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Platinum group metal-free (PGM-free) catalysts are a low-cost alternative to expensive PGM catalysts for polymer electrolyte fuel cells. However, due to the low volumetric activity of PGM-free catalysts, the catalyst layer thickness of the PGM-free catalyst electrode is an order of magnitude higher than PGM based electrodes. The thick PGM-free electrodes suffer from increased transport resistance and poor water management, which ultimately limits the fuel cell performance. This manuscript presents the study of water management in the PGM-free electrodes to understand the transport limitations and improve fuel cell performance. In-operando neutron imaging is performed to estimate the water content in different components across the fuel cell thickness. Water saturation in thick PGM electrodes, with similar catalyst layer thickness to PGM-free electrodes, is lower than in the PGM-free electrodes irrespective of the operating conditions, due to high water retention by PGM-free catalysts. Improvements in fuel cell performance are accomplished by enhancing water removal from the flooded PGM-free electrode in three ways: (i) enhanced water removal with a novel microporous layer with hydrophilic pathways incorporated through hydrophilic additives, (ii) water removal through anode via novel GDL in the anode, and (iii) lower water saturation in PGM-free electrode structures with increased catalyst porosity.

Abstract Proton exchange membrane fuel cell (PEMFC) electrodes have a complex structure of carbon supported platinum (Pt) nanoparticles intermixed with proton conducting ionomer. This structures creates a porous network facilitating transport of electrons, protons, and reactants. These catalyst layers have been shown to exist with a non-homogenous distribution of ionomer, with aggregates and agglomerates of carbon and ionomer. To better understand the catalyst layer structure, we have applied atomic force microscopy (AFM) to investigate fuel cell catalyst layer properties at high resolution. Various ionomer-carbon (I/C) ratios were analyzed in order to identify correlations between carbon agglomerate size, pore distribution, and structure. Surface and cross-sectional images of topography, adhesion, and roughness were collected. Scanning electron microscopy (SEM) was performed to allow for additional depictions of the samples. Several other techniques are planned to further explore structures. Electrochemical capabilities of the AFM will be utilized to inspect the effects of current on the aforementioned properties and provide better control in measurements. Additionally, samples will be placed in a water-filled holder to study effects of swelling through in situ AFM. Transmission electron microscopy (TEM) will also be performed to complement these measurements. A Bruker AFM was used in PeakForce tapping mode with a ScanAsyst Air tip to obtain the images in Figure 1. Figure 1a shows the height map of a sample with an I/C ratio of 1. Figure 1b is the adhesion map of the same sample, showing clear spatial adhesion contrast. In this map, lighter colors correspond to the more adhesive ionomer. Darker areas signify hard materials like carbon. Figure 1c depicts the adhesion map overlaid on the height map, indicating how the adhesion and topography relate. The arrows in Figure 1c illustrate where the ionomer thickness can be measured between carbon aggregates. Figure 2a summarizes the average ionomer thickness calculated for a range of I/C ratios from 0.25 to 1.25. There is a positive correlation between I/C ratio and average ionomer thickness, suggesting that increased availability of ionomer leads to thicker surface coatings and more separation between the carbon aggregates. Figure 2b shows the relationship between I/C ratio and aggregate diameter. These values were measured by processing the images in ImageJ and reducing aggregates to approximated spheres. Although no clear relationship is apparent now, further investigation with differing I/C ratios may reveal a trend. Acknowledgements : This research is supported by the U.S. Department of Energy Fuel Cell Technologies Office, through the Fuel Cell Performance and Durability (FC-PAD) Consortium (Technical Development Manager: Greg Kleen and Fuel Cells Program Manager: Dimitrios Papageoropoulos). This work was also performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science. Figure 1

