Nanocrystalline Surface Layer of WO₃ for Enhanced Proton Transport during Fuel Cell Operation

Abstract

High ionic conductivity in low-cost semiconductor oxides is essential to develop electrochemical energy devices for practical applications. These materials exhibit fast protonic or oxygen-ion transport in oxide materials by structural doping, but their application to solid oxide fuel cells (SOFCs) has remained a significant challenge. In this work, we have successfully synthesized nanostructured monoclinic WO₃ through three steps: co-precipitation, hydrothermal, and dry freezing methods. The resulting WO₃ exhibited good ionic conductivity of 6.12 × 10⁻² S cm⁻¹ and reached an excellent power density of 418 mW cm⁻² at 550 °C using as an electrolyte in SOFC. To achieve such a high ionic conductivity and fuel cell performance without any doping contents was surprising, as there should not be any possibility of oxygen vacancies through the bulk structure for the ionic transport. Therefore, laterally we found that the surface layer of WO₃ is reduced to oxygen-deficient when exposed to a reducing atmosphere and form WO_3−δ/WO₃ heterostructure, which reveals a unique ionic transport mechanism. Different microscopic and spectroscopic methods such as HR-TEM, SEM, EIS, Raman, UV-visible, XPS, and ESR spectroscopy were applied to investigate the structural, morphological, and electrochemical properties of WO₃ electrolyte. The structural stability of the WO₃ is explained by less dispersion between the valence and conduction bands of WO_3−δ/WO₃, which in turn could prevent current leakage in the fuel cell that is essential to reach high performance. This work provides some new insights for designing high-ion conducting electrolyte materials for energy storage and conversion devices.

1. Introduction

Fuel cells (FCs) provide a clean and efficient way to generate electricity from H₂ and hydrocarbon fuels. Due to high operating temperatures, solid oxide fuel cells (SOFCs) can also be used in combined heat and power applications. One of the main challenges for SOFCs is developing high oxygen ion (O²⁻) conductivity of electrolyte materials. Structural doping has been remained a general methodology for developing high ionic conductivities [1,2]. In this methodology, the host cations could often be replaced by a lower valence state, which produces an oxygen-deficient structure to conduct O²⁻ (For example, Zr⁴⁺ or Ce⁴⁺ are replaced with Y³⁺ and Sm³⁺) [3,4]. However, this approach does not significantly enhance fuel cell performance at low operating temperatures due to limited ionic conductivity [4].

Protonic conducting fuel cells (PCFCs), a sub-class of SOFCs, have emerged with great potential for lowering the operating temperature to the range of 400–700 °C. PCFCs can achieve higher fuel utilization by producing the water on the cathode side and avoiding fuel dilution effectively [5]. Another advantage of PCFCs is that many proton-conducting electrolytes such as BaZr_0.8Y_0.2O_3−δ, BaZr_0.1Ce_0.7Y_0.2O_3−δ, and Yb-doped BaZr_0.1Ce_0.7Y_0.2O_3−δ offer high protonic conductivity at an intermediate operating temperature of 600–800 °C [6,7,8]. Further reduction of the operating temperature, (e.g., 400–600 °C) reduces the power output sharply. If PCFCs could operate in the range of 400–600 °C, it would help prolong the lifetime. The main critical obstacle to achieving high-performance low-temperature PCFCs is the limited proton conductivity of the electrolyte and cathode materials for oxygen reduction reaction (ORR), both of them perform poorly at low temperatures [9,10,11,12,13,14].

