New Energy And Industrial Technology Development Organization Research Articles

FC-Platform Activities and Post Analysis of Durability Tested MEA

FC-Platform, specialized in PEFC's analysis and evaluation, has been launched as a NEDO's funded project. FC-Platform consists of 5 groups of the electrochemical evaluation, simulation, materials analysis, materials informatics and management, involves several institutions. The numbers of various materials research themes have been also running in the NEDO's project. FC-Platform provides better directions of improving for these materials research through analysis and evaluation of their materilas. Moreover, Materials informatics (MI) have been applyed to the development of thier matrials in several themes. Visualizing and predicting phenomena in the PEFCs are essential for realizing high performance and durability at a material level. Therefore, we are focusing on advanced visualization techniques and trans-scale stack modeling development. We also had post analysis and evaluation of the MEAs under long term road-cycle test.This work was supported by a NEDO FC-Platform project commissioned by th e New Energy and Industrial Technology Development Organization (NEDO).

Self-Supporting Microporous Layer for Polymer Electrolyte Fuel Cells

Introduction Improving power density is desired in research and development of polymer electrolyte fuel cells (PEFC) [1-5]. A reduction of gas diffusion resistance by down-thinning the gas diffusion layer (GDL) less than 50 μm is desirable [2]. Whilst both the electrocatalyst layers and electrolyte membrane are around 10 μm, the GDL is still relatively thick at about 150 μm [1]. The previous study investigated the possibility and technical issues in using high-strength and thin, porous metallic GDLs [3]. Based on such previous studies, this study aims to develop a self-supporting thin MPL/GDL with improved power density by creating conductive paths and reducing contact resistance with microporous layer (MPL) deposited on various GDLs. Experimental In this study, MPL was screen-printed on various mesh GDLs as substrates. Cells were fabricated by using these materials, and current-voltage characteristics and microstructural observations were performed. Stainless steel (SUS316 977mesh, 28 μm thick), titanium (200mesh, 152 μm thick), and carbon fiber (45 μm thick) were used as mesh GDL materials. The MPLs were composed of carbon black, carbon nanotubes, polytetrafluoroethylene (PTFE), and solvent (a mixture of polyethylene glycol 600 and pure water), which were mixed, screen printed, and heat treated at 300°C for 30 min to fabricate the MPL/GDL. Conventional materials such as Pt/C (TEC10E50E, Tanaka Kikinzoku, Japan) were used for preparing membrane-electrode-assemblies (MEAs). In the electrochemical characterization of this study, current-voltage characteristics were evaluated using an electrochemical impedance analyzer (SAS SP-240, Bio-Logic Science Instruments, France), and various overvoltages were separated. Results and discussion Figure 1 shows the current-voltage characteristics of cells using various MPL/GDLs. The thickness of MPL/GDLs fabricated was ranging from 62 to 180 μm. For comparison, the characteristics of a cell using a standard material, commercially available MPL/GDL (22BB, SGL Carbon, Germany), are also shown. The results indicate that the use of any kind of these MPLs improves cell performance. The electrochemical performance with metallic GDLs was considerably improved by applying MPLs. In particular, GDLs with metallic meshes made of SUS316 show a significant improvement in current-voltage characteristics by depositing MPLs. The activation overvoltage and ohmic overvoltage were comparable to those of 22BB, the commercial GDL/MPL. This is because the contact resistance between the GDL and the catalyst layer is reduced by the carbon component in the MPL, and the electrical conductivity is improved. The concentration overvoltage was also suppressed by depositing the MPL. The cell performance with MPL3/SUS316 977mesh was comparable to that of 22BB. In the future, we aim to further reduce the concentration overvoltage by optimizing e.g., the content of hydrophilic components in MPLs. Acknowledgment This paper is based on results obtained from a project, JPNP20003, commissioned by the New Energy and Industrial Technology Development Organization (NEDO). References Takahashi, T. Ikeda, K. Murata, O. Hotaka, S. Hasegawa, Y. Tachikawa, M. Nishihara, J. Matsuda, T. Kitahara,M. Lyth, A. Hayashi, and K. Sasaki, J. Electrochem. Soc., 169, 044523, (2022).New Energy and Industrial Technology Development Organization (NEDO), Roadmap of Fuel Cells for Heavy-Duty Vehicles, 48, (2022), (in Japanese), https://www.nedo.go.jp/content/100944011.pdf Yamamoto, M. Yasutake, Z. Noda, S. M. Lyth, J. Matsuda, M. Nishihara, A. Hayashi, and K. Sasaki, ECS Trans., 109 (9),265 (2022).Larminie and A. Dicks, Fuel Cell Systems Explained, 2nd ed., John Wiley & Sons, England, 2003.Kitahara, T. Konomi, H. Nakajima, and J. Shiraishi, Kikai B, 76 (761), 101, (2010). Figure 1

Degradation Mechanisms of All-Solid-State Lithium-Ion Batteries with Sulfide-Type Electrolytes Based on High-Temperature Studies

Introduction Sulfide-based all-solid-state Li-ion batteries (LIBs) are potential next-generation energy storage devices due to their increased safeties and energy densities compared to those of conventional LIBs with liquid electrolytes. However, our understanding of the durabilities and degradation mechanisms of such batteries remains limited. Via industry-academia collaboration, a project that aims to develop a sulfide-based all-solid-state LIB for use in electric vehicles (NEDO “Development of Fundamental Technologies for All-Solid-State Battery Applied to Electric Vehicles (SOLiD-EV),” 2018-2023) has been undertaken in Japan. Under this project, we conducted high-temperature cycling and storage studies using a SOLiD-EV prototype cell to gain insight into the durability and degradation mechanism of the sulfide-based all-solid-state LIB. [1] Experimental The SOLiD-EV prototype cell, which was a developed laminate-type single-layer all-solid-state LIB with a capacity of approximately 8 mAh, was assembled by combining LiNbO3-coated LiNi0.5Co0.2Mn0.3O2 (NCM, Sumitomo Metal Mining, Tokyo, Japan) as the positive active material, graphite (Mitsubishi Chemical Holdings, Tokyo, Japan) as the negative active material, and argyrodite-structured Li(7–x)PS(6–x)Cl x (x ≈ 1.2 × 10−3 S cm−1; Mitsui Mining & Smelting, Tokyo, Japan) as the solid electrolyte (SE). All assembly processes were conducted in a glove box under an Ar atmosphere with a dew point below −76 °C and an O2 concentration of <1 ppm.In the high-temperature cycling study, the prototype cell was charged at a rate of C/10 to 4.2 V in constant current mode and further charged in constant potential voltage (CC-CV) mode. It was then discharged at a rate of C/3 to reach 3.0 V in constant current (CC) mode at 100 °C. In the high-temperature storage study, the prototype cell was stored in a thermostatic chamber at 60, 80, or 100 °C after CC-CV charging. In both studies, the following performance measurements were conducted at intervals of 7 days: charge/discharge evaluations at the same current condition as that used in the high-temperature cycling study at 25 °C, charge/discharge evaluations at a rate of C/20 at 45 °C, and electrochemical impedance spectroscopy at a depth of discharge (DOD) of 50% at 25 °C. Both studies were conducted for a total of 28 days.X-ray photoelectron spectroscopy (XPS), scanning electron microscopy (SEM), and inductively coupled plasma atomic emission spectrometry (ICP-AES) component analysis of the SE at each electrode were performed using the prototype cells that were degraded in the high-temperature studies as disassembly analyses. Results and discussion Each evaluation reduces the capacity of the SOLiD-EV prototype cell, and the loss of capacity at high temperatures is faster than that at low temperatures, as shown in Figure 1. This reduction in the capacity during the cycling study at 100 °C is lesser than that during the storage study at 100 °C in the first 7 days; however, the reduction during the cycling study increases after 14 days. The degradation of the electrodes was caused by the repeated charging and discharging of the prototype cell. Increases in internal resistance are also confirmed, and the trend of the temperature dependence is almost identical to the capacity trend. The partial substitution of PS4 3− in Li(7–x)PS(6–x)Cl x by P–S–P, –S–S–, and S components, as revealed by XPS, and the decrease in the Li/P ratio, as revealed by ICP-AES, indicate that the SE (Li(7–x)PS(6–x)Cl x ) is oxidatively decomposed at the positive electrode due to its (noble) potential. S, which is a product of oxidative decomposition, is a poor conductor of Li+ ions and may increase the charge transfer and path resistances in the positive electrode. The substitution of PS4 3− in Li(7–x)PS(6–x)Cl x with Li3P and Li2S and the increase in the Li/P ratio indicate that the SE (Li(7–x)PS(6–x)Cl x ) is reductively dissolved at the negative electrode due to the negative electrode (low) potential. The capacity degradation of the cell is due to the loss of active Li+ ions caused by the reduction of the sulfide electrolyte at the negative electrode. Furthermore, the change in the composition of the SE, as determined via ICP-AES (Figure 2) before and after the high-temperature storage study, suggests that the SE degradation reactions at the positive and negative electrodes are temperature-dependent and roughly follow the Arrhenius law, with different activation energies. Acknowledgement This study was supported by the “Development of Fundamental Technologies for All Solid State Battery Applied to Electric Vehicles” project (SOLiD-EV, JPNP18003) commissioned by the New Energy and Industrial Technology Development Organization (NEDO). Reference [1] K. Ando, T. Matsuda, T. Miwa, M. Kawai, D. Imamura, Battery Energy, 20220052 (2023). Figure 1

