ECS Meeting Abstracts

Optimizing the structure of the catalyst layer is critical to improving the proton exchange membrane fuel cell (PEMFC) performance at low Platinum (Pt) loading. A cell with reduced cathode Pt-loading for the cost reduction suffers from a limited performance at high current density, and it is a barrier to practical applications of PEMFC [1]. In addition to Pt-loadings, the types of carbon supports affect cell performance [2]. In the prior work, the nano-scale scanning transmission electron microscopy computed tomography (STEM-CT) was applied to visualize the individual Pt particle distributions in 3D at two Pt-loadings on two types of carbon supports [3]. This imaging showed that with medium surface area, low porosity carbon support (e.g., Vulcan), a large majority of Pt particles are on the surface of the carbon support. In contrast, STEM-CT showed that with a high-surface-area carbon (HSC), a majority of Pt particles are within the porous carbon support. These internal Pt particles are not in contact with the acidic electrolyte, therefore it is unclear how protons are transported to the Pt surface and how the oxygen reduction reaction (ORR) occurs on these Pt particles. In addition, prior experimental results suggest that the utilization of inner Pt particles is improved at high relative humidity with HSC since the micropores in the carbon support are likely to be filled with condensed water [3].

Here, to investigate the characteristics of those catalyst materials, we developed a transport and ORR model on 3D structures obtained from the STEM-CT images. We solve the Poisson-Nernst-Planck equations to model the electromigration and surface charge effects on the proton concentration in the porous carbon supports with a continuum approximation. Figure 1 shows the ORR current density on the Pt particles in the HSC support at the electrode potential, 0.7 V. The simulation result suggests that negative surface charge on Pt particles may be necessary for a high proton concentration and activity of inner particles in water-filled micropores at high relative humidity. This model can be used to characterize the performance of the catalyst layers in different configurations.

Acknowledgement

This work was partially supported by the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy under grant DE-EE0007271.

References

[1] A. Kongkanand and M. F. Mathias, "The Priority and Challenge of High-Power Performance of Low-Platinum Proton-Exchange Membrane Fuel Cells," J. Phys. Chem. Lett., vol. 7, no. 7, pp. 1127–1137, Apr. 2016.

[2] V. Yarlagadda et al., "Boosting Fuel Cell Performance with Accessible Carbon Mesopores," ACS Energy Lett., vol. 3, no. 3, pp. 618–621, 2018.

[3] E. Padgett et al., "Editors' Choice—Connecting Fuel Cell Catalyst Nanostructure and Accessibility Using Quantitative Cryo-STEM Tomography," J. Electrochem. Soc., vol. 165, no. 3, pp. F173–F180, 2018.

Figure 1

In recent years, the lithium iron phosphate battery is widely used in the fields of electric vehicles and energy storage because of its high energy density, long cycle life and safety^[1], but the existing battery technology was not enough to meet the requirements of electric vehicles^[2]. So it is of great importance to research performances of battery.

In this paper, the influence of different depth of discharge (DOD) on the cycle life of the battery was investigated. The specific research process is as follows, three kinds of LiFePO₄ batteries of the same type were charged and discharged at three different discharge depths (30% DOD, 50% DOD and 100% DOD) under constant conditions of 40℃and 1C (1.3A), and the discharge capacity decay curve and decay rate curve were measured after a certain number of cycles.