However, obtaining the benchmark in proton conductivity (H⁺) for electrolytes and their chemical stability is a big challenge. The best proton-conducting electrolytes, such as doped barium cerates (BaCeO_3−δ), lead to poor CO₂ tolerance and rapid decomposing due to forming carbonates at a temperature higher than 600 °C, which is a significant limitation for their use in the device [15]. Recently, next-generation proton-conducting electrolytes such as MTO₄, where M = lanthanide/alkali metals and T = Nb, W, Mo, Mn, shows high CO₂ tolerance [16,17]. Moreover, the semiconductor-based electrolyte has been reported to deliver the best performance at a low operating temperature [10,11,12,13,14]. The doped BaCeO_3−δ only exhibits dominant H⁺ conductivity under wet conditions or reduced atmosphere and shows significant p-type conductors under oxidizing conditions [15]. WO₄/WO₃ based materials almost remain pure H⁺ conductors in both wet and oxidizing conditions, although they exhibit moderate H⁺ conductivity. However, WO₃ to be a direct bandgap semiconductor with less disseminative valence and conduction bands has shown considerable interest for multiple applications in energy devices. WO₄/WO³⁻ type oxides exhibit a high oxide ion conduction, e.g., Pb_0.9Sm_0.1WO_4+δ shows a conductivity of ∼2 × 10⁻² S cm⁻¹ at 800 °C, which is comparable to that of YSZ (3.6 × 10⁻² S cm⁻¹ at 800 °C) [18].

Controlling grain boundary conduction (GBs) into nanocrystalline material is an emerging field. Therefore, we have synthesized nanocrystalline WO₃ in three steps to obtain the fine morphology to effectively modulate its electrical properties. We have synthesized nano-structure WO₃ by combining three steps following one by one, such as (i) co-precipitation, (ii) hydrothermal, and (iii) dry freezing method. The structural, chemical and morphological analysis of prepared WO₃ powders is analyzed. However, in situ formations of WO_3−δ/WO₃ heterostructure when exposed to H₂- side in the fuel cell that could help migrate ions accompanied with the creation of the internal cavity (charge accumulation and depletion region) could be explained by the “surface locking” effect. The prepared WO₃ spontaneously facilitated ionic transport exhibiting high ionic conductivity of 6.12 × 10⁻² S cm⁻¹ and excellent power density of 478 mW cm⁻² at 550 °C. The results demonstrate that the approach proposed is helpful for developing new materials with unique functionalities for advanced PCFCs/SOFCs.

2. Materials and Methods

2.1. Material Preparation Methods

The powders of WO₃ were prepared in three steps. Initially, an appropriate amount of ammonium tungstate with a chemical formula of (NH₄)10H₂(W₂O₇)₆ was bought from Shanghai Macklin Biochemical Co., Ltd., Shanghai, China (purity 98.5%) was dissolved into deionized water to prepare the 0.5 mol/L solutions. In parallel, the 1 mol/L of Na₂CO₃ was prepared into 100 mL water; afterward, the Na₂CO₃ solution was added dropwise to the above solution under continuous stirring resulted in white milky precipitates. The obtained precipitates were moved to a Teflon autoclave bottle for hydrothermal treatment at 180°C for 6 h at 180 in a vacuum oven. Afterward, the precipitates were collected to wash and filter numerous times with distilled water and absolute ethanol to eliminate surface-adsorbed water. The resultant precursors were dry freezes at −40 °C for 4 h. Furthermore, it was followed by a vacuum at 1.0 P to cool down to room temperature to remove the absorbed water. Moreover, obtained powders were calcined at 750 °C for 6 h to obtain WO₃ nanoparticles. Furthermore, tungsten trioxide (WO₃) purchased from Shanghai Macklin Biochemical Co., Ltd., Shanghai, China (99.8%) was used for comparative study to synthesize WO₃.

