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Hydrogen fuel is receiving a lot of attention owing to its high efficiency and eco-friendliness. The reforming of natural gas is conventionally used to generate hydrogen gas; however, the process generates carbon dioxide causing global warming. To solve the problem, water electrolysis, which produces hydrogen and oxygen gas from water, is actively studied in order to design a cheaper and more effective hydrogen generation system. The major research gap of water electrolysis is the high cost due to the rarity of catalyst materials such as platinum [1,2]. As a substitute for a platinum catalyst, transition metal phosphides represented by nickel phosphide (Ni-P) has shown the comparable catalytic activity to platinum, despite the much lower cost compared to platinum. The nickel phosphide can be fabricated with a facile electroplating process which can be more suitable synthesis process of electrocatalyst, because it is not necessary to mix with a polymer such as Nafion, disturbing emission of generated gas and decreasing surface area [3]. On the other hand, the nickel phosphide has showed low stability in acid media. According to the previous researches [4], the main reason of low stability is due to the difference of corrosion potential between nickel metal and Ni-P phases. Furthermore, they suggest three strategies to increase the corrosion resistance: (i) well-alloying of nickel and phosphorus, (ii) formation of amorphous film at grain boundary, and (iii) formation of phosphate passive film on the surface [4]. In this context, we adopted the electrochemical surface treatment for the electroplated Ni-P to improve the activity and stability of Ni-P/carbon paper (CP) catalyst. The surface treatment (“dealloying surface treatement”) is to remove the ‘isolated Ni’ which has less bonding with P, which is composed of the CV cycles inducing the repetition of oxidation and reduction of Ni-P/CP catalyst. Furthermore, it is expected to form well-alloying between nickel and phosphide (recrystallization of Ni-P). As a result, the morphology of Ni-P/CP after CV 20 cycles was similar with raw Ni-P/CP as shown in Fig. 1a, b. However, an area of P-rich portion (dark part) increased after dealloying cycles. Therefore, an atomic ratio of phosphorus increased about 186% higher after 20 cycles of CV than the Ni-P/CP which did not conduct surface treatment. In addition, the more invasion of phosphorus atom aroused a decrease of grain size of nickel metal (Fig. 1c). According to XPS data of Ni state in Fig. 1d, an intensity of peak at a binding energy of 853 eV which means a bonding of Ni2P increased after the dealloying process. That is, the bonding between Ni and P atoms increased after the treatment. An atomic ratio of P at the surface was also increased to 36.54at%, since Ni-P/CP catalyst showed 24.88at% of P. As a result, a repetition of reduction and oxidation of Ni-P during the dealloying surface treatment decreases the cystalline size of Ni and increase the P content of the Ni-P catalyst. Due to such the changes, the Ni-P/CP showed a higher activity and stability after the dealloying surface treatment. Specifically, Ni-P/CP after CV 20 cycles showed the highest HER activity, with the overpotential of 124 mV at the current density of 10 mA cm-2, since raw Ni-P/CP had the overpotential of 179 mV (Fig. 2a). And then, the current density for Ni-P/CP after CV 20 cycles was maintained at approximately 80.4% at −0.2 VRHE for 6 hours, while raw Ni-P/CP maintained a current density of 35.6%. It is presumably because the nickel atoms can use enough electrons for the HER, which is provided from well –bonded phosphorus atoms [5]. References [1] M. Balat, Potential importance of hydrogen as a future solution to environmental and transportation problems, Int. J. Hydrogen Energy 33 (2008) 4013–4029. [2] X. Zou, Y. Zhang, Noble metal-free hydrogen evolution catalysts for water splitting, Chem. Soc. Rev. 44 (2015) 5148–5180. [3] K. Shinozaki, J. Zack, S. Pylypenko, B. Pivovar, S. Kocha, Oxygen Reduction Reaction Measurement on Platinum Electrocatalysts Utilizing Rotating Disk Electrode Technique II.Influence of Ink Formulation, Catalyst Layer Uniformity and Thickness, J. Electrochem. Soc. 162 (2015) F1384-F1396. [4] A. Kucernak, V. Sundaram, Nickel phosphide: the effect of phosphorus content on hydrogen evolution activity and corrosion resistance in acid medium, J. Mater. Chem. A 2 (2014) 17435-17445. [5] Y. Pan, Y. Lin, Y. Chen, Y. Liu, C. Liu, Cobalt phosphide-based electrocatalysts: synthesis and phase catalytic activity comparison for hydrogen evolution, J. Mater. Chem. A 4 (2016) 4745–4754. Figure 1