Discharge capacity decay curve and decay rate curve measured under different discharge depth are shown in Figure 1. As can be seen from the left graphic, the discharge capacity of the battery will have a slight increase in the early, this is because the battery anode material has not been fully activated during the initial cycle, as the cycle progresses, the electrolyte gradually penetrates into the interior of the electrode material, and the lithium ions smoothly migrate to the inside of the electrode material and undergo reversible deintercalation reaction, resulting in an increase in the capacity of the battery. In addition, At the beginning of the cycle, the depth of discharge has little effect on the capacity of the three groups of batteries. When the cycle continues, the discharge capacity of the LiFePO₄ battery gradually decreased, the attenuation of battery capacity by the depth of discharge is more and more obvious. The right capacity fading rate curve shows that battery capacity decay rate remained the same at the beginning of the cycle. At this time, the influence of the battery capacity by depth of discharge is almost independent. After the initial cycle, the deeper the depth of discharge, the faster the cell capacity decays, and there is a significant the positive correlation between the depth of discharge and the decay rate of battery capacity. Due to the different depth of discharge, the internal structure of the electrode material will occur to different degrees of deterioration.

Fig. 1 the discharge capacity decay curve and decay rate curve under different discharge depth.

It can be seen from the above studies that the effect of the battery cycle life by depth of discharge is various in different cycle stages. In the early cycle, LiFePO₄ battery capacity at different depth of discharge changes in the same law, indicating that the depth of discharge has no effect on the battery life in the early cycle. But as the cycle continues, the greater depth of discharge, the faster decay of battery capacity, the battery cycle life decline faster.

Acknowledgements

This work is financially supported by the National High Technology Research and Development Program of China (863Program, no.2015BAG01B01).

References

[1] Forgez C, Vinh Do D, Friedrich G, et al. Thermal modeling of a cylindrical LiFePO4/graphite lithium-ion battery[J]. Journal of Power Sources, 2010,195(9):2961-2968.

[2] Ritchie A, Howard W. Recent developments and likely advances in lithium-ion batteries[J]. Journal of Power Sources, 2006,162(2):809-812.

Figure 1

Abstract not Available.

We calculated the total cost of ownership (TCO) of fuel cell electric vehicles (FCEVs) in 2017, 2020, 2035, and 2050. Our TCO model incorporates proton exchange membrane fuel cell (PEMFC) cost and durability data that we obtained during our expert elicitation interviews [1]. Our assumptions, including hydrogen storage system and fuel costs, are consistent with those published by the U.S. Department of Energy [2, 3]. We characterized the uncertainty associated with FCEV life cycle costs, and we compared our FCEV projections to the DOE's 2035 projections for internal combustion engine vehicles (ICEVs), battery electric vehicles (BEVs), hybrid-electric vehicles (HEVs), and plug-in hybrid electric vehicles (PHEVs). In our sensitivity analyses, we studied the dependence of FCEV TCO on vehicle lifetime, hydrogen fuel cost, and hydrogen storage system cost. All monetary values are expressed in 2017 USD.

On median, we estimated the FCEV TCO to be $0.42/mile in 2017, $0.33/mile in 2020, $0.19/mile in 2035, and $0.18/mile in 2050 (Figure 1). Our 2017 and 2020 estimates ranged widely. Our 2017 estimates ranged from $0.40 to $0.75/mile and 2020 estimates ranged from $0.31 to $0.51/mile. By 2035, FCEVs could be competitive with ICEVs, BEVs, HEVs, and PHEVs, as long as the PEMFC stack lasts sufficiently long. If the FCEV's lifetime falls below 11 yrs, FCEVs could become more expensive than competing vehicles. At about 11 yrs, the FCEV's TCO exceeds that of the BEV. If stack durability remains a challenge, replacing the stack could be economical. Achieving a hydrogen fuel cost below $4/kg H₂ and hydrogen system cost below $390/kg H₂ could further improve FCEVs' competitiveness.

References

[1] Whiston, M. M., Azevedo, I., Litster, S., Whitefoot, K. S., Samaras, C., and Whitacre, J. F. "Expert Assessments of the Cost and Expected Future Performance of Proton Exchange Membrane Fuel Cells for Vehicles." Proceedings of the National Academy of Sciences of the United State of America. Manuscript submitted.

[2] Nguyen, T. and Ward, J. "Life-Cycle Cost of Mid-Size Light-Duty Vehicles." Record No. 16009. U.S. DOE, Washington, D. C.