2.2. Material Characterizations and Electrochemical Measurements

The Bruker D8 advanced X-ray diffractometer (Bruker, Hanau, Germany, with Cu Kα radiation, λ = 1.5418 Å) was employed to study the crystalline structure of the synthesized WO₃ powders in 2θ ranges from 10–70°. Furthermore, crystal structure, microstructure, and chemical composition of prepared WO₃ were studied via FEI Tecnai GI F30 and JEOL JSM7100F (resolution transmission electron microscopy; HR-TEM and field emission scanning electron microscope). X-ray photoelectron spectroscopy (XPS, Physical Electronics Quantum 2000, Al Kα X-ray source) was engaged for chemical oxidation states where the raw data was analyzed through CASA XPS software. Raman spectra were carried out on NT-MDT (Russia) Raman spectrometer at room temperature, with 532 nm solid-state laser as the excitation source and laser power of 20 mW. The UV-Vis 3600 spectrophotometer was used to measure the absorbance spectrum. The electron spin resonance (ESR) measurements were performed using JES X320 (JEOL). Keithley 2400 source meter was used to measure dc-four probed method conductivity.

2.3. Fuel Cells Fabrication and Electrochemical Measurements

We utilized the dry-pressing method to fabricate the SOFC cell with dense electrolyte and porous electrodes. The WO₃ powder was pressed between the symmetrical electrodes of Ni_0.8Co_0.15Al_0.05LiO_2−δ (NCAL) to form the electrolyte membranes as thin as 285 µm using low filling density powders. In detail used as an electrode was purchased from Bamo Sci. & Tech. Joint Stock Company Ltd., Tianjin, China. In detail, a slurry of NCAL was prepared and painted on porous Ni-foam followed by drying at 120 °C for 2 h. The synthesized and commercially purchased WO₃ powders were pressed between the prepared NCAL electrodes under 240 MPa to produce the solid cells. The designed fuel cell pellet’s diameter, thickness, and active area were about 13 mm, 1.5 mm, and 0.64 cm², respectively. DC electronic load instrument (ITECH8511, ITECH Electrical Co., Ltd., New Taipei, Taiwan) was employed to determine the fuel cell performance of fabricated cells under the H₂ as fuel and air as an oxidant with a flow rate of 100–110 mL/min and 100 mL/min, respectively. The I-V (current density-voltage) and I-P (current density–power density) were recorded to present the electrochemical characteristics of the prepared FC devices. Electrochemical impedance spectroscopy (EIS) was measured by Gamry Reference 3000, USA workstation under the fuel cell open-circuit voltage (OCV) conditions in the frequency range of 0.1 to 106 Hz with 10 mV of dc signal. ZSIMPWIN software was used to fit the model circuit with obtained EIS data.

3. Results & Discussion

3.1. Structural and Compositional Study

Figure 1a shows the XRD pattern of WO₃, whose main diffraction peaks are located at 2θ° of the 22.8, 23, 24, 26, 28, 33, 33.5, 34.5, 41, 50.5, 56, which are corresponding to (002), (020), (200), (120), (112), (202), (022), (220), (222), (232) and (114) planes of monoclinic crystal structure of WO₃ (JCPDS no 43–1035), with space group Pi (C~), and lattice of a = 7.309, b = 7.522, c = 7.678, α = 88.81, β = 90.92, γ = 90.93. There are no extra peaks observed in the patterns, eliminating the possibility of additional phase formation except the monoclinic WO₃ phase. To verify crystallography of synthesized WO_3, Nano-STAR (Bruker-AXS was used to measure the wide-angle scattering. Figure 1b,c shows a wide-angle plot in the 2θ range from 2°–45° and its corresponding 3-D mapping image, respectively. The wide-angle confirms the high purity monoclinic crystal structure of WO₃ measured by XRD. Moreover, the crystallographic structure of WO₃ is confirmed by HR-TEM, as shown in Figure 1d,e. The d-spacing values calculated using a digital micrograph are 0.236 and 0.213 nm, which can be nominated to (200) and plane (020) planes, respectively. These planes could also be confirmed by a selected area electron diffraction (SAED) pattern, as shown in Figure 1f [19,20].

