binding energy. The increased concentrations of Ti³⁺ and Ti²⁺ species in the R-TiO₂/Ag nanocomposite system suggest enhanced defect states, which can influence charge transport dynamics [30].

Figure 6C(a) presents the O 1s spectrum of pristine R-TiO₂, deconvoluted into three distinct peaks at 527.5 eV, 529.3 eV, and 531.1 eV. Based on the literature, the binding energies (B.E.) for O 1s in R-TiO₂ correspond to near-surface oxygen occupancy (527.0 eV), bridging oxygen (531.0 eV), and on-top oxygen (532.2 eV) [31]. The observed peak at 527.5 eV matched near-surface oxygen occupancy, indicating the presence of O-Ti³⁺ bonds. The peak at 529.3 eV was attributed to threefold near-surface oxygen occupancy, corresponding to O-Ti⁴⁺ bonds, while the peak at 531.1 eV was associated with bridging oxygen, representing adsorbed hydroxyl (OH) groups or the influence of oxygen vacancies [31]. Figure 6C(b,c) presents the O 1s spectra of R-TiO₂ after Ag NP deposition. The intensities of the O-Ti³⁺ and OH components increased, while the O-Ti⁴⁺ component slightly decreased, indicating modifications in the surface chemistry and electronic structure induced by Ag NPs. These changes are consistent with the Ti 2p XPS data, suggesting that Ag NPs promote the partial reduction of Ti⁴⁺ to Ti³⁺ and enhance surface reactivity in R-TiO₂.

The high-resolution Ag 3d spectra for the R-TiO₂/Ag NP samples, shown in Figure 6D(b,c), revealed two peaks at approximately 368.8 eV and 374.5 eV, corresponding to Ag 3d₅/₂ and Ag 3d₃/₂, respectively [28]. Deconvolution of the Ag 3d₅/² peak revealed two components at 367.0 eV and 368.8 eV, assigned to Ag⁰ and Ag⁺, respectively. Similarly, the Ag 3d₃/₂ peak was resolved into two subpeaks at 372.8 eV and 374.8 eV, corresponding to Ag⁰ and Ag⁺ [29]. These results confirm the coexistence of both metallic (Ag⁰) and oxidized (Ag⁺) states in the R-TiO₂/Ag NP samples. The presence of Ag⁰ nanoparticles facilitates electron transfer to the R-TiO₂ surface, promoting the partial reduction of Ti⁴⁺ to Ti³⁺ and Ti²⁺, while Ag⁺ ions interact with the TiO² surface, stabilizing oxygen vacancies and defect sites. This synergy between Ag⁰ and Ag⁺ enhances the electronic structure and surface reactivity of the composite [32]. In conclusion, the XPS analysis demonstrates that the R-TiO₂/Ag NPs composites exhibited tailored surface chemistry, making them promising candidates for various applications.

3.3. UV-Vis and Fluorescence Spectroscopy

Diffuse reflectance spectroscopy was conducted, and the resulting spectra was transformed into absorbance using the Kubelka–Munk (F(R)) equation given as:

$$F\left(R\right)={\displaystyle \frac{{\left(1-R\right)}^2}{2\times R}}$$

(7)

where R represents the reflectance. Figure 7A presents the Kubelka–Munk (F(R)) function for pristine R-TiO₂ and Ag NP-decorated R-TiO₂ nanostructures. Pristine R-TiO₂ absorbed only the UV portion of the UV-Vis spectrum (300–800 nm), corresponding to the intrinsic excitonic absorption of R-TiO₂. In contrast, the F(R) spectra of R-TiO₂ decorated with Ag NPs showed an improvement in UV-Vis absorption across the entire spectral range (300–800 nm). Notably, reflectance decreased with increasing Ag NP wt.%, thereby justifying the observed improvement in absorption throughout the UV-Vis range [33]. Additionally, a localized enhancement in absorption around 510 nm was observed in the spectrum of the R-TiO₂ + Ag 1.5% wt. sample. This feature was attributed to the localized surface plasmon resonance (LSPR) effect of Ag NPs anchored on R-TiO₂ NP surfaces [34].