[3] U.S. DOE. "Multi-Year Research, Development, and Demonstration Plan." U.S. DOE, Washington, D. C. Accessed April 3, 2018. Available at https://www.energy.gov/eere/fuelcells/downloads/fuel-cell-technologies-office-multi-year-research-development-and-22

Figure 1

LiCoO₂ has conquered the cathode market for lithium ion batteries; recently the mixed transition metal oxide lithium nickel manganese cobalt dioxide have gained attention due to their superior electrochemical properties, better thermal stability, and lower cost. Nickel-rich compositions of lithium nickel manganese cobalt oxide (NMC) have been developed as promising high-energy, high-voltage and improved rate cathode materials capable of significantly increasing the energy density of lithium-ion batteries.

Here we report a systematic comparative study on rate capability, charge discharge behavior and nickel/lithium mixing of the nickel-rich layered oxides, NMC622, NMC811 and NCA. The results showed that both NMC622 and NMC811 have pure hexagonal layered phase with Ni/Li mixing below 4%. Both samples have similar specific capacity of about 220 mAh/g at current densities, C/10; however, NMC811 demonstrated superior electrochemical performance compared to NMC611 at higher rates. The 622 sample showed a strong dependence on the electrode loading, higher electrode loading above 15 mg/cm² had poor performance compared to lower loading at the same current density in mA/cm² while NMC811 sample does not show such loading dependence. This work was supported by the DOE-EERE-Battery500 consortium.

Abstract not Available.

Lithium-ion batteries can age non-uniformly posing additional challenge in managing larger battery cells. For instance, a non-uniform distribution of solid electrolyte interphase (SEI) or plated lithium has been observed in cylindrical cells along the jelly roll length ^(1-2). The authors have suggested pressure distribution as a cause of this non-uniform ageing. This necessitates investigation of the effect of pressure on ageing. With different goals in mind, the effect of pressure on the rate of lithium-ion battery ageing has been studied previously ^(3-4).The work by Rubino et al.⁽³⁾ indicates  that the higher capacity fade in prismatic cells as compared to cylindrical cells of the same chemistry is due to the lower pressure in the former. However, the pressure was not a directly controlled parameter in this study. In contrast, the work by Arnold et al. ⁽⁴⁾ on commercial pouch cells shows that high stack pressure causes higher capacity fade, and that a small stack pressure is important to extend the life-time. However, the pouch cells were constrained along the thickness inducing non-uniform pressure on the different electrode layers and varying amplitudes of cyclic stack pressure during charge and discharge that is proportional to the external pressure. Hence, the above discrepancy of the effect of pressure calls for further investigation on ageing of model systems with single layer cells subjected to different levels of constant pressure. As an additional contribution, postmortem analysis is performed in order to understand how pressure level affects ageing mechanism.

Here we present a study on commercial NMC positive and graphite negative electrodes in small laboratory pouch cells. The pouch cells are subjected to different stack pressure levels at controlled temperature using spring loaded stainless steel plates. Additionally, in order to monitor any differences in current distribution due to pressure and ageing, groups of cells are connected in parallel as the current distribution is measured using high precision shunt resistors and a Keithly differential multimeter. Capacity and impedance are measured periodically in order to track the performance changes with aging.

A preliminary result indicates higher capacity fade in cells subjected to lower stack pressure. Results from EIS test indicate an increased ohmic resistance for high stack pressure cells, while low stack pressure cells show an increase in charge transfer resistance. Furthermore, different stack pressure on the cells of a parallel connected cell configuration results in current distribution with the cell subjected lower pressure taking slightly larger share of current. Details of the experimental techniques, and results from electrochemical and post mortem analysis will be presented and discussed during the meeting. Conclusions from this study will contribute to better understanding of the causes of non-uniform aging and suggestion for solutions to this problem through better cell design.