Figure 2 shows nano-scaled HR-TEM and s STEM images for synthesized WO₃, where sophisticated and fine nano-particles with a particle size of <50 nm can be seen clearly. The restrain of agglomerates and the growth of particles in WO₃ can be attributed to the purification of the synthesis method we have used. Moreover, a line scan using high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) was used to confirm the elemental distribution at the particle level. Figure 2c shows a HAADF-STEM image with elements mapping of W and O, where a homogenous chemical distribution can be observed. Figure 2d shows the chemical distribution along with the line scan compared to the HAADF image in the 0–100 nm range. Figure 3f shows the energy dispersive spectroscopy (EDS) image measured using the HAADF-STEM image. The results mentioned above manifest that the particles are at nano-level and are well connected, which leads to enhancement of the surface area with more active sties to enhance the overall performance of fuel cell devices [21,22].

3.2. Electrochemical Performance Measurements

The electrochemical performance of the prepared WO₃ in comparison to commercially purchased WO₃ in a fuel cell as an ion-conducting electrolyte was evaluated. Figure 3a displays typical current-voltage (I–V) and corresponding power density (I-P) characteristics of fuel cells using commercial WO₃ and our synthesized WO₃ material as an electrolyte at 450–550 °C. The OCV of 1.03 V and the maximum power density (Pmax) of 418 ± 2% mW cm⁻² were achieved for our prepared WO₃ electrolyte, while commercially purchased WO₃ exhibited (Pmax) of only 260 mW cm⁻². The OCV (e.g., 1.03 V at 550 °C) agrees with the Nernst theoretical potential at this temperature, showing that the ionic transfer number is near unity and WO₃ acts electronically insulating [23,24,25]. Moreover, our synthesized WO₃ electrolyte fuel cell displayed good electrochemical performance even at low temperatures, as shown in Figure 3b. It exhibited peak power densities of 360 and 310 mW cm⁻² at 500 °C and 450 °C, respectively. The enhanced electrochemical performance of synthesized WO₃ over commercially purchased WO₃ suggests the vital role of synthesis method and nanostructured morphology, where the surface layer of WO₃ can easily be reduced and produce the strong WO_3−δ/WO₃ heterostructure for creating the surface path for easy transport of H⁺ when it’s exposed to H₂ in fuel cell conditions [20]. During the online electrochemical process, the prepared WO₃ material could facilitate the better intercalation of protons because of the high surface-active area and easy reduction of the surface layer. Therefore, obtained results suggested that the nanostructured stoichiometry synthesized by three steps synthesis could be helpful to control the surface structural properties. In this way, the WO₃ surface kinetically favors the heterolytic dissociation of H₂ to form W–H and O–H species. The resulting W–H further evolves to the thermodynamically more stable O–H species, accompanied by the reduction of W⁶⁺ to W⁵⁺ into the surface layer.

Additionally, Figure 3c shows a typical cross-sectional scanning electron microscopy (SEM) image of the fuel cell with WO₃ acquired after online sintering and test. The SEM image indicates that the WO₃ electrolyte layer appears fully dense, without noticeable connected pores, and appears well-adhered to the anode substrate, without any cracking or delamination after fuel cell testing as compared with before testing (Figure 3d,f). Such adhesive and dense structure of WO₃ electrolyte guarantee better fuel cell performance [23,26].