The optical bandgap Eg values were estimated using the Wood and Tauc method [35], based on the relation:

$$F\left(R\right)·h\nu ={\left(h\nu -E_g\right)}^n$$

(8)

where n = 2 for indirect transition [36]. The results are shown in Figure 7B. The pristine R-TiO₂ NPs exhibited an E_g = 3 eV, which decreased to 2.90 eV and 2.82 eV for R-TiO₂/Ag nanostructures with. 0.3 wt.% and 1.5% wt.%, respectively. The E_g value for pristine R-TiO₂ aligned closely previously reported values [37]. Meanwhile, the reduced E_g values for R-TiO₂/Ag nanostructures highlight the synergistic effects of Ag NP decoration, which enhance Ti³⁺ and/or oxygen vacancy defects. These defects introduce intermediate energy levels within the band gap, leading to a reduction in the E_g value.

Figure 8 shows the fluorescence spectra of the examined samples at room temperature. The fluorescence of pristine R-TiO_2, shown in Figure 8a, exhibited a weak, broad band in the UV-Vis region (390–500 nm), with peaks at 394, 409, and 433 nm. The peak at 394 nm was attributed to recombination from conduction to valence band transition [38]; the peaks at 409 and 433 nm were related to emission from shallow defects (oxygen vacancies) as well as deeper defect states associated with Ti³⁺ structural distortions [39]. The XRD (Figure 3) and XPS (Figure 6) results confirmed the presence of oxygen vacancies and Ti³⁺ in the pristine R-TiO₂ NPs. The low-intensity emission peaks at 409 and 433 nm reflected the high crystallinity of the R-TiO₂ NPs, as inferred from Figure 3. The Ag NP-modified R-TiO₂ nanostructure exhibited enhanced fluorescence intensity compared with that of pristine R-TiO₂. This enhancement was due to the surface plasmon resonance (SPR) of the Ag NPs, which improves light absorption and promote electron–hole pair generation in TiO₂ surface. The fluorescence peak intensity of the R-TiO₂/Ag nanostructures was dependent on the Ag wt.%, as seen in Figure 8b,c. The fluorescence intensity of the R-TiO₂/Ag (1.5 wt.%) nanocomposite was lower than that of the R-TiO₂/Ag (0.3 wt.%) nanocomposite, with the decrease in intensity at higher Ag NP content attributed to the absorption of 532 nm laser light by the Ag NPs decorating the surface of the R-TiO₂ NPs [14]. Furthermore, new fluorescence peaks at 381 nm, 522 nm, and 555 nm were observed in the R-TiO₂/Ag nanocomposite samples. The emission at 381 nm was due to near-band-edge emissions, while the green emissions at 522 and 55 nm were associated with deep-level defects due to oxygen vacancies, Ti³⁺ trap levels, and LSPR-enhanced recombination [38,40].

3.4. Dielectric Properties

The study of dielectric properties of materials is essential for understanding their behavior under electric fields, leading to improved designs and applications in electronics, telecommunications, biophysics, and energy storage [41]. The dielectric constant, or relative permittivity, quantifies how much a material diminishes the electric field in its surroundings. It is defined as the ratio of the material’s permittivity to the permittivity of free space. In other words, it measures a material’s capacity to store energy [42,43].

Figure 9a–c shows the variation in the dialectic constant (ε′) for R-TiO₂, R-TiO₂/0.3%Ag, and R-TiO₂/1.5%Ag as a function of frequency (4–8 MHz) at different temperatures (20–230 °C). The figure demonstrates the typical behavior of dielectric materials, where the dielectric constant decreases with increasing frequency [44]. Dielectric polarization consists of diverse types, including ionic, dipolar, electronic, and interfacial polarizations. Each type responds differently to the frequency of the applied electric field [45]. At lower frequencies, all types of polarization can effectively respond to the periodically varying electric field. However, as the frequency increases, certain types of polarizations fail to follow the frequency of the electric field, so the net polarizations of the material diminish, resulting in a decreased dielectric constant [46]. Additionally, Figure 9a–c shows that the dielectric constant increased (ε′) with increasing temperature, reaching a maximum value at 80 °C before decreasing. Studying the imaginary part of the dielectric constant (dielectric loss factor ε″) is crucial for enhancing devices’ performance and understanding the underlying mechanisms of energy dissipation in materials [46].