Effect of Electrolyte Anions on the Activity of Iridium Oxide Catalysts for Water Electrolysis

Developing environmentally friendly and renewable energy sources is essential due to the environmental and energy security issues caused by traditional combustion of fossil fuel-based energy sources. Therefore, as a clean and sustainable energy carrier, hydrogen has attracted great attention in recent years. In this regard, water electrolysis utilizing renewable energy sources to produce hydrogen is considered one of the most promising technologies. It is believed elucidation of the reaction mechanism of highly active iridium oxides is promising for the further development of cost-efficient PEM water electrolysis catalysts. The properties of the MEA could be affected by many parameters such as the catalysts, operating temperature, and pressure1,2. Another important factor is the cation impurities that can be found in the circulating water, which could severely degrade the MEA performance during the operation of a PEM water electrolyze. However, there are less reports on anion impurities such as SO4 2- and PO4 3-. Therefore, in this study, the relationship between electrolyte anions and OER activity was investigated for IrOx catalysts with different crystal structure. By using operando X-ray absorption spectroscopy (XAS), we investigated the electronic state of Ir and O under operation potential.IrOx with different degree of crystallinity (SA3.5 and SA58) were provided by Tanaka Kikinzoku Kogyo K.K., Chemical & Refining Company and named after their BET surface areas. These two samples were then tested for electrochemical properties in different concentrations of phosphoric acid and perchloric acid, respectively. The electrochemical measurements were performed using a standard three-electrode cell connected to an MPG-205-NUC system (Bio-Logic). The cyclic voltammograms (CV) were scanned from 0.1 V to 1.2 V vs. RHE at 50 mV s-1. The electrochemical activities of the catalysts were subsequently investigated using linear sweep voltammetry (LSV) in the range of 1.1 to 1.7 V vs RHE at 10 mV s−1. operando XAS were collected using a home-made flow-type cell, where Pt wire and RHE worked as counter and reference electrodes, respectively. The catalyst ink was coated onto an Nafion membrane to prepare both the X-ray window and working electrode. The electrolyte was circulated at the flow rate of 100 mL min-1.The catalytic performance of IrOx was investigated at different types and content of electrolyte under H3PO4 and HClO4. Figure 1 shows the CVs of SA3.5 and SA58 under different types and content of electrolyte. In the presence of phosphate anions, SA3.5 showed a tendency to decrease Ir redox between 0.6 and 1.1 V with increasing electrolyte concentration. On the other hand, there were no obvious changes in SA58. It was found that IrOx catalysts have an great effect of specific adsorption of anions and low symmetry of monoclinic phase catalysts (SA 3.5) are more susceptible to anion poisoning. Figure 2 shows the result of examining the poisoning mechanism using soft X-ray XAS. O K-edge results present the formation of µ2-O (O-O) bond at around 528.7 eV, which was previously observed by Nong, H. N. et. al. through operando O-K edge measurements during OER process3. The peak intensity at 528.7 eV of the amorphous samples was in the same order as the OER activities. According to potential dependence change of the pre-edge peak, under the presence of phosphate anions, the hole formation to Ir-O hybrid orbitals was suppressed in SA3.5. This result indicates that the anion adsorbs to the active site, poisoning the catalyst surface and inhibiting the reaction. Acknowledgements This work is based on results obtained from a project (JPNP14021) commissioned by the New Energy and Industrial Technology Development Organization (NEDO) of Japan. References Zhang, Y.; Zhu, X.; Zhang, G.; Shi, P.; Wang, A.-L.; et al. J. Mater. Chem. A: 2021, 9, 5890-5914.Chen, Z.; Duan, X.; Wei, W.; Wang, S.; Ni, B.-J.; et al. Nano Energy: 2020, 78, 105270.Nong, H. N.; Falling, L. J.; Bergmann, A.; et al. Nature 2020, 587 (7834), 408-413. Figure 1

Self-Repairing Ability of Hybrid Cobalt Hydroxide Nanosheets As Oxygen Evolution Catalysts in Alkaline Electrolytes

Introduction Water electrolysis is a key technology to produce H2 from renewable energy. Because proton exchange membrane water electrolyzers (PEMWE) require noble metals as catalysts, alkaline-type electrolyzers, such as alkaline water electrolyzers (AWE) and anion exchange membrane water electrolysis (AEMWE), have attracted much attention for the use of nonnoble-metal-based catalysts. One of the most important issues of AWE is the degradation of catalysts due to generation of reverse current upon shutdown and power fluctuation. Reverse current hardly generates in AEMWE in principle; however, it possibly generates if an electrolyte is supplied instead of pure water.We have previously reported the self-repairing anode catalysts for AWE [1, 2]. A hybrid cobalt hydroxide nanosheet (denoted as Co-ns, Figure 1a), a Brucite-type cobalt hydroxide modified with tripodal ligands (tris(hydroxymethyl)aminomethane), was dispersed in an alkaline electrolyte. Co-ns is deposited on the anode during the electrolysis at 800 mA/cm2 to form a catalyst layer. After the degradation of the catalyst by potential cycling, Co-ns is deposited to retain the catalytic activity. To develop catalysts with higher activity, durability, and repairing ability, the understanding of the electrochemical deposition of Co-ns is quite important. Here, we demonstrate the mechanism of the electrochemical deposition of Co-ns.ExperimentalCo-ns was synthesized according to our previous report [1]. The Co-ns was dispersed in 1 M KOH electrolyte at varied concentrations of Co-ns. Ni electrodes were used as both working and counter electrodes. Reversible hydrogen electrode (RHE) was used as a reference. Constant current electrolysis at 800 mA/cm2 was applied for 30 min, followed by cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) to evaluate the oxygen evolution reaction (OER) activity and the amount of deposited Co. These processes ware repeated for several times to characterize the deposition behaviors. Similar experiments were performed, using Co(OH)2 or CoOOH instead of Co-ns.Results and DiscussionCo-ns was transformed into CoOOH by the decomposition of the tripodal ligand. The deposited amount of Co was evaluated by CV, in which the charge of the cathodic currents at 0.8–1.6 V vs. RHE (denoted as Q c) was proportional to the amount of Co measured by ICP-AES. Figure 1b shows the plots of Q c vs. duration of electrolysis. The amounts of deposited CoOOH depended on the concentration of Co-ns. The deposition rates (slope of the curves) gradually decreased along with the progress of the catalyst deposition. During the electrolysis, the electrode potential decreased because of the increased OER activity due to the catalyst deposition. Assuming that the deposition of CoOOH is a pseudo first-order reaction, its rate constant was calculated by dividing the deposition rate by the concentration of Co-ns. The rate constants are linearly correlated with the electrode potential (Figure 1c); thus, the deposition was confirmed to be an electrochemical process. The present results mean that the electrochemical deposition of CoOOH is accelerated when the electrode potential increases due to the degradation of catalysts. Therefore, the self-repairing process is automatically controlled by the condition of the electrode.The electrochemical deposition proceeds with Co(OH)2, whereas CoOOH hardly deposited at the same conditions, indicating that the oxidation of Co2+ to Co3+ is an essential step of the electrochemical deposition. The electrochemical deposition of Co(OH)2 was stopped after 60 min of electrolysis because of the autooxidation of Co(OH)2 by dissolving O2. It was shown that Co-ns is stable in an electrolyte saturated with O2 probably because the organic modification suppresses the autooxidation. Thus, it is shown that the organic modification of cobalt hydroxide is crucial to retain the self-repairing ability in an electrolyzer.ConclusionIn conclusion, the potential dependence of the rate constant and the essential step of electrochemical deposition of Co-ns were elucidated through systematic studies using Co-ns and related metal hydroxide catalysts. These findings lead to the design of self-repairing electrocatalysts with high activity, long lifetime, and high self-repairing ability using organic functionalization.Acknowledgement This work was supported by the JSPS KAKENHI (grant number 20H02821). Part of this study uses outcomes of the development of fundamental technology for the advancement of water electrolysis hydrogen production in the advancement of hydrogen technologies and utilization projects (grant number JPNP14021) commissioned by the New Energy and Industrial Technology Development Organization (NEDO) in Japan.