3.3. Electrochemical Impedance and Electrical Conductivity

Further, EIS characterizations were performed for the cell with synthesized WO₃ electrolyte in both air and H₂/air atmosphere at 450–550 °C under OCV conditions. Figure 3a,b shows Nyquist curve of the measured EIS spectra and Fitted data using ZSIMPWIN software with equivalent circuit modeling of R_o-(R₁-CPE₁) − (R₂-CPE₂), where R_o, is the ohmic resistance from the electrolyte, R₁ and R₂ belong to charge transfer and mass transfer losses of the fuel cell with WO₃ electrolyte, respectively. It can be seen from the EIS spectra of Figure 4a,b that R_o of the fuel cell with WO₃ dramatically reduced in H₂/air (fuel cell operating conditions) as compared to in the air. For example, fuel cell with WO₃ electrolyte exhibited R_o of only 1.09 Ω cm² in H₂/air at 550 °C, while in the air it shows 12.85 Ω cm² Similarly, a clear difference in R_o values at a low operating temperature of 500 and 450 °C also can be visualized. These results describe that when WO₃ is exposed to the H₂ atmosphere, its surface could be reduced to W⁵⁺ to form WO_3−δ/WO_3, and protons can easily be transported through this layer as reported for CeO₂. Afterward, it breaks the accumulation layer at the interface of electrode/electrolyte and hence reduces the charge transfer resistance (R₁), mass transport resistance (R₂) of fuel cells, as shown in Figure 4a,b. The decrease in charge transfer resistance at the interface of electrolyte/electrode for each cell may be followed by value capacitance for that cell (i.e., R₁∼C₁, R₂∼C₂), where capacitance can be determined by $C_i=\left(\frac{R_i{Q}_i^{\frac{1}{n}}}{R_i}\right)$, where R is corresponding resistance and n to Frequency power [0 ≥ n ≤1] of the Q’s values [11]. The cell’s charge transfer resistance in H₂/air compared to EIS spectra in air decreases from 10.5 to 0.17 Ω cm², followed by a decline in space charge capacitance from 1.231 × 10⁻³ F to 1.713 × 10⁻⁵ F at 550 °C [24,25,27].

The ionic conductivity of synthesized WO₃ in both air and H₂/air was calculated using R_o values of the fitted data. As shown in Figure 4c, the ionic conductivity of synthesized WO₃ in H₂/air is much higher than only in air atmospheres. The synthesized WO₃ exhibits the ionic conductivity of 6.12 × 10⁻² S cm⁻¹ in H₂, while in the air, only 1.21 × 10⁻³ S cm⁻¹. The conductivity results support our findings that the WO₃ surface could be reduced to deficient WO_3−δ and support fast protonic transport. Moreover, cell ionic conductivity was measured over 24 h to check its behavior, and it shows the stable ionic conductivity as can be seen in Figure 4d. Furthermore, the proton conductivity of WO₃ was measured using Ag current collectors in H₂ without using the NCAL electrode. In these conditions, WO₃ shows a little lower proton conductivity as compared to fuel cell conditions (Figure 4c). However, these results indicate that WO₃ synthesized by three-step methods could be a good candidate for electrolyte application in advanced SOFCs [20].

3.4. Spectroscopic Analysis

Moreover, different spectroscopic techniques, such as Raman, UV-visible, and X-ray photoelectron spectroscopies, were employed to study further structural properties of WO₃ powders before and after fuel cell measurements. Figure 5a displays specific Raman shifts bands for synthesized WO₃ centered at 266, 322, 709, and 803 cm⁻¹ links to the stretching vibrations of (O-W-O), (O-W-O), (W₂O₆ & W₃O₈), and (W-O-W) from 200–2400 Raman shifts/cm⁻¹. After performing fuel cell measurements, a small redshift of 0.5 cm⁻¹ in the Raman shift band of WO₃ was observed. However, overall, WO₃ shows good structure stability after fuel cell measurements. Moreover, UV-Visible absorbance spectra of WO₃ for before and after fuel cell performance measurements are presented in Figure 5b. There is just a little bit of difference in absorbance spectra that can be observed. The difference in absorbance spectra is an obvious indicator for lower down in the energy band gap of WO₃, which only could be due to the reducing the surface and producing oxygen vacancies. The low energy gap will help to improve the ionic conductivity and, hence, better fuel cell performance [21,24,25,28].