References

[1] M. Klett, R. Eriksson, J. Groot, P. Svens, K.S. Högstöm, R.W. Lindström, H. Berg, T. Gustafson, G. Lindbergh,

K. Edström , J.Power Sources 257 (2014) 126-137.

[2] M. Petzl, M. Kasper, M.A. Danzer, J.Power Sources 275 (2015) 799-807. 

[3] R.S. Rubino, H. Gan, E.E. Takeuchi, J. Electrochem. Soc. 148 (2001) A1029-A1033.

[4] J. Cannarella, C.B. Arnold, J.Power Sources 245 (2014) 745-751.

Abstract not Available.

Ferroelectric materials have been utilized for non-volatile memory applications [1]. Recently, Al_1-xSc_xN (ASN) films were reported to exhibit ferroelectricity with a high remnant polarization of over 100 µC/cm² which is attractive for future non-volatile memory [2]. The feasibility of FeRAM application has been confirmed based on material and electrical properties [3]. One of the issues of ASN films is the lack of reliability and endurance properties, as the number of switching cycles of 10⁴ to 10⁵ is the typical endurance of the ASN films before breakdown [4]. In this study, we infer the gradual change in the ASN film upon switching cycles.

A 50-nm-thick ASN film with x of 0.22 was characterized with TiN electrodes at top and bottom. The as-fabricated film is self-polarized in an upward direction; the negative voltage applied to the top electrode triggers the switching reversal. Capacitance-voltage (C-V) characteristics, double sweeping the voltage to the negative, and then to the positive directions, were used to observe the switching voltage. A sinusoidal wave with an amplitude of 28 V at 2 kHz was used to switch the capacitor, ending with a positive voltage applied to the top electrode.

The C-V curve for the as-fabricated sample (fresh), shown in figure 1, showed only a single peak capacitance for each sweep. The voltage at the peak capacitance during the negative sweep shifted from -20.8 to -17.4 V from the first to third sweep, larger than the shift during the positive sweep. This indicates that the switching voltage, or the coercive field, was reduced during the first negative sweep. After stressing the sample by 10³-cycle switching, the C-V curve revealed two capacitance peaks in the negative direction, as shown in figure 2; the peak at -20.2 and -15.2 V are referred to as the peaks obtained in the first and third sweeps for the fresh device. The capacitance peak at -20.2 V gradually decreased moving to merge to the peak at -15.2 V. During this transition, little change of the peak in the positive region was observed, indicating that a gradual transformation in the ASN film upon switching reversal is happening only when the film is switched from upward to downward direction. We hypothesized the creation of nitrogen-vacancy (V_N) during the switching, facilitating the movement of N atoms to reduce the coercive field. One can understand that the transformation is completed during the C-V measurement for the fresh sample as the time for the measurement is long.

In conclusion, we observed a gradual change in the switching voltage upon switching reversal. A lower coercive field was found to gradually appear after switching, decreasing, and shifting the initial coercive field.

References

[1] T. S. Böscke, et al., Appl. Phys. Lett., 99, 102903 (2011). [2] S. Fichtner, et al., J. Appl. Phys., 125, 114103 (2019). [3] C. Liu, et al., IEDM Tech. Dig. (2021). [4] S.-L. Tsai, et al., Jpn. J. Appl. Phys. 60, SBBA05 (2021).