Quantification Method of Mass Transfer Resistance in Cathode Catalyst Layer in PEFC

IntroductionTo spread PEFCs, it is important to optimize the cathode catalyst layer (CCL) because the oxygen reduction reaction (ORR) in CCL is sluggish. Our previous study proposed that the mass transfer resistance in the through-plane direction in CCL can be quantified by measuring ORR rate dependency on oxygen partial pressure, effective oxygen diffusion coefficient, ORR rate constant, or CCL thickness[1]. In this study, dependency of the polarization curve on the oxygen partial pressure was measured with commercial supported-platinum catalysts at fixed RH and temperature. The objective of this study is to establish the quantification methods of mass transfer resistances in CCL.TheoryIn our previous study[2, 3], the current density, i is expressed as i = 4F δ (C) k vcm p Oc F e [A/m2], (1)where δ (C) is the cathode catalyst layer (CCL) thickness, k vc is the reaction rate constant per unit volume of CCL [mol/(Pa·m3·s)] which does not include through-plane mass transport resistance. p O is the oxygen partial pressure [Pa]. The subscripts m and c denote PEM–CCL boundary and CCL–gas diffusion layer boundary, respectively. F e is the effectiveness factor, a function of the following 4 dimensionless moduli, M Om , M pm , P Om , and y Oc . F e is defined as the ratio of the observed reaction rate to the reaction rate which does not have the influence of through-plane mass transfer resistance. M Om denotes the ratio of O2 diffusion resistance to the reaction resistance. M pm denotes the ratio of proton transfer resistance to the reaction resistance. P Om is defined as the ratio of O2 diffusion resistance to the convection resistance. y Oc denotes the oxygen mole fraction. Definitions of the dimensionless four moduli are expressed as M Om =δ (C) (k vcm RT/D eO)0.5, (2) M pm = δ (C) (4F k vcm RT/(σ ep b c))0.5, (3) P Om = δ (C) N gm /(C g D eO), (4) y Oc = p Oc /P, (5)where R, T, σ ep, N gm and P denotes the gas constant [J/(mol K)], cell temperature [K], effective proton conductivity [S/m], total gas flux [mol/(m2 s)] and total pressure [Pa], respectively. When the cathode emf at the PEM–CCL boundary, E cm , T, and relative humidity (RH) are fixed, M Om is fixed. F e and M pm are proportional to i/p Oc and p Oc 0.5, respectively. The relationship between F e and M pm shown in Fig. 1 corresponds to that between i/p Oc and p Oc 0.5. Thus, F e and M pm can be determined just by fitting experimental data of i/p Oc plotted against p Oc 0.5 to the theoretical curves.Results and DiscussionPolarization curves shown in Fig. 2 were measured at 80 °C and 0.77 of inlet RH at varied p Oc . As expected from Eq. (1), current densities at a fixed E cm increases when p Oc increases. The current densities at the fixed E cm are plotted against p Oc as shown in Fig. 3. The current density is sublinear against p Oc . This is because M pm expressed in Eq. (3) increased when p Oc is higher. The ORR rate is lower than the value expected from the Tafel equation when mass transfer resistance increases.The dimensionless moduli were determined by fitting the theoretical relationship to the measured i/p Oc vs. p Oc 0.5 graph. From the determined dimensionless moduli, the ORR rate and mass transport parameters, k vc⊕, D eO and σ ep were also obtained and these values are shown in Table 1 with values in literature[4, 5].Effective proton conductivity, σ ep is relatively low compared with values in literature. This resulted from an experimental error in the current density at p Oc of 16 kPa. Lowering the current density at high p Oc leads to decrease in σ ep because M pm increases when p Oc is raised. By extensively measuring the polarization curves dependency on p Oc , the ORR rate parameters and mass transport parameters can be determined more accurately. Recalculated results using the obtained mass transport parameters are shown in Figs. 3 and 4. The recalculated results appropriately reproduced the ORR rate dependency on E cm and p Oc .ConclusionsThe through-plane mass transfer resistance in CCL was quantified by measuring the ORR rate dependency on the oxygen partial pressure. From the determined dimensionless moduli, mass transport parameters were obtained. Recalculated results appropriately reproduced the ORR rate dependency on cathode emf and oxygen partial pressure.AcknowledgementThis work was supported by the FC-Platform Program: Development of design-for-purpose numerical simulators for attaining long life and high performance project (FY 2020–2023) conducted by the New Energy and Industrial Technology Development Organization (NEDO), Japan. Figure 1

Enhanced Water Management in Fuel Cell Cathodes by the Application of Superhydrophobic Fluorinated Carbon Microporous Layers

Under high current density operation, the efficiency of polymer electrolyte fuel cells (PEFCs) can dramatically decrease. This is due to water accumulation at the cathode side, preventing oxygen diffusion to the electrocatalyst. As such, effective water management is of vital importance by use of a suitable gas diffusion layer (GDL) and/or microporous layer (MPL). MPLs generally consist of carbon black as the porous electron conducting phase, and polytetrafluoroethylene (PTFE) as a hydrophobic binder. Here, we instead use superhydrophobic fluorinated carbon powder in the MPL as a novel material to decrease the required PTFE content. It is confirmed that the water contact angle of the MPL can be increased from 131° to 151° by using fluorinated carbon. Moreover, the fluorinated carbon MPL shows lower oxygen transport resistance at high humidity. Furthermore, in single fuel cell tests at various temperatures and relative humidity values, the I–V performance is significantly and consistently better than for the conventional MPL. These results confirm that fluorinated carbon is a promising new material for water management in the MPLs of PEFCs.This paper is based on results obtained from project JPNP20003, commissioned by the New Energy and Industrial Technology Development Organization (NEDO).1. Tailored Wettability in Fluorinated Carbon Nanoparticles Synthesized from Fluorotelomer Alcohols, E. Can, M. Nishihara, K. Sasaki, S. M. Lyth, Applied Surface Science, 157136 (2023)2. Superhydrophobic Fluorinated Carbon Powders for Prevention of Flooding in PEFCs, E. M. Can, A. Mufundirwa, T. Kitahara, M. Nishihara, K. Sasaki, S. M. Lyth, Journal of Power Sources, 548, 232098 (2022)3. Solvothermal Synthesis of Superhydrophobic Hollow Carbon Nanoparticles from Fluorinated Alcohol, S. M. Lyth, W. Ma, J. Liu, T. Daio, K. Sasaki, A. Takahara, and B. Ameduri, Nanoscale, 7, 16087 (2015)

Highly Durable Membrane–Electrode-Assemblies Using High M w and Ether-Free Polyfluorene-Based Electrolytes for Anion Exchange Membrane Water Electrolyzer

Water electrolysis technology, which produces hydrogen through electrolysis of water using electricity derived from renewable energy, is attracting attention as a way to realize a carbon-neutral society. Among water electrolysis systems, anion exchange membrane water electrolysis (AEMWE) can achieve high energy conversion efficiency and quick response to fluctuating power supplies by using thin polymer electrolyte membranes, and operates in an alkaline environment, allowing the use of inexpensive non-precious metals for bipolar plates and electrodes. However, conventional AEMs and ionomers exhibit low durability in high-temperature and alkaline environments, which is a critical issue in realizing high-performance and highly durable membrane–electrode assemblies (MEAs) for AEMWE.Our group developed high-molecular-weight (M w) and ether-free, polyfluorene-based anion conducting polyelectrolyte (PFT-Cx-TMA, Cx represent the number of carbon atoms in the alkyl side chain) as shown in Fig. 1a.1 These polyelectrolytes show high OH− conductivity over 100 mS cm− 1, high chemical durability (stable in 8 M NaOH at 80 °C for 168 h), and high tensile strength (25–42 MPa, comparable to a Nafion membrane). This study applied PFT-Cx-TMA to AEMs and ionomers in MEAs to evaluate the durability against various AEMWE operations (constant current operation and start-stop operation in different voltage ranges), demonstrating MEAs with both high performance and high durability.Fig. 1b shows the constant current durability of the MEA with PFT-C8-TMA as AEM and ionomer.2 This test was performed at the constant current density of 0.2 A cm− 2 and 80 °C using a 1 M KOH aqueous solution fed to the anode. Here, to investigate the effects of AEMs and ionomers on MEA durability, conventional precious metal catalysts (anode: IrO2, cathode: Pt/C or PtRu/C) were used in the MEA. The MEA with PFT-C8-TMA exhibited a current density of 1.0 A cm− 2 at 1.79 V and there was almost no voltage change even after the constant current operation for over 100 h. It is noteworthy that the MEAs using PFT-C8-TMA and PFT-C10-TMA demonstrated high performance and high durability not only in the alkaline-fed operation but also in the pure-water-fed operation.Next, this study investigated the start-stop durability of MEAs using voltage cycles simulating the start-stop operation of water electrolyzer. Here, square-wave voltage cycling of holding at 1.8 V and 0.1 V for 1 min each was used for the start-stop test at 80 °C with a 1 M KOH solution fed to the anode. As a comparison, the MEA was prepared using a commercial AEM (Sustainion X37-50, 50 μm in thickness) and ionomer (Sustaionion XB-7). As shown in Fig. 1c, the MEA using a PFT-C10-TMA AEM (56 μm in thickness) and ionomer was highly durable in the start-stop test compared to the MEA using commercial Sustaionion. During the test, the elution of black solid particles from the catalyst layer and the increase in charge transfer resistance (Rct) were observed in the MEA with Sustaionion, whereas no elution of catalyst materials and no change in Rct were confirmed in the MEA with PFT-C10-TMA. These results suggest that nonlinear mechanical stress via rapid bubble formation and shutdown during start-stop cycles caused the leaching of catalysts, resulting in the performance loss of the MEA with Sustainion. The commercial ionomer (Sustainion XB-7) exhibits a large dimensional change (156%) compared to the PFT-C10-TMA membrane (34%) at 80 °C in a 1 M KOH aqueous solution. Catalyst particles covered with a highly swollen ionomer are easily detached from a catalyst layer by the nonlinear mechanical stress. On the other hand, the PFT-C10-TMA ionomer has a high molecular weight (M w = 240000 g mol−1), low swelling degree, and well-entangled polymer chains, hence preventing leaching of the catalyst particles.In summary, this study demonstrates that MEAs using high M w and ether-free polyfluorene-based electrolytes with both high chemical and mechanical durability have excellent durability for AEMWE operations including constant and dynamic (start-stop voltage cycling) operations. Acknowledgement This presentation is based on results obtained from a project, JPNP14021, commissioned by the New Energy and Industrial Technology Development Organization (NEDO). References (1) S. Miyanishi and T. Yamaguchi, Polym. Chem., 2020, 11, 3812–3820.(2) R. Soni, S. Miyanishi, H. Kuroki, and T. Yamaguchi, ACS Appl. Energy Mater., 2021, 4, 1053–1058. Figure 1