Polymer electrolyte membrane fuel cells (PEMFC) have received much attention as a next generation energy conversion system due to the pollutant-free byproduct (water), and high efficiency. PEMFC fundamentally consists of many components such as membrane electrode assembly (MEA) gas diffusion layers (GDLs), and bipolar plates (GDLs). Among them, MEA is directly correlated to the performance and durability of PEMFC since the MEA components such as catalyst materials, supporting carbon, and an ionomer (binder) can be degraded under corrosive environment during operating Hence, it is important to develop the meticulous analysis method to diagnose the degraded condition of PEMFC; thus, fault detection have to be properly identified before severe deterioration of MEA happens. Among an in situ diagnosis tool to investigate the degraded components of MEA, the electrochemical impedance spectroscopy (EIS) technique is well-known electrochemical method to classify each resistance such as high frequency resistance (HFR, Ohmic resistance), charge transfer resistance, and Warburg impedance, separately; thus, it makes it possible to trace the degraded components of MEA. Among them, the performance of PEMFC is closely related to the activity of catalyst that is reflected as a charge transfer resistance, and therefore, usual impedance studies are mostly focused on the catalyst [1]. In fact, however, the degradation behavior of a fuel cell system is very complicated since it originated from a complex chain and combination of corrosion issues [2], such as ionomer degradation, membrane thinning, corrosion of the supporting carbon, and Pt corrosion, which are the crucial degradation factors that affect both performance and durability of PEMFC. Hence, to elaborate the impedance measurement, it is necessary to focus on the inherent properties of all components of MEA. In this context, we performed the impedance measurements under low humidity (RH 30 %) condition to emphasize the ionic resistance during PEMFC operation. By deliberately limiting the ionic conductivity, the dehydrated ionomer accentuated the influence of ionic resistance among various resistance components inside the MEA, and thus the degraded MEA can be analyzed with respect to the ionomer surrounding catalyst layer (CL), as shown in Figure 1. Moreover, the complex capacitance adopting transmission line model (TLM) circuit under non-faradaic process facilitates to investigate the double layer capacitance (EDLC) of degrading CL [3, 4] as shown in Figure 2. To verify the feasibility of this approach, we studied degradation mechanism of MEA under dehydration and cathode flooding condition, respectively, since these problems related to water management are typical issues to be solved in systemic point of view [5]. As a result, under the dehydration degradation, the increase of gas permeation induced the mixed potential formation at the anode, causing the cathode CL deterioration with surface variation of cathode CL. It could be monitored though the volume fraction of supporting carbon and ionomer. In addition, under the flooding degradation, the unbalanced current distribution caused the irregular anode deterioration, and phase differences between transported protons were reflected at horizontal aspect at low frequency region in Warburg impedance spectra. The detailed mechanism based on the aforementioned EIS measurement will be discussed in this presentation. References [1] J.-R. Kim, J.S. Yi, T.-W. Song, Journal of Power Sources, 220 (2012) 54-64. [2] M. Kandidayeni, A. Macias F, L. Boulon, S. Kelouwani, IEEE Transactions on Vehicular Technology, 68 (2019) 11462-11472. [3] D. Malevich, B.R. Jayasankar, E. Halliop, J.G. Pharoah, B.A. Peppley, K. Karan, Journal of The Electrochemical Society, 159 (2012) F888-F895. [4] S. Yoon, J.H. Jang, B.H. Ka, S.M. Oh, Electrochimica Acta, 50 (2005) 2255-2262. [5] V.A. Sethuraman, J.W. Weidner, A.T. Haug, L.V. Protsailo, Journal of The Electrochemical Society, 155 (2008) B119-B124. Figure 1