Figure 1

The Distribution of Relaxation Times (DRT) is an important analysis tool that is capable of giving initial information from Electrochemical Impedance Spectra (EIS) with respect to the number of relaxation processes occurring in the system and their corresponding relaxation frequencies. From this the possible number of circuit elements for the equivalent circuit model in the nonlinear least square fit of the data can be selected. There are several techniques and possibilities to perform the transformation [1–3], nevertheless every possibility has its limitation and/or will remain an ill-posed problem [4]. For our EIS data, the DRT transformation with the Tikhonov regularization is used for further analysis. Hence, we want to understand and more precisely describe the effects of this transformation together with occurring pitfalls on the most commonly implemented circuit elements used to describe EIS data for characterizing Solid Oxide Fuel or Electrolysis Cells (SOFC/SOEC) operating at high temperatures. Therefore, a database with the future possibility of pattern recognition is created for DRT transformations on selected circuit elements like a resistor in series with a finite length diffusion or Warburg-short element (R-Ws) or a resistor in parallel with a constant phase element (RQ). A parameter interval is selected for each occurring element. The origin for these data comes from a series of experiments where quite well fitting Equivalent Circuit Models (ECM) were analyzed with respect to their minimum and maximum values per element. Depending on the coverage of the interval, useful intermediate numbers are selected in order to perform parameter sweeps for every possible combination together with a downstream DRT transformation step. We obtain the behavior of the DRT transformation as a function of the individual circuit element parameters and can evaluate which occurring peaks belong to a circuit element and which ones are artefacts or numerical issues in the transformation. Additionally, the regularization parameter (λ) is taken as a sweep parameter as well to see its influence on the transformation. From both perspectives, we can state a range of optimal numbers of recorded data points per decade of frequency for experiments. We also can give a direction for selecting λ depending on what exactly is to be evaluated with the respective DRT transformation. Additionally, we are able to categorize peak-forming behavior of the transformation since the selected ECMs are known and so, in theory, its transformation result. Knowledge for the DRT behavior is also useful for modeling the electrolysis process where theoretical characteristics (i-V, EIS data) are calculated in order to validate and evaluate the model.

Figure caption: Figure 1: The DRT transformation of the model RQ-RQ is converted into the frequency domain with a variation of α_P1 from 0.5 to 1.0. The fixed parameters are q₁ = 10^-1 S, r₁ = 0.303 Ω, q₂ = 10^-2 S, r₂ = 0.303 Ω, α_P2 = 0.8, λ = 10^-7.

References

[1] T. Hörlin (1998), Deconvolution and maximum entropy in impedance spectroscopy of noninductive systems, Solid State Ionics, vol. 107, no. 3–4, pp. 241–253

[2] S. Hershkovitz, S. Baltianski, Y. Tsur (2011), Harnessing evolutionary programming for impedance spectroscopy analysis: A case study of mixed ionic-electronic conductors, Solid State Ionics, vol. 188, no. 1, pp. 104–109

[3] A. L. Gavrilyuk, D. A. Osinkin, D. I. Bronin (2017), The use of Tikhonov regularization method for calculating the distribution function of relaxation times in impedance spectroscopy, Russ. J. Electrochem., vol. 53, no. 6, pp. 575–588

[4] E. Tuncer, J. R. MacDonald (2006), Comparison of methods for estimating continuous distributions of relaxation times, J. Appl. Phys., vol. 99, no. 7, pp. 074106–074106-4

Figure 1

One hundred years ago, the "century of cinema" began, with images far superior to still images. However, even in the 21st century, molecular formulae and still images still play a leading role in the world of chemistry. In 2007, I reported an electron microscopic image of a process in which a hydrocarbon molecule undergoes a conformational change in a carbon nanotube (Science, 316, 853 (2007)), and this was the first movie capturing the motions of a molecule with a single-molecule atomic-resolution real-time electron microscopic (SMART-EM) technique. This marked the beginning of the era of "Cinematic Chemistry", where we can see with our eyes how chemistry takes place at an atomistic level (Acc. Chem. Res., 50, 1281-1292 (2017). In this lecture, I will report on the new paradigms of "cinematic chemistry", where carbon nanotube plays an essential role as a substrate on or in which the chemistry takes place. Examples will include single-molecule thermodynamics and kinetics based on ultra-fast imaging of individual molecules and reaction events, in situ structural and statistical analysis of crystal growth (JACS, 143, 1763-1767 (2021).