Highly Durable and Active Electrocatalysts Using Pt Nanorod Catalysts Supported on Nb Doped SnO2 for Polymer Electrolyte Fuel Cells Operating at Wide Temperature Range

The improvement of the durability and activity in cathode catalyst is required toward the wide spread use of polymer electrolyte fuel cells. One of the interesting catalysts is the Pt catalysts supported on SnO2 (Pt/M-SnO2, M = Nb, Ta) without carbon additives [1,2]. The durability (startup/shutdown, load cycling) and oxygen reduction reaction (ORR) activity of the Pt/M-SnO2 is superior to those of commercial Pt catalysts supported on carbon black (Pt/CB) etc. [3-8]. The non-carbon support of M-SnO2 nanoparticles has a unique carbon-like microstructure of a fused-aggregate network structure, which supply the essential function of the electronically constructing pathways via necking of each support particles and the gas diffusion pathways via the open pores surrounded by particles. The Pt nanorod catalyst on M-SnO2 also improves the ORR activity as shown in the catalyst design concept (Fig. 1)[9]. The IV performance of the single cell using the Pt nanorod/Nb-SnO2 cathode catalyst layers at operating temperatures from 80oC to 120oC is quite high to approach the NEDO target performance (Fig. 2) with keeping high durability. The catalyst is considered to have possibilities for application of the heavy-duty vehicles operating at wide-temperature range (< 120oC)[10]. Acknowledgments This work was partially supported by funds for the project “Electrolytes, Catalysts and Catalyst layers with Extraordinary Efficiency, power and Durability for PEFCs-2030 (ECCEED’30) from the New Energy and Industrial Technology Development Organization (NEDO) of Japan, and JSPS KAKENHI Grant Number (23H02059) from the Ministry of Education, Culture, Sports, Science and Technology. References A. Masao, S. Noda, F. Takasaki, K. Ito, K. Sasaki, Electrochem. Solid-State Lett.,12, B119 (2009).E. Fabbri, A. Rabis, R. Kötz, T.J. Schmidt, Phys. Chem. Chem. Phys., 16, 13672 (2014).K. Kakinuma, M. Uchida, T. Kamino, H. Uchida, M. Watanabe, Electrochim. Acta, 56, 2881 (2011).Y. Senoo, K. Kakinuma, M. Uchida, H. Uchida, S. Deki, M. Watanabe, RSC Adv., 6, 321800 (2014).Y. Chino, K. Taniguchi, Y. Senoo, K. Kakinuma, M. Watanabe, M. Uchida, J. Electrochem. Soc., 162, F736 (2015).K. Kakinuma, R. Kobayashi, A. Iiyama, M. Uchida, J. Electrochem. Soc., 165, J3083 (2018).K. Kakinuma, K. Suda, R. Kobayashi, T. Tano, C. Arata, I. Amemiya, S. Watanabe, M. Matsumoto, H. Imai, A. Iiyama, M. Uchida, ACS Appl. Mater. Interfaces 11, 34957 (2019).K. Kakinuma, M. Hayashi, T. Hashimoto, A. Iiyama, M. Uchida, ACS Appl. Energy Mater., 3, 6922 (2020).G. Shi, T. Tano, D.A. Tryk, A. Iiyama, M. Uchida, K.Kakinuma, ACS Catal., 11, 5222 (2021).K. Kakinuma, H. Taniguchi, T. Asakawa, T. Miyao, M. Uchida, Y. Aoki, T. Akiyama, A. Masuda, N. Sato, A. Iiyama J. Electrochem. Soc. 169 (2022) 044522. Figure 1

Improvement of Current Efficiency of PEM Electrolyzer for Organic Chemical Hydride Energy Carrier Production with Water Electrolysis by Humidified Air Supply to Anode

Introduciton To improve hydrogen energy carrier synthesis, we have studied toluene electrohydrogenation and water decomposition using proton exchange membrane (PEM) electrolysis. The cathode reaction is electrohydrogenation of toluene to synthesis methylcyclohexane instead of hydrogen evolution reaction of water electrolysis. Only side reaction of the cathode is hydrogen evolution reaction, which was affected by electroosmotic water from the anode to cathode. Sulfuric acid solution feeds as anolyte that is not only reactant reduces water transportation and increases current efficiency [1, 2]. As practical system, water transportation from anode to cathode should be suppressed without sulfuric acid utilizationIn this study, to improve current efficiency by water management, we have investigated humidified air feed to reduce water transfer to cathode compare with liquid water feed. Experimental A single cell electrolyzer made of titanium with 16 cm2 of projected electrode area was used. Figure 1 shows the schematic drawing of toluene direct hydrogenation electrolyzers with humidified air feed. Membrane-electrodes assembly was consisted of Nafion® 117 (DuPont), 0.6 mg cm-2 of PtRu/C (TEC61E54, TKK, I/C = 0.8) cathode catalyst layer and 2.0 mg cm-2 of IrOx (TEC77100, TKK, I/IrOx = 0.28). Cathode backing was a carbon paper (10BC, SGL) of 0.02 mg cm-2 of Pt loading, which works flow filed of toluene-methylcyclohexane, too. Anode baking as a Pt plated sintered Ti web (Bekaert), and anode flow filed was serpentine type. 10 mL min-1 of toluene or 10% toluene – methylcyclohexane was fed to cathode. 5.0 L min-1 (313 mL min-1 cm-2) of fully humidified air or liquid water was fed to cathode. Operation temperature was 80oC. Internal resistance was determined by high frequency intercept of real axis of the electrochemical impedance spectroscopy in the current density region between 0.1 to 0.6 A cm-2. The water transportation through the membrane was determined by separated water in outlet liquid from cathode during constant current electrolysis, and the Faraday efficiency was determined by generated gas volume from the cathode, simultaneously. Results and discussion Figure 2 shows the cell voltage and current efficiency as a function of the current density for 10% toluene-methylcyclohexane or 100% toluene to cathode and 5.0 L min-1 of fully humidified air or liquid water to anode feed. Cell voltages were almost same for all experiments in less than 0.3 Acm-2 of current density region, while cell voltages of humidified air feed were higher than that of liquid water feed in higher current density region. Here, internal resistance of humidified air and liquid water feed were 0.29 and 0.21 Ω cm2. Therefore, membrane condition of the humidified air feed was drier than that of the liquid water feed. Water transportation through the membrane around 0.2 mg min-1 cm-2, which was less than 1/10 of the liquid water or 1 M H2SO4 feed electrolysis. The current efficiency for 100% toluene feed was more than 95% at 1 A cm-2, and slightly decreased with current density. In this case, the current efficiency of the liquid water feed was a little higher than that of the humidified air feed. The current efficiency of the humidified air feed was more than 95% up to 0.5 Acm-2, while the current efficiency of the liquid water feed decreased 95% around 0.2 Acm-2. In addition, the current efficiency of 10% toluene-methylcyclohexane significantly decrease with current density.The current efficiency would be affected by the cathode potential and toluene transportation in the catalyst layer. For the 100% toluene feed, toluene transportation would less affect and cathode potential would much affect to the current efficiency. Larger internal resistance of the humidified air feed would correspond to dry condition of cathode ionomer, and it would decrease the cathode potential. Therefore, the current efficiency of humidified air was slightly lower than liquid water feed. On the other hand, the current efficiency of 10% toluene feed would be controlled by toluene transportation in the cathode catalyst layer and is significantly affected by the water transportation through the membrane. Therefore, the current efficiency of the humidified air feed significantly higher than that of the liquid water feed. Acknowledgement This study was based on results obtained from the Development of Fundamental Technology for Advancement of Water Electrolysis Hydrogen Production in Advancement of Hydrogen Technologies and Utilization Project (P14021) commissioned by the New Energy and Industrial Technology Development Organization (NEDO). References Nagasawa, K. Tanimoto, J. Koike, K. Ikegami, S. Mitsushima, J Power Sources, 439, 227070 (2019).Oi, K. Nagasawa, T. Takamura, Y. Misu, K. Matsuoka, S. Mitsushima, ECS Meeting s, MA2021-02, 1737 (2021). Figure 1