With regard to increasing energy consumption, a polymer electrolyte membrane fuel cell (PEMFC) has received lots of attention due to its pollutant-free fuel and byproduct (water). There are lots of works to improve the performance and stability of PEMFC, however, there remains unsolved problem of long-term durability directly related to cost, which are caused by the corrosive environments inside the membrane-electrode assembly (MEA). In particular, both electrochemical corrosion and chemical deterioration result in a fast degradation of PEMFC, especially under the specific degradation conditions such as fuel starvation, flooding, dehydration, and on-off cycle. Among them, a study on dehydration condition becomes the most important as next-generation fuel cell system demands high temperature (low RH) in stack operation for better tolerance to CO and the simplified water management systems in PEMFC [1]. However, the previously reported degradation mechanisms under dehydration are mostly focused on the membrane decay and the electrochemical analyses were too conventional to elucidate the mechanism. Hence, in the present study we aim at performing sequential deconvolution of the degradation mechanism of PEMFC operating at 300 mA cm-2 under low humidity of RH 30%. During the dehydration degradation process, we conducted electrochemical measurements under both humidity condition: normal condition (RH 100%) and dry condition (RH 30%). In particular, electrochemical measurement under low humidity condition allows us to exclude the influence of the water as a proton transport medium (See Figure 1(a)) and subdivide specific resistance components from the profile of the Nyquist plots by increasing the ionic resistance deliberately in order to focus on only the intrinsic properties of membrane itself as shown in Figure 1(b). As shown in Figure 2(a-b), the current density at 0.7 V of cell voltage under RH 100 % condition decreases from 620.0 to 518.8 cm-2 during initial 300 h of degradation test and additional 300 h of procedure causes additional 17.2% decrement with the increase of charge transfer resistance. Even though these results show the degraded MEA under corrosive condition, these are not enough to identify the successive degradation components. As shown in Figure 2(c), the i-V curves measured under RH 30 % shows the performance also decrease during initial 300 h of degradation process with the increase of ionic resistance from 1.15 to 3.15 ohm cm2. However, after additional 300 h (total 600 h) the abnormal phenomenon showing the temporarily improved performance was confirmed with the decrease of ionic resistance from 3.15 to 1.9 ohm cm2 as shown in Figure 2(d). The unique characteristic could seem to be induced by a shorten proton diffusion path by the membrane thinning [2]. After operating at 300 mA cm-2 for more than 600 h at RH 30%, TEM surface analyses were conducted. As shown in Figure 3(a-e), the Pt particles were grown more largely on cathode than anode, which implies that reverse current decay (cathode carbon corrosion) occurs even in the dehydration degradation. In summary, the degradation mechanism under dehydration can be separately identified in sequential order: (1) ionomer degradation, (2) membrane thinning, and (3) reverse current decay. References [1] V.A. Sethuraman, J.W. Weidner, A.T. Haug, L.V. Protsailo, Journal of the Electrochemical Society, 155 (2008) B119-B124. [2] X.-Z. Yuan, S. Zhang, H. Wang, J. Wu, J.C. Sun, R. Hiesgen, K.A. Friedrich, M. Schulze, A. Haug, Journal of Power Sources, 195 (2010) 7594-7599. Figure 1