Figure 10a–c illustrates the variation of dielectric loss factor (ε″) for R-TiO₂, R-TiO₂/0.3%Ag, and R-TiO₂/1.5%Ag as a function of frequency (4–8 MHz) at various temperatures (20–230 °C). As shown, the dielectric loss factor (ε″) decreased with increasing frequency. Dielectric materials consist of molecules possessing electric dipoles that orient in response to an applied electric field. At low frequencies, these dipoles can adequately respond to the alternating field, resulting in energy absorption and high dissipation. However, as the frequency increases, the dipoles may not be able to keep pace with the rapidly changing field, leading to reduced energy loss [47]. Furthermore, Figure 10a–c demonstrates that the dielectric loss factor (ε″) increased with temperature, reaching a maximum value at 80 °C, and then decreased.

The variation in conductivity as a function of frequency at the different temperatures for R-TiO₂, R-TiO₂/0.3%Ag, and R-TiO₂/1.5%Ag is illustrated in Figure 11a–c. The experimental data in Figure 11a–c indicate the presence of two distinct regions based on the curves. The first region, observed at low frequencies, was characterized by a plateau representing the direct current conductivity. In the second region, at high frequencies, conductivity gradually increased with increasing frequency [48,49]. The behavior of conductivity versus frequency can be explained by the nature of charge carriers (electrons or ions) in many materials, especially oxides. These carriers are often not completely free but instead localized or trapped in specific regions. When an alternating current (AC) field is applied, these charge carriers can oscillate or hop between localized states. As the frequency of the field increases, the carriers have less time to get trapped, resulting in increased mobility and higher conductivity [49]. Additionally, polarization can reduce conductivity at low frequencies by hindering the transport of charge carriers. However, as the frequency increases, polarization effects may not be able to keep pace with the applied field, leading to reduced hindrance and, consequently, increased conductivity [49]. The electrical conductivity can be described by

$${\sigma }_{Tot}={\sigma }_{DC}+\mathsf{A}{\omega }^n$$

(9)

${\sigma }_{DC}$ expresses direct current (DC) conductivity, while $\mathsf{A}{\omega }^n$ represents AC conductivity (${\sigma }_{AC}$). Here, $\mathsf{A}$ is a constant specific to a given temperature, and n is an exponent dependent on frequency and temperature. Analyzing the behavior of the exponent n versus temperatures provided insights into the conduction mechanisms and their corresponding models. In the small polaron tunneling model (SPTM), the value of n increased with temperature because of thermally assisted hopping of small polarons [46,50]. According to the overlapping large polaron tunneling (OLPT) mechanism, the exponent n is influenced by both temperature and frequency. It decreases with a rising temperature until reaching a minimum, after which it increases as the temperature continues to rise [50,51]. In contrast, the correlated barrier hopping (CBH) model suggests that n decreases with increasing temperature, representing charge carrier hopping between localized sites over a potential barrier [46,50]. The quantum mechanical tunneling model (QMTM) indicates that n is approximately 0.8 and either slightly increases with temperature or remains temperature independent [50]. As shown in the insets in Figure 11a–c, the frequency exponent n decreased with increasing temperature up to 80 °C, after which it increased. This suggests that the conduction mechanism may correlate barrier hopping (CBH) from room temperature to 80 °C while the temperature range greater than 80 °C is governed by the small polaron tunneling model (SPTM) mechanism or an overlapping large polaron tunneling (OLPT) mechanism.