Utilization of Pt Nanoparticle Catalysts during Capillary Condensation Phenomenon in PEMFC: A Lattice DFT Simulation Based on 3D Reconstruction of TEM Images

The use of porous carbon support particles in Cathode Catalyst Layer (CCL) of Polymer Electrolyte Membrane Fuel Cell (PEMFC) has gained significant interest due to their ability to enable proper dispersion of Platinum (Pt) nanoparticles [1]. This offers a key advantage of avoiding the Pt nanoparticles from potential poisoning by sulfonate anion functional groups in the ionomer that is deposited on the surface of the support particles [2]. Furthermore, the presence of condensed liquid water is essential for effective utilization of the Pt nanoparticles that are dispersed both inside and outside of the porous carbon support particles. This is because water plays a crucial role as a proton pathway for hydrogen ions during the electrochemical reaction, thus a requisite amount of condensed liquid water is also needed to promote efficient catalytic activity and maximize the performance of Pt nanoparticle catalysts. It has been demonstrated that the microstructure characteristics of the porous carbon support particles have a significant impact on the performance of Pt utilization, which is attributed to the varying behavior of water condensation [2]. Microstructure properties such as pore size, surface area, and wettability properties, can influence the distribution of condensed liquid water, thereby affecting the electrochemical performance of Pt nanoparticles in the catalyst layer of PEMFC. Understanding the effect of capillary water condensation in the porous carbon support particles is therefore crucial for maximizing the utilization and catalytic activity of Pt nanoparticles in fuel cell systems.To investigate the behavior of water condensation in the CCL particles of PEMFC, we conducted simulations based on reconstructed structural data obtained from Transmission Electron Microscopy (TEM) images, as shown in Figure 1. The CCL particle used in this study is taken from the Toyota MIRAI Gen 2 model, an emerging commercial PEMFC-based vehicle in Japan. Here, we used the lattice Density Functional Theory (DFT) to simulate the process of capillary water condensation within the pore network of the particle. The lattice DFT method, which is based on the mean-field lattice gas model, can be interpreted as a coarse-grained equivalent of the conventional DFT [3]. The significant advantage of using lattice DFT is to provide a simple model and computationally efficient approach for studying the behavior of capillary water condensation. To validate the parameters used in our lattice DFT simulations, we compared the isotherm curve obtained from our simulation results with experimental data. Through this comparison, we were able to fine-tune the parameters of the lattice DFT simulation to fit the simulated isotherm curve with that of the experiment. When the parameters are optimized, the resulting water distribution obtained from the lattice DFT simulation, as illustrated in Figure 1, serves as a crucial input to investigate the utilization/activation of Pt nanoparticles within and outside the carbon support particle under varying relative humidity (RH) conditions. Since the resulting water distribution depends on RH, we expect that the utilization/activation of Pt nanoparticles also depend on RH. In this study, we defined utilization/activation of Pt nanoparticle when there is a continuous path of condensed water connecting the Pt nanoparticle to the surface of the carbon support particle. This criterion is similar to the case when we assume that the ionomer is distributed uniformly on the surface of the carbon support particle. Our investigation shows that the utilization/activation behavior of Pt nanoparticles follows closely the isotherm curve of the capillary water condensation.The complex internal structure of CCL particles in PEMFC makes it challenging to visualize and understand their workings through experiments alone. Therefore, our approach of using lattice DFT simulations offers valuable insights to the role of capillary water condensation behavior in the utilization/activation of Pt nanoparticles. This approach may also pave the way for future research in optimization strategies for PEMFC, ultimately contributing to the advancement of clean energy technologies.This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) of Japan (Grant number JPNP20003).[1] Islam, M.N., et al. (2022), Nat. Commun. 13, 6157.[2] Ramaswamy, N., et al. (2020) J. Electrochem. Soc. 167 064515.[3] Monson, P. (2012) Micropor. Mesopor. Mater. 160, 47–66.Figure 1. Reconstructed structural data obtained from TEM images of a PEMFC CCL particle used in Toyota MIRAI Gen 2 model, with visualization of condensed water obtained from a sample lattice DFT simulation. Figure 1

Structural Effects on ORR Performances of Carbon-Free Connected Platinum–Cobalt Catalysts with Chemically Ordered Superlattice Structures

Polymer electrolyte fuel cells (PEFCs) have been considered as the promising next-generation energy-conversion technology, owing to their high energy efficiency and clean emissions. However, a sluggish oxygen reduction reaction (ORR) rate and the low durability of conventional catalysts in PEFCs limit the wider commercialization. To address these issues, our group has developed a carbon-free, connected Pt1–Fe1 catalyst with a porous hollow capsule structure. This catalyst comprises the nano-sized network formed by the connection of Pt1–Fe1 nanoparticles and exhibits an enhanced ORR specific activity (9 times higher than a commercial Pt/C) as well as excellent durability against start-stop operations due to the carbon-free structure.[1,2] However, during the load cycle operation, iron is dissolved from the catalyst, leading to the loss of ORR activity and the degradation of polymer electrolytes caused by the Fenton reaction.In this work, carbon-free and iron-free connected Ptx-Co1 catalysts with chemically ordered superlattice structures have been developed for the first time, as shown in Fig. 1(a). In addition, the metal compositions and chemically ordered (superlattice) degrees in the catalysts were controlled to investigate the structural effects on ORR activity and durability of the connected Ptx-Co1 catalysts.Herein, we employed the silica-coating method for the catalyst synthesis, which involves the combination of silica coating and high-temperature annealing. As shown in Fig. 1(b), the surface of the Ptx-Co1 nanoparticles/SiO2 sample was coated with thin silica layers to prevent the detachment of Ptx–Co1 nanoparticles from the SiO2 template as well as the large catalyst agglomeration and coalescence during the annealing process. We controlled the chemically ordered degrees of the connected Pt1–Co1 and Pt3–Co1 catalysts by using different annealing temperatures. The XRD patterns of the catalysts showed the peaks corresponding to L10 type and L12 type chemically ordered structures for the connected Pt1–Co1 and Pt3–Co1 catalysts annealed at 600°C, respectively. These two catalysts possessed ordered degrees (S) of ca. 50–60%. The TEM images of the connected Ptx–Co1 catalysts showed the formation of connected Ptx–Co1 nanonetwork and porous hollow capsule structures.The electrochemical measurements of ORR activity and durability were conducted in a 0.1 M HClO4 electrolyte solution. The connected Pt1–Co1 and Pt3–Co1 catalysts achieved 6- and 10-times higher ORR specific activities than that of Pt/C, respectively. Subsequently, the ORR activity and load cycle durability (0.6 V for 3 s Û 1.0 V for 3 s at 60 °C) of the connected Pt3–Co1 catalyst with different ordered degrees (S = 0%, 28%, and 58%) were evaluated. As the ordered degrees in the connected Pt3–Co1 catalysts were higher, the ORR specific activity was increased. Also, the retention of the specific activity of the ordered catalyst (S = 58%) was higher than that of the disordered catalyst (S = 0%) after 10,000 load cycles. The STEM-EDX line mapping results showed that 83% of the alloyed Co retained in the catalyst with a higher ordered degree even after 10000 load cycles, whereas the disordered and lower ordered catalysts retained lower Co ratios after the test, indicating the suppression effect of an ordered structure on metal dissolution. In addition, the L12 ordered Pt3–Co1 structure remained in the connected catalyst (S = 58%) after load cycle test, confirmed by the TEM diffraction analysis.In summary, we successfully developed the connected Ptx-Co1 catalysts with chemically ordered superlattice structures. In addition, we clarified the structural effects of metal composition and ordered degree in connected Ptx-Co1 catalysts on ORR activity and load cycle durability. These results provide useful guidelines for the design of advanced ORR catalysts towards high-performance PEFCs. Acknowledgement The part of this presentation is based on results obtained from a project, JPNP20003, commissioned by the New Energy and Industrial Technology Development Organization (NEDO).