Figure 1

The intermetallic Co₇W₆ catalysts have shown great success in chirality-specified growth of single-walled carbon nanotubes (SWCNTs). Electron microscopic techniques, particularly the in-situ techniques enable us to study the growth mechanism and the behavior of catalysts at atomic scale. It was found that the structure of Co₇W₆ nanocrystals were stable at the temperature of 1100 °C under carbon feeding condition. No carbon dispersion within the nanocrystal happened. These observations illustrate why such kind of catalysts can act as the structural template to grow SWCNTs with specified chiralities. Due to the less efficient carbon diffusion and supply on the surface of catalysts in such a vapor-solid-solid process than that in a vapor-liquid-solid process, SWCNTs normally nucleate on larger catalyst nanocrystals. All these behaviors are distinctly different from those of normal metallic catalysts, demonstrating the uniqueness of intermetallic Co₇W₆ catalysts in chirality-selective growth of SWCNTs. These results can help us to further understand the mechanism of the origination of chirality selectivity in SWCNT growth, benefiting the rational design of catalysts.

We discuss the FC-CVD synthesis of SWNTs, especially the tuning of tube atomic structure i.e. (n,m) distributions and subsequently the thin film color. By introducing various amount of CO₂ with CO as the carbon source and in-situ ferrocene decomposition generated Fe catalyst nanoparticles, we directly produced SWNT films with tunable (n,m) i.e. helicity distribution as well as tunable colors. When operating the FC-CVD reactor at the ambient pressure and at 850 ^oC temperature with 0.25 and 0.37 volume percent of added CO₂, the SWNT films display green and brown colors, respectively. We ascribed various colors to suitable diameter and narrow (n,m) distributions, which were determined in detail using the electron diffraction. We calculate the color maps for all SWNT (n,m) structures.

N. Wei et al. Adv. Mater. 2021, 33, 2006395.

Single-walled carbon nanotubes have been a candidate for outperforming silicon in ultrascaled transistors, but the realization of nanotube-based integrated circuits requires dense arrays of purely semiconducting species. In order to directly growth such nanotube arrays on wafers, control over kinetics and thermodynamics in tube-catalyst systems plays a key role, and further progress requires the comprehensive understanding of seemingly contradictory reports on the growth kinetics. Here, we propose a universal kinetic model that decomposes the growth rates of nanotubes into the adsorption and removal of carbon atoms on the catalysts, and provide its quantitative verification by ethanol-based isotope labeling experiments. While the removal of carbon from catalysts dominates the growth kinetics under a low supply of precursors, our kinetic model and experiments demonstrate that chiral angle-dependent growth rates emerge when sufficient amounts of carbon and etching agents are co-supplied. The kinetic maps, as a product of generalizing the model, include several kinetic selectivities that emerge depending on the balance of gases with opposing effects. Our findings not only resolve discrepancies existing in literature, but also offer rational strategies to control chirality, length, and density of nanotube arrays for practical applications.

Part of this work was supported by JSPS (KAKENHI JP20K15137, JP20H00220) and JST (CREST JPMJCR20B5).

Self-supporting, sponge-like paper of carbon nanotubes (CNTs) can be a new platform for high energy/power density batteries. Presently, battery electrodes are built on metallic foils having only two faces and large mass with the aid of polymeric binders and conductive fillers. CNT paper can replace all these components. Furthermore, it captures nanomaterials at high loading, activates insulative materials, and moderates volumetric change of active materials.

We have developed semi-continuous fluidized bed chemical vapor deposition method [1], which yields submillimeter-long few-wall CNTs with carbon purity >99 wt% and metal impurity <0.1 wt% [2]. By capturing capacitive particles (99 wt%) with CNTs (1 wt%), a lithium ion full cell with graphite-CNT anode and LiCoO₂-CNT cathode is demonstrated (Figure 1b) [3]. This platform is also effective to create electrodes of the emerging active materials such as Si anode [4] and S cathode [5] with practically high gravimetric, areal, and volumetric loadings, and realized Li_xSi-S and Li-Li₂S_x full cells [6,7]. In addition, by combining with the insulative sponge of boron nitride nanotubes (BNNTs), the graphite-CNT|BNNT| LiCoO₂-CNT full cell with high thermal stability was developed [8].