Due to the extra-high capacity (3860 mAhg-1) and the lowest electrochemcial equilibrium potential (-3.040 VSHE), lithium metal has been considered as a promising candidate for high-energy lithium ion battery anode [1]. However, the undesired dendrite growth results in detrimental effects in the cell performance, such as cell short-circuit, low coulombic efficiency, and dead lithium, which have hampered widespread exploits of lithium metal-based batteries. Therefore, in recent year, numerous strategies to prevent or mitigate the problematic dendrite growth, including protective layers on Li metal surface, solid electrolytes, artificial solid electrolyte interphase (SEI) layer and so forth has been reported. Among them, one of the effective methods is to provide heterogeneous lithiophilic nucleation sites on the surface of current collectors as first suggested by Yan et al., [2]. They found that heterogeneous metal seeds, which have solubility of lithium in binary phase diagram, could induce selective deposition of lithium on its surface. Recently, Hou et al., reported lithiophilic Ag nanoparticles as nucleation sites by using electroless deposition for stable electrodeposition of lithium [3]. However, from a heterogeneous nucleation point of view, the morphology of electrodeposited lithium metal will be strongly affected by uniformity of nucleation seed layers, in other words, the more uniform distribution of nucleation seeds, the more stable morphology of lithium. In this study, we introduce a thermal evaporation method to deposit Ag nanoparticles as well-distributed nucleation sites. The Ag seeds with various evaporation thicknesses (1, 3 and 5 nm) were deposited on the commercial copper foils (thickness ≈ 25 um) to study the effect of evaporation thickness on the morphology of electrodeposited lithium metal. To electrodeposit lithium on Ag-seeded Cu foil and investigate the electrochemical properties of deposited Li metal electrodes, CR 2032-type coin cells were assembled in an Ar-filled glove box (CE: Li strip, WE: Ag-seeded Cu foil, and Celgard separator). We used 1M lithium bis(trifluoromethane)sulfonamide (LiTFSI) dissolved in 1,3-dioxolane/1,2-dimethoxyethane (DOL/DME, 1:1 by volume) as electrolyte without any additives. Figure 1 shows SEM surface morphologies of a pristine copper foil and Ag-seeded Cu foils with different evaporation thicknesses of Ag. After evaporation of Ag 1 nm on the pristine Cu foil (Fig. 1-(a)), it is showed that nano-sized Ag particles uniformly distributed on the surface of Cu foil (Fig. 1-(b)). In addition, from Fig. 1-(b) to Fig. 1-(d) it is clearly confirmed that the sizes of Ag particles increased with increasing evaporation thicknesses. Figure 2 shows the morphologies of electrodeposited lithium on Ag-seeded layers with different evaporation thickness. Figure 2-(a) shows SEM image of lithium metal deposited on a pristine Cu foil, which has well-known dendritic structure. In case of electrodeposited lithium on the Ag 1 nm-seeded Cu foil (Fig. 2-(b)), it also shows dendrites on its surface although heterogeneous nucleation sites are provided. However, the morphology of lithium on the Ag-3nm seeded Cu foil has distinct differences in shape and size of lithium deposits (Fig. 2-(c)). The width size of lithium deposits was smaller than those of Ag 1 nm-seeded Cu foil and pristine Cu foil, and they have interlacing vine-like structure. Further increased evaporation thickness of Ag nucleation layer to 5 nm leads to large morphological changes of electrodeposited lithium metal (Fig. 2-(d)). It shows highly dense structure with less pore and flat surface, meaning that dendrite-free lithium film was deposited. In summary, we used thermally evaporated Ag nanoparticles as heterogeneous nucleation sites to electroplate uniform lithium metal. From the results, it was concluded that through a facile thermal evaporation controllable and well-distributed heterogeneous nucleation metal seeds can be formed on the commercial Cu current collectors, leading to a dendrite-free deposition of lithium. References [1] X.-B Cheng, R. Zhang, C. -Z. Zhao, and Q. Zhang, “Toward Safe Lithium Metal Anode in Rechargeable Batteries : A Review,” Chem. Rev, 117, 10403, 2017. [2] K. Yan, Z. Lu, H. -W Lee, F. Xiong, P. -C. Hsu, Y. Li, J. Zhao, S. Chu and Y. Cui, “Selective Deposition and Stable Encapsulation of Lithium throuh Heterogeneous Seeded Growth,” Nature Energy, 1, 16010, 2016. [3] Z. Hou, Y. Yu, W. Wang, X. Zhao, Q. Di, Q. Chen, W. Chen, Y. Liu, and Z. Quan, “Lithiophilic Ag Nanoparticle layer on Cu Current Collector toward Stable Li Metal Anode,” ACS Appl. Mater. Interfaces, 11, 8148, 2019. Figure 1