A Study on Electrochemical Properties of Fuel-Electrode-Supported Reversible Solid Oxide Cells

Introduction Reversible solid oxide cells (r-SOCs) enable both power generation and steam electrolysis, and have attracted attention as a solid state electrochemical energy device for a decarbonized society (1). R-SOCs have the advantage of high efficiency due to high temperature operation, but there exist various technical issues remained in materials selection, start-up and shutdown, and long-term durability. Therefore, reducing operating temperature is desirable for practical applications (2,3). Electrode-supported cells, in which the electrode acts as a structural support, enable lower temperature operation by reducing the thickness of the electrolyte, which has certain electrical resistivity. Here in this study, the current-voltage characteristics and reversible cycle durability under r-SOC operating conditions are evaluated using fuel-electrode-supported cells, with the aim of developing r-SOCs capable of highly efficient fuel cell power generation and steam electrolysis. Experimental For the experiments, fuel-electrode-supported cells were fabricated using fuel electrode-supported half-cells (Japan Fine Ceramics, Japan), schematically shown in Fig. 1. The half-cell consists of scandia-stabilized zirconia (ScSZ: 10 mol%Sc2O3-1mol%CeO2-89 mol%ZrO2) or yttria-stabilized zirconia (YSZ: 8 mol%Y2O3-92 mol%ZrO2) electrolyte and a Ni-cermet fuel electrode, such as Ni-ScSZ or Ni-YSZ. (La0.6Sr0.4)(Co0.2Fe0.8)O3 (LSCF) was applied for the air electrode, and Gd0.1Ce0.9O2 (GDC) was inserted between the electrolyte and the air electrode to suppress interdiffusion and chemical reactions. Four types of cells were used with different combinations of the electrolyte component (YSZ or ScSZ) in the electrolyte layer or the supporting fuel electrode.In electrochemical tests, 50%-humidified hydrogen (100 ml min-1) was supplied to the fuel electrode, and air (150 ml min-1) was supplied to the air electrode. R-SOC initial performance tests were conducted using an electrochemical analyzer (1255B, Solartron). Cell voltage and impedance were measured at 700-800°C at current densities ranging from -0.5 A cm-2 to +0.5 A cm-2. Positive current density means the value in SOFC mode, while negative current density means the value in SOEC mode. In r-SOC 1,000 cycle durability tests, cycles of switching between SOFC and SOEC operation were repeated 1,000 times by varying current density. The range of current density in the cycle tests was between -0.2 A cm-2 and +0.2 A cm-2. The cell impedance was measured before the cycling test and every 100 cycles. After each test, the cells were analyzed by using a focused-ion beam scanning electron microscopy (FIB-SEM) to observe and evaluate the electrode microstructure. Results and discussion Figure 2 (a) shows the r-SOC initial performance of each fuel-electrode-supported cell. The cell with the Ni-YSZ fuel electrode and the YSZ electrolyte showed better current voltage characteristics than other cells. Figure 2 (b) shows the r-SOC 1000-cycle durability of the cell with the Ni-YSZ fuel electrode and the YSZ electrolyte. The results showed a decrease in power generation and electrolysis performance with increasing the number of cycles in both SOFC and SOEC modes. Possible degradation mechanisms will be discussed. Acknowledgments This paper is based on results obtained from a project (Research and Development Program for Promoting Innovative Clean Energy Technologies Through International Collaboration), JPNP20005, commissioned by the New Energy and Industrial Technology Development Organization (NEDO). Collaborative support by Prof. H. L. Tuller and Prof. B. Yildiz at Massachusetts Institute of Technology (MIT) is gratefully acknowledged. References (1) Venkataraman, M. Pérez-Fortes, L. Wang, Y. S. Hajimolana, C. Boigues-Muñoz, A. Agostini, S. J. Mcphail, F. Mar échal, J. V. Herle, and P. V. Aravind, J. Energy Storage, 24, 100782 (2019).(2) Subotić, S. Pofahl, V. Lawlor, N. H. Menzler, T. Thaller, and C. Hochenauer, Energy Proc., 158, 2329 (2019).(3) B. Mogensen, Current Opinion Electrochem., 21, 265 (2020). Figure 1

Influence of Ethanol Decomposition on Dispersion of PEFC Catalyst Ink

To achieve high power density operation of polymer electrolyte fuel cells (PEFCs), it is required to realize high-performance catalyst layer with low oxygen transport resistance, high proton and electron conductivities, and high electrochemical surface area (ECSA) with low Platinum loading. Because dispersion structure of catalyst ink strongly affects porous structure of the catalyst layer, it is crucial for realization of high-performance catalyst layer to understand the dispersion mechanism of the catalyst ink. Our group has reported that decomposition of ethanol, as a solvent, by platinum catalyst significantly affects aggregation of the catalyst ink. [1, 2] But, the mechanism of the catalyst ink aggregation was not fully understood. In this study, effect of ethanol decomposition on the aggregation of the catalyst ink was investigated.The catalyst ink was fabricated by mixing platinum-supported carbon (TEC10V30E, Tanaka Kikinzoku), ionomer (DE1021, Sigma-Aldrich), and water/ethanol solvent (water/ethanol: 60/40 wt%). Decomposition product of ethanol within the catalyst ink solvent was analyzed by GC/MS (GCMS-QP2020NX, SHIMADZU), and presence of acetaldehyde (C2H4O) and acetic acid (CH3COOH) was detected. In order to investigate the influence of these products on aggregation of the catalyst ink, acetaldehyde or acetic acid was added to the catalyst ink. The particle size distribution was evaluated by using laser diffraction type particle size distribution meter (LA-960V2, HORIBA) without dilution, and it was confirmed that the acetaldehyde-added catalyst ink showed larger particle size and the acetaldehyde caused the aggregation of the catalyst ink (Figure 1). In addition, dispersion of the catalyst ink was observed by using an optical microscopy, and aggregation of the catalyst ink by adding the acetaldehyde was clearly observed (Figure 2). To understand the aggregation mechanism by adding acetaldehyde, ionomer adsorption fraction on the platinum-supported carbon was measured by using centrifugation method, and the decrease in ionomer adsorption fraction by adding acetaldehyde was confirmed. From these results, it was confirmed that acetaldehyde generated by the decomposition of ethanol in the catalyst ink leads to a decrease in the ionomer adsorption fraction. It is reported that the ionomer promotes the dispersion of platinum-supported carbon due to electrostatic repulsion forces, and it is suggested that the decrease in ionomer adsorption fraction results in the aggregation of the catalyst ink. Therefore, to realize high-performance catalyst layer, influence of acetaldehyde should be minimized and further research is requested to understand it. Acknowledgement This presentation is based on results obtained from a project commissioned by the New Energy and Industrial Technology Development Organization (NEDO).

Power Generation Charactristics of Polymer Electrolyte Fuel Cells with Electrocatalysts Supported on SnO2 in High Current Density Range

Introduction Considerable efforts have been made for FCEV commercialization, for which sufficient performance and durability have to be achieved1. SnO2-supported electrocatalysts with carbon backbone have a potential to achieve both high catalytic activity and high cycling durability2-3. However, it is necessary to further optimize cell performance up to high current densities for such alternative electrocatalysts. By controlling the thickness of the seal material and optimizing the Nafion content of the electrocatalyst layer, successful reduction of various overvoltages and improved power generation performance can be achieved, especially in the high current density region4. This study aims to further improve cell performance at high current densities by varying electrocatalyst layer fabrication conditions and selecting suitable settings for the MEAs. Experimental Using mesoporous carbon (MC) as a support backbone and SnO2 as a support surface layer on the cathode side5, membrane-electrode-assemblies (MEAs) with Pt/Sn0.92Nb0.08O2/MC electrocatalysts were fabricated. On the anode side, a standard Pt/C catalyst (TEC10E50E, Tanaka Kikinzoku Kogyo, Japan) was used. Cell performance and overvoltages up to high current densities were measured by varying the size of electrocatalyst particles. Start-stop cycle tests up to 60,000 cycles were also conducted. Results and discussion By ball milling before the Pt loading process in the preparation of the SnO2-supported electrocatalyst Pt/Sn0.98Nb0.02O2/MC, various overvoltages could be reduced, and cell performance became almost comparable to that of the standard Pt/C catalyst. Figure 1 shows the comparison of IV characteristics and Figure 2 shows the comparison of microstructures of electrocatalyst layers. Ball milling resulted in a more appropriate microstructure of the electrocatalyst layer due to the homogenization of the electrocatalyst particle size. As shown in Figure 3, the ECSA retention after 60,000 start-stop cycles was about 50%. Even when MC was used as the support framework, high durability could be achieved due to the presence of SnO2, which was further improved by ball milling. Without ball-milling, the retention was about 45%.In the future, we aim to further improve cell performance by controlling the porous microstructure and water management of the electrocatalysts layer. We will fabricate cells under various conditions, conduct performance tests, and analyze overvoltages. The start-stop and load cycle durability will also be evaluated. Based on these efforts, the latest results of cell performance and durability tests will be presented. Acknowledgment This paper is based on results obtained from a project, JPNP20003, commissioned by the New Energy and Industrial Technology Development Organization (NEDO). References Takahashi, T. Ikeda, K. Murata, O. Hotaka, S. Hasegawa, Y. Tachikawa, M. Nishihara, J. Matsuda, T. Kitahara, S. M. Lyth, A. Hayashi, and K. Sasaki, J. Electrochem. Soc., 169, 044523 (2022).Nakazato, D. Kawachino, Z. Noda, J. Matsuda, S. M. Lyth, A. Hayashi, K. Sasaki, J. Electrochem. Soc., 165, F1154 (2018).Matsumoto, M. Nagamine, Z. Noda, J. Matsuda, S. M. Lyth, A. Hayashi, K. Sasaki, J. Electrochem. Soc., 165, F1164 (2018).Ogawa, Y. Inoue, K. Yamamoto, M. Yasutake, Z. Noda, S. M. Lyth, J. Matsuda, M. Nishihara, A. Hayashi, K. Sasaki, ECS Trans., 109 (9), 241 (2022).Inoue, M. Yasutake, Z. Noda, S. M. Lyth, M. Nishihara, A. Hayashi, J. Matsuda, K. Sasaki, ECS Trans., 109 (9), 413 (2022). Figure 1