[1] D.Y. Kim, et al., Carbon 49, 1972 (2011).

[2] Z. Chen, et al., Carbon 80, 339 (2014).

[3] K. Hasegawa and S. Noda, J. Power Sources 321, 155 (2016).

[4] T. Kowase, et al., J. Power Sources 363, 450 (2017).

[5] K. Hori, et al., J. Phys. Chem. C 123, 3951 (2019).

[6] K. Hori, et al., Carbon 161, 612-621 (2020).

[7] Y. Yoshie, et al., Carbon 182, 32-41 (2021).

[8] K. Kaneko, et al., Carbon 167, 596-600 (2020).

Figure 1. (a) Semi-continuous production of submillimeter-long few-wall CNTs by fluidized bed [1]. (b) Lithium ion full cell made of graphite-CNT anode and LiCoO₂-CNT cathode [3].

Figure 1

Optimizing the structure of the catalyst layer is critical to improving the proton exchange membrane fuel cell (PEMFC) performance at low Platinum (Pt) loading. A cell with reduced cathode Pt-loading for the cost reduction suffers from a limited performance at high current density, and it is a barrier to practical applications of PEMFC [1]. In addition to Pt-loadings, the types of carbon supports affect cell performance [2]. In the prior work, the nano-scale scanning transmission electron microscopy computed tomography (STEM-CT) was applied to visualize the individual Pt particle distributions in 3D at two Pt-loadings on two types of carbon supports [3]. This imaging showed that with medium surface area, low porosity carbon support (e.g., Vulcan), a large majority of Pt particles are on the surface of the carbon support. In contrast, STEM-CT showed that with a high-surface-area carbon (HSC), a majority of Pt particles are within the porous carbon support. These internal Pt particles are not in contact with the acidic electrolyte, therefore it is unclear how protons are transported to the Pt surface and how the oxygen reduction reaction (ORR) occurs on these Pt particles. In addition, prior experimental results suggest that the utilization of inner Pt particles is improved at high relative humidity with HSC since the micropores in the carbon support are likely to be filled with condensed water [3].

Here, to investigate the characteristics of those catalyst materials, we developed a transport and ORR model on 3D structures obtained from the STEM-CT images. We solve the Poisson-Nernst-Planck equations to model the electromigration and surface charge effects on the proton concentration in the porous carbon supports with a continuum approximation. Figure 1 shows the ORR current density on the Pt particles in the HSC support at the electrode potential, 0.7 V. The simulation result suggests that negative surface charge on Pt particles may be necessary for a high proton concentration and activity of inner particles in water-filled micropores at high relative humidity. This model can be used to characterize the performance of the catalyst layers in different configurations.

Acknowledgement

This work was partially supported by the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy under grant DE-EE0007271.

References

[1] A. Kongkanand and M. F. Mathias, "The Priority and Challenge of High-Power Performance of Low-Platinum Proton-Exchange Membrane Fuel Cells," J. Phys. Chem. Lett., vol. 7, no. 7, pp. 1127–1137, Apr. 2016.

[2] V. Yarlagadda et al., "Boosting Fuel Cell Performance with Accessible Carbon Mesopores," ACS Energy Lett., vol. 3, no. 3, pp. 618–621, 2018.

[3] E. Padgett et al., "Editors' Choice—Connecting Fuel Cell Catalyst Nanostructure and Accessibility Using Quantitative Cryo-STEM Tomography," J. Electrochem. Soc., vol. 165, no. 3, pp. F173–F180, 2018.

Figure 1