Figure 11a–c shows that conductivity increased with temperature, reaching a maximum value at 80 °C before decreasing. The influence of temperature on the AC conductivity of R-TiO₂/Ag nanocomposites is crucial for their application in electronic devices. As temperatures increase, the mobility of charge carriers improves, enhancing AC conductivity via mechanisms such as thermally activated hopping and variations in ionic mobility [49]. At lower temperatures, AC conductivity exhibits frequency independence due to restricted charge mobility. However, as thermal energy exceeds lattice barriers at elevated temperatures, a notable increase in conductivity occurs, associated with increased hopping activities among charge carriers. Furthermore, the presence of impurities and defect states in R-TiO₂/Ag nanocomposites significantly influences AC conductivity at higher temperatures. The dynamics of these defects can significantly influence charge transport.

Impedance spectroscopy is a widely employed method for investigating the electrical properties of nanomaterials. The process of complex impedance spectroscopy is employed to differentiate the reactive and resistive components of electrical characteristics. The complex impedance of a material can be expressed as the sum of its real and imaginary components, which can be written as

$$Z=Z^{\prime }+Z^{″}$$

(10)

The variation of complex impedance with frequency at different temperatures for pure R-TiO₂, R-TiO₂/0.3%Ag, and R-TiO₂/1.5%Ag nanocomposites is depicted in Figure 12. It is evident that as frequency increased, Z decreased for all samples. The behavior of complex impedance with temperature was the same for ac conductivity, where it decreased with temperatures to 80 °C and thereafter increased.

Figure 13a–c shows the composition dependence of ε′, ε″, and AC conductivity at room temperature as a function of frequency. All parameters, ε′, ε″, and AC conductivity, decreased with increasing Ag content in R-TiO₂/Ag nanocomposites. The decreases in ε′, ε″, and AC conductivity with the incorporation of nano-Ag particles can be explained via the Coulomb blockade effect [52]. Coulomb blockade islands are small conductive areas (often nanoscale) that are electrically isolated from their surroundings by insulating barriers [52,53]. These islands can affect a material’s electrical and dielectric characteristics, especially in nanostructured systems such as composites of TiO₂ and silver nanoparticles [52].

4. Conclusions

Morphological analysis through FESEM and HR-TEM confirmed the successful deposition of Ag nanoparticles (NPs) on the R-TiO₂ surface, with an average size of 29.8 nm. The nanocomposites exhibited distinct nanodisc or nanoplate morphology, ranging from 128.7 to 181.9 nm, with a lattice fringe pattern indicating reduced interplanar spacing due to the influence of Ag NPs. XRD analysis indicated the polycrystalline nature of rutile-phase TiO₂ and demonstrated that increasing the Ag NP content led to reduced crystallinity and peak intensity due to enhanced scattering effects. Williamson–Hall plot analysis revealed an increase in crystallite size and a reduction in tensile strain with higher Ag content, indicating that Ag NPs contribute to relieving structural stress and promoting grain growth. Raman spectroscopy confirmed the structural integrity of R-TiO₂, albeit with quenching and broadening of modes due to Ag NPs, linked to oxygen vacancy content and particle size changes. XPS analysis provided evidence of successful plasma-assisted deposition of Ag NPs and the creation of Ti²⁺ and Ti³⁺ states, enhancing the material’s electronic properties. UV-Vis spectroscopy demonstrated a broadened absorption spectrum and a bandgap reduction attributed to the localized surface plasmon resonance effect of Ag NPs. The dielectric constant and dielectric loss factor decreased with frequency, while AC conductivity increased. The dielectric constant, dielectric loss factor, and AC conductivity were enhanced with increasing temperatures up to 80 °C, and after that, they decreased. The results demonstrate that the decreased dielectric constant, dielectric loss factor, and AC conductivity with the incorporation of nano-Ag particles can be explained as the Coulomb blockade effect. Coulomb blockade islands at the nanoscale are small conductive regions that are electrically isolated from their surroundings by insulating barriers. These islands can affect a material’s electrical and dielectric characteristics. These findings collectively underscore the potential of R-TiO₂/Ag nanocomposites for applications in photocatalysis, sensors, and electrochemical systems, where efficient electron transport and tailored conductivity are critical.