Comparison of MEA Degradation through FCV Actual Driving Test and Load-Cycle Durability Test

Introduction In research and development of high-durability polymer electrolyte fuel cells (PEFCs) for fuel cell vehicles (FCVs), accelerated stress tests (ASTs) are employed to evaluate durability of materials with a short time [1-3]. In order to validate an AST, it is necessary to confirm that the degradation mechanisms in the AST and in actual FCV operation are consistent, and to elucidate quantitative correspondence of their degradation rates. In this study, we analyzed degradation of a membrane electrode assembly (MEA) after FCV real-world driving and compared the degradation factor with that of a load-cycle durability test in a laboratory. Experiments MEA samples were extracted from an FC stack of a Toyota MIRAI 2014 model which was operated in real-world driving of 200,000 km, about 6000 h. MEA performance and electrochemical properties of them were investigated using a small differential cell with a straight flow field. As a comparison, another MEA at beginning of life (BOL) was subjected to a load-cycle durability test following a protocol encouraged by New Energy and Industrial Technology Development Organization (NEDO), Japan. Cell temperature and relative humidity were controlled at 80°C and 100%, respectively. A square-wave potential cycle shown in Figure 1 was adopted. MEA performance was investigated after cycles of 500, 1000, 3000, 10,000, and 30,000. Structures of catalyst layers and catalyst particles were analyzed for the MEAs after driving test and 30,000 cycles of the load-cycle test using scanning electron microscope (SEM), transmission electron microscope (TEM), and X-ray absorption fine structure (XAFS). Results and discussion The performance of the MEA after 200,000 km driving is close to that of the MEA after 30,000 cycles of the load-cycle test as shown in Figure 2. The performance degradation is mainly attributed to degradation of the catalyst layer, and properties of a gas diffusion layer and a proton exchange membrane are unchanged from those of the BOL sample. Degradation mechanisms of the catalyst layers in the driving test and the load cycle test are well corresponding. First, mass activities of PtCo catalyst decay to about a half of the BOL value mainly due to the decrease in electrochemically effective surface areas (ECSAs) as represented in Figure 3. Decrease rates of the ECSAs are consistent with the increase in average particle sizes of the catalyst. The catalyst size distributions of the two samples are also comparable. Then, oxygen transport resistance in the catalyst layers is increased. The resistance of the samples after the driving test and 30,000 cycles of the load-cycle test are twice as large as the BOL value. As shown in Figure 4, the transport resistance obviously exhibits inverse proportion to the catalyst surface area. This relationship represents that the primary factor of the resistance is one at vicinity of catalyst particles including ionomer permeation [4], and increase in it is attributed to decrease in the catalyst surface area. A performance simulation based on a simple 1-dimensioal MEA model supports that the decrease in the ECSA causes most of the MEA performance degradation through the deterioration in the catalyst activity and the oxygen transport resistance. Conclusions The MEA performance degradation after the 200,000 km real-world driving with the system of MIRAI 2014 model is primarily caused by catalyst metal coarsening and is well simulated by NEDO load-cycle durability test of 30,000 cycles. This result has validated the test protocol and should contribute to a target-setting of durability of future MEAs.

Suppression of PEFC Membrane Degradation By Using SnO2 As Electrocatalyst Support

Introduction Long-term durability of polymer electrolyte fuel cells (PEFCs) is required for various applications such as automobiles. It has been suggested that certain ions dissolved from electrocatalyst materials accelerate radical formation and subsequent chemical degradation of polymer electrolyte membrane (PEM) by the Fenton reaction [1]. Membrane degradation increases hydrogen and oxygen permeability and decreases proton conductivity and mechanical strength [2]. These phenomena significantly affect the electrochemical performance and durability of PEFCs. Therefore, in case metal oxides are used as catalyst support, it is necessary to check the dissolution of metal ions and to evaluate their effects on membranes. Our research group has been studying Nb-doped SnO2 (Sn(Nb)O2), showing improved durability of the electrocatalysts [3]. Here in this study, we evaluate the effects of Sn(Nb)O2 on the chemical degradation of PEMs. Experimental A commercial Pt/C catalyst (TEC10E50E, Tanaka Kikinzoku Kogyo, Japan) was used as the anode catalyst, while three types of cathode catalysts were applied and compared: 1) the commercial Pt/C; 2) Pt/MC where Pt was directly deposited on mesoporous carbon (MC); and 3) Pt/Sn(Nb)O2/MC where Pt was deposited on the Sn(Nb)O2 decorated on the MC. The durability test at open circuit voltage (OCV) was performed for 100 h. Before and after the durability test, cell performance, hydrogen crossover current density, and electrochemical surface area (ECSA) were characterized. The degree of membrane degradation was evaluated by measuring fluorine emission rate (FER) by ion chromatography, and membrane thickness and elemental distribution by energy-dispersive x-ray spectroscopy (EDS) coupled with field-emission scanning electron microscopy (FESEM). Results and discussion It was found that the use of Pt/Sn(Nb)O2/MC led to better durability in the OCV durability test. The decrease in OCV using Pt/Sn(Nb)O2/MC was smaller than that using Pt/C and Pt/MC, less than half of that Pt/C, as shown in Figure 1. The membrane thickness with Pt/Sn(Nb)O2/MC after the durability test was the same as that before the durability test, indicating that SnO2 actually suppressed the membrane degradation as shown in Figure 2. EDS elemental analysis confirmed no Sn ion dissolution from the cathode side as shown in Figure 3, also suggesting that SnO2 at the cathode suppressed the membrane degradation. In this presentation, we will quantitatively analyze the membrane degradation in using SnO2 and discuss possible mechanisms. Acknowledgment This paper is based on results obtained from a project, JPNP20003, commissioned by the New Energy and Industrial Technology Development Organization (NEDO).

Hydrogen and Oxygen Permeability through PEFC Membrane and Membrane Electrode Assembly