High-resolution XPS spectra of as-synthesized WO₃ powders are shown in Figure 6a,b. After subtracting Shirley’s background, high-resolution XPS spectra were fitted by the mixture function of Lorentzian and Gaussian. Our focus was to observe the chemical and electronic state configuration changes of W-4f and O1s spectra before and after electrochemical performance measurements. Figure 6a shows the XPS spectra of 4f-W⁶⁺ (5/2, 7/2) and 4-f W⁵⁺ (5/2, 7/2) that appear at 35.32/37.52 and 35.7/38.05 eV, whereas in after fuel cell performance measurements at 35.12/37.22 and 35.9/37.85 eV. These downshifts in B.E of W⁶⁺ 4f (5/2, 7/2) and W5 + 4f (5/2, 7/2) manifest a reduction in WO₃ after fuel cell measurements. Moreover, O1s spectra of the material also influence the ionic conductivity of a material [29,30]. The O1s spectra after fuel cell measurements contain lattice oxygen (lattice O²⁻) and oxygen vacancy peaks, as shown in Figure 6b. The O1s spectra of WO₃ after fuel cell measurements display two partially superimposed peaks (Figure 6b). There are two significant excitations: the first includes O1s of lattice oxygen bands and the second, WO_3−δ band with binding energy (BE) ranging from 528 to 533.5 eV [28,29]. The low BE peak at 529.2 can be ascribed to the lattice oxygen (O Lattice), higher at 531.4- to extra oxygen vacancies. The increased area percentage ratio of O_lat/O_vac of WO₃ after fuel cell measurements indicates high oxygen vacancies concentration, which plays an essential role in high fuel cell performance [21,28]. However, these results provide clear evidence that an in situ surface reduction produces oxygen vacancies for the fast protons transports, as shown in Figure 6c. Therefore, our developed strategy could help create high-performance LT-SOFCs electrolyte materials in a new way [31]. The XPS data shows that W4f and O1s percentage is different (W4f−18.37; O1s 56.65%) as compared to as synthesized sample (W4f−20.5; O1s 54.26). It means the after the fuel cell performance measurements the WO₃ has more oxygen vacancies than as pristine WO₃. Furthermore, electron spin resistance (ESR) configures the change in structural properties after fuel cell measurements. Figure 6c shows the full spectrum of proton unirradiated WO₃ and proton irradiated WO3 phases. Where proton irradiation-induced defect phase lines and the well-resolved group of the ESR spectrum can be seen after fuel cell measurements. However, ESR results confirm our proposed mechanism that the surface layer of WO₃ is reduced during the in-situ operation of the fuel cell, and it facilitates proton transport.

The schematic diagram of the process involved for pristine WO₃ to deficient WO_3−δ layer formation for migration of oxygen and proton ions is shown in Figure 7. The proton transport is accompanied by a surface layer of WO₃/WO_3−δ as shown in the last step of Figure 7.

4. Conclusions

In summary, we have successfully synthesized and characterized WO₃ nanostructured material by combining three synthesis methods. Furthermore, the synthesized nanostructured WO₃ with a monoclinic structure demonstrated excellent proton conductivity during fuel cell operation. The ionic conductivity reached 6.12 × 10⁻² S cm⁻¹ in an H₂ atmosphere as compared to 1.21 × 10⁻³ S cm⁻¹ in the air. The fabricated fuel cell using prepared WO₃ as elect4rolyte exhibited a high-power density of 418 mW cm⁻² at 550 °C. Furthermore, we used different microscopic and spectroscopic analyses to study the mechanism behind the drastic increase in ionic conductivity of WO₃ during fuel cell operation. We found that synthesized nanostructured WO₃ surfaces can easily be reduced to form an oxygen-deficient layer and facilitate protons transport effectively when exposed to a reduction in the atmosphere. The ex-situ spectroscopies that included Raman, UV-visible, XPS, and ESR clearly described our findings and the structural change properties of WO₃ during the fuel cell operation. In conclusion, this method could form the basis of interest to develop new WO₃ based proton-conducting electrolytes, which could be useful for all energy devices and material systems.