IntroductionThe permeation of feed gas through a proton exchange membrane (PEM) affects a PEFC performance [1, 2]. The temperature and relative humidity (RH) dependences of the hydrogen and oxygen permeability through the perfluorinated sulfonic-acid (PFSA) membrane were investigated. A membrane is considered to consist of a bulk layer sandwiched between skin layers [3]. In this study, by measuring the hydrogen and oxygen permeation flux through a membrane and plotting the permeation resistance against the membrane thickness, the bulk layer effective diffusion coefficients and the skin layer mass transport coefficients of PEM and membrane electrode assembly (MEA) were examined.ExperimentalNafion™ NR-211, NR-212, and N115 (Chemours) with different thickness (at RH of 50 %), 25.8, 51, and 127 μm, respectively, were used as PEM. Catalyst layers made of Pt-loaded carbon (Pt 50 wt%) and Nafion ionomer (I/C = 1.0) with a thickness of 8.5-10 μm on both sides of PEMs were prepared as MEA. They were installed separately in a Japan Automobile Research Institute (JARI) standard cell whose active area was 5.0 cm×5.0 cm. The cell was placed in the constant temperature and humidity oven (ESPEC, LHL-113) and the relative humidity of the supplied gas was controlled by the bubbler temperature. The hydrogen and oxygen fluxes through a membrane were measured at cell temperatures of 60-90 °C and RH of 0.2-1.0 (moisture content λ = 1.7-10.5) with different partial pressures of nitrogen diluent gas. Permeate gas was quantified by a gas chromatograph (Agilent 990 micro GC, column: MSieve5A, detector: TCD). The schematic drawing of the measurement system is shown in Fig. 1.Results and discussionThe permeate gas flux through a membrane, Nj (j = H, O), is expressed asNj = k pj Δpj , where k pj is the gas permeance and Δpj is the partial pressure difference between the supply side and the permeate side.Fig. 2 shows the temperature and RH dependences of O2 permeance through PEM and Fig. 3 shows the permeation resistance against the membrane thickness. k pO increases with increasing RH and 1/k pj becomes lower at higher temperature. The permeation resistance through MEA tends to be lower than PEM. By assuming that Nafion has a bulk layer and skin layers [3], the bulk layer and skin layer resistances were separated from a slope and an intercept of trend lines in Fig. 3. The permeation resistances through MEA do not have y-axis intercept. The fused layer can be created by hot press process in manufacturing MEA and an x-axis intercept might be a fused layer thickness. It is easier to permeate gas through the fused layer than the skin layer of PEM. The temperature dependences of the bulk layer effective diffusion coefficient, D ej (B), are shown in Fig. 4. D eH (B) and D eO (B) of PEM and MEA are on the same line, with activation energies, E a, of 23.9 and 20.6 kJ/mol, respectively at λ = 6. D ej (B) can be formulated as follows: D eH (B) = (4.26×10-11 + 6.50×10-11RH) exp(-E aH/R (1/T-1/T ref)) m2/s D eO (B) = (8.16×10-12 + 3.39×10-11RH) exp(-E aO/R (1/T-1/T ref)) m2/s T ref = 353.15 K, E aH = 0.65 λ + 20.0 kJ/mol, E aO = 21.2 kJ/mol,Reproduced D eH (B) and D eO (B) are shown in Fig. 5. The skin layer mass transport coefficient, k pj (S), can be obtained from the y-axis intercept of a trend line in Fig. 4 and it shows that only PEM has skin layers. It is necessary to consider the gas permeation through MEA not through PEM in designing PEFCs because using the transport properties through PEM may lead to overestimate the performance.ConclusionsBy measuring the hydrogen and oxygen permeation flux through PEM and MEA in PEFC, the bulk layer effective diffusion coefficients and the skin layer mass transport coefficients of hydrogen and oxygen were obtained. The bulk layer effective diffusion coefficients of PEM and MEA were found to be the same and the temperature and RH dependences of D ej (B) were formulated as a function of temperature and RH. The skin layer mass transport coefficient appears only in PEM.AcknowledgementThis work was supported by the FC Platform Program: Development of design-for-purpose numerical simulators for attaining long life and high-performance project (FY 2020–FY 2023) conducted by the New Energy and Industrial Technology Development Organization (NEDO), Japan.

Corrosion Behavior of Al Alloy A6061 Jointed with High-Tensile Strength Steel By Friction Stir Welding in NaCl Solution

In recent years, weight reduction of automobile bodies has been effective in reducing carbon dioxide emissions. For this reason, the replacement of steel, which has a large specific gravity, with aluminum alloys and carbon fiber reinforced plastics is being promoted for the right materials in the right places. In this case, an appropriate dissimilar material joining method is needed in terms of ease of assembly, joining performance, and cost. Friction stir welding (FSW) has emerged as a promising joining method for a variety of engineering applications. FSW is a solid-phase joining technique that has been shown to have significant advantages over conventional welding in terms of the strength and quality of the joints produced. However, during joining of metallic materials by FSW, microstructural changes may occur that affect their corrosion degradation. In this study, the corrosion behavior of A6061 jointed with high-strength steel (SPCC980) by FSW in aqueous NaCl solution and the effect of Al alloy microstructural changes on it were investigated.In this study, FSW was performed on A6061 and SPCC980 steel plates, and galvanic corrosion behavior was investigated by immersion tests and surface observations, as well as electrochemical measurements of Al alloys. Immersion tests were performed in 1 M NaCl solution. Micro electrochemical polarization tests of A6061 were also performed to investigate the corrosion behavior.In the immersion test of the galvanic couple of A6061 and SPCC980 steel plates in NaCl solution, galvanic corrosion occurred with the Al alloy as the anode and the steel plate as the cathode, and hydrogen evolution was observed mainly from the BM site on A6061 in the initial stage of immersion. This is due to the hydrogen evolution reaction occurring inside the pitting corrosion of the A6061. The pits was formed near Fe-rich inclusions in the BM site. In the HAZ site, corrosion propagated from Fe-rich inclusions along grain boundaries, but no pitting corrosion was formed. In the NZ site, minor localized corrosion was developed. These results indicate that the susceptibility to localized corrosion and its morphology differ in each site on the A6061. The difference in localized corrosion susceptibility on the A6061 was considered to be caused by the difference in the distribution of Fe-rich inclusions. However, the fact that the size and number of Fe-rich inclusions in each site did not differ suggests that the difference in anodic dissolution activity near Fe-rich inclusions affects pitting corrosion susceptibility in each site. A6061 alloy is a precipitation-strengthened Al alloy reinforced by β'' phase, which is a metastable phase of Mg2Si, and dense precipitation of β'' phase was observed in the BM site. Therefore, the reason for the higher localized corrosion susceptibility of BM is related to the precipitation of the β'' phase.This paper is based on results obtained from a project, JPNP14014, commissioned by the New Energy and Industrial Technology Development Organization (NEDO).

Iridium Oxide Nanorods Supported on Sb-Doped SnO2 As Highly Active Oxygen Evolution Catalysts for PEM Water Electrolysis

With the increasing installation of renewable energy generation systems such as solar and wind power, efficient energy storage devices are required to overcome the problem of fluctuations in the availability of electric energy. In particular, the splitting of water by electrolysis to produce storable hydrogen and oxygen appears to be one of the most promising approaches for large-scale energy storage. Proton exchange membrane water electrolysis (PEMWE) has received a great deal of attention due to its ability to respond quickly to renewable energy fluctuations and to operate at high current densities. However, a key limitation of PEMWE is the high cost associated with the use of high loadings of noble metal Ir catalysts at the anode electrode, which is necessary due to the slow kinetics of the oxygen evolution reaction (OER). Therefore, reducing the Ir loading at the OER electrode is essential to realize large-scale applications of PEMWE. This requires the development of highly active OER catalysts to accelerate the reaction or reduce the overpotential. Herein, we report a remarkably active OER catalyst of Ir oxide with nanorod-like structure supported on antimony-doped tin oxide (Sb-SnO2), showing an order of magnitude increase in Ir mass-specific activity than a commercial IrOx catalyst at 1.5 V vs. RHE. We also investigated the temperature dependence of the OER activity, as a way to elucidate the mechanism of enhanced catalysis, from both experimental and theoretical perspectives.The Sb-SnO2 support material (Sb dopant content 4 at.%) with fused-aggregate network structure was prepared by flame pyrolysis method.1 The as-prepared support powder was annealed at 700 oC for 2 h in air using a rotary kiln furnace. Sb-SnO2 supported IrOx nanorods (NRs) catalyst was synthesized by a colloidal method.2 The obtained IrOx NRs/Sb-SnO2 powder was heat-treated at 350 °C in N2 atmosphere for 2 h and at 150 °C in 1% H2 (N2 balance) for 2 h. The Ir loading was determined to be 49.3 wt.% by inductively coupled plasma optical emission spectrometry (ICP-OES). The OER activity was evaluated in 0.1 M HClO4 at 25 to 80 oC using the rotating disk electrode (RDE) technique.2,3 From Figure 1a, it can be seen that the Sb-SnO2 support particles present an interconnected network structure, which is considered to be beneficial to enhance the electron conduction and gas diffusion in the electrocatalytic reaction. On the support, the majority of IrOx nanoparticles are shown to have coalesced, forming rodlike structures. Figure 1b shows OER polarization curves (iR-corrected and normalized to the Ir mass) of the IrOx NRs/Sb-SnO2 catalyst compared to commercial IrOx catalyst (TANAKA Precious Metals). It is clear that the IrOx NRs/Sb-SnO2 catalyst significantly outperforms the IrOx reference catalyst over the entire potential range, and the OER onset potential of the IrOx NRs/Sb-SnO2 catalyst is lower. At 1.5 V (Figure 1c), the IrOx NRs/Sb-SnO2 catalyst exhibits a mass activity (MA) of 0.40 A mg-1 Ir, which is about 10 times higher than that of the IrOx reference catalyst. This dramatically enhanced activity makes it possible to significantly reduce the Ir loading on the anode of PEMWE. Experimental observations and density functional theory (DFT) calculations suggest that the nanorod geometry and strong metal-support interactions may play an important role in lowering the energy barrier for the crucial first step of the OER, thereby enhancing the OER activity. Acknowledgement This work was partly supported by the JSPS KAKENHI (23H02059) and by the project from the New Energy and Industrial Technology Development Organization (NEDO) of Japan.