Cerium ions enhance the chemical stability and lifetime of polymer electrolyte membrane (PEM) fuel cell components by rapidly and reversibly scavenging degrading radical species which are generated during operation [1]. However, during cell fabrication and discharge, these ions readily migrate between the membrane and catalyst layers (CLs) of the membrane electrode assembly (MEA) [2]. Complete washout of cerium from the MEA has also been observed [3]. It is necessary to understand the mechanisms and magnitude of cerium migration during fuel cell operation, since cerium ions are ineffective outside of the catalyzed area of the MEA. These results can be applied to improve cerium stability in the active area of the MEA and localize it to areas of high radical generation, which can further extend the lifetime of PEM fuel cells. Table 1: AST conditions and flow fields Test Time (h) RH (%) Flow field I 500 30 50 cm2 single serpentine II 2,000 100 25 cm2 tri-serpentine III 2,000 100 25 cm2 single serpentine Membrane chemical stability accelerated stress tests (ASTs) [4] were performed on cerium-containing MEAs at 90°C in single-cell hardware (Fuel Cell Technologies), compressed with 8 x bolts at 50 in-lb of torque, using the conditions and hardware shown in Table 1. Nafion XL (DuPont) membranes were used, which contain a nominal cerium loading of 6 μg Ce/cm2. Carbon-supported platinum electrodes (TKK, 48% Pt, 0.1-0.2 mg Pt/cm2loading) and Sigracet 25BC GDLs (SGL) were also used. X-ray fluorescence (XRF) was performed on MEA components before and after the ASTs in order to measure in-plane cerium content in the membrane and CLs. After 500 hours of OCV operation at 30% RH, cerium moved uniformly from the membrane of MEA I into the CLs (not shown). Here, migration is attributed cell component hot pressing and interactions with the CLs [2]. Membrane cerium was reduced to 3.7± 0.68 μg Ce/cm2, while anode and cathode CL concentrations were increased to 2.3 ± 0.06 and 3.4 ± 0.22 μg Ce/cm2, respectively. At 100% RH, cerium also remained in the active area, however, membrane concentration increased from inlet to outlet (Figure 1a). This gradient may arise due to the increased presence and flow of liquid water during 100% RH operation. Only trace amounts of cerium remained in the CLs, except near the outlet, which suggests that under humidified conditions, the effects of CL interactions on cerium migration are reduced. After 2,000 hours of operation at 100% RH, flow field compression was observed to have implications on cerium migration out of the active area. The cerium profile of MEA III (Figure 1b) shows that it migrated from areas of high compression in the active area (shown in red) into low compression regions outside of the active area. In contrast, compression was higher around the active area of MEA II (Figure 1a), which prevents cerium from leaching from it. However, concentration was non-uniform, as discussed above. These preliminary results indicate that cerium migration and leaching out of the active area are affected by membrane water content and cell clamping pressure. It is believed that other factors such as electrical potential and temperature influence migration, as well. The authors wish to acknowledge the financial support of the Fuel Cell Technologies Program and the Technology Development Manager Nancy Garland. References [1] F. D. Coms, H. Liu, J. E. Owejan, ECS Trans., 16, 1735-1747 (2008). [2] S. M. Stewart, D. Spernjak, R. L Borup, A. Datye, F. Garzon, ECS Lett., 3, F19-F22 (2014). [3] M. V. Lauritzen, S. Knights, T. Cheng, D. W. Banham, E. Kjeang, A. Sadeghi Alavijeh, Fuel Cells 2014 Science & Technology, April 2014, Amsterdam, The Netherlands [4] U.S. DOE, Cell Component Accelerated Stress Test Protocols for PEM Fuel Cells, 2010. Figure 1

PEM fuel cells (PEMFCs) show great promise to increase the fuel efficiency for transportation applications; however, for this application, they must show performance and durability with the requirements for transportation. For transportation applications, the fuel cell will be subjected to frequent power cycling. For example, the DOE/Fuel Cell Tech Team (FCTT) protocol for durability includes load cycling from 0.02 A/cm2 to 1.2 A/cm2 every 0.5 min. The cathode catalyst and catalyst layer have been shown as susceptible to degradation causing loss of performance due to both loss of kinetics for the oxygen reduction reaction and loss of mass transport. Catalyst support-carbon corrosion can result in thinning of the catalyst layer contributing to degradation in performance. To examine the effect of power cycling in situ on carbon corrosion and electrode degradation, we directly measured the catalyst support degradation by measuring CO2 in the cathode outlet by NDIR (Non-Dispersive Infra-Red) while operating a single-cell fuel cell. CO2 present in air was removed by a lime bed prior to introduction to the fuel cell. We operated with a modified DOE/FCTT durability protocol using controlled voltage, and varied the potential limits to explore the effects of the upper potential limit, lower potential limit, the potential step size and time at potential. The upper potential limit was varied from 0.95 to 0.55V; the lower potential limit from 0.40V to 0.80V, with times ranging from 0.5 min to 5 min. The corrosion of three different types of carbon were explored, high surface area (E), vulcan (V), and graphitized (EA). The catalyst support carbon corrosion occurs under normal fuel cell operating conditions and is exacerbated by the voltage cycling inherent in these steps in potential. A series of carbon corrosion spikes during potential cycling is shown in Figure 1 for E-type carbon, varying the upper potential from 0.95 V to 0.60V while keeping the lower potential constant at 0.40V. Sharp spikes in the carbon corrosion rate are observed during a step increase in cell potential with the magnitude of the spikes decreasing as the high cell potential is reduced from 0.95 V to 0.6 V. The carbon corrosion rate at high cell potential (0.95V) decreases with time at potential, indicating formation of passivating carbon surface oxides. Carbon corrosion was measured during the drive cycle measurements for all three types of carbon, with the relative carbon corrosion rates of E>V>EA. The series of step potential carbon corrosion spikes where the potential was varied for the lower potential from 0.40 V to 0.60V while keeping the upper potential constant at 0.95V shows similar results in terms of the carbon corrosion. The magnitude of the spikes decrease as the lower cell potential is raised. These results indicate that the size of the step in potential has a more significant impact on the carbon corrosion rate than does the absolute value of the potential for normal cathode operating potentials. The peak in CO2 evolution occurs when the cell potential increases from high power operation to low power near open circuit. This correlates with when CO2 evolution is observed during cyclic voltammograms, which occurs during the positive sweep at ~ 0.55 to 0.60 V. The evolution of this CO2 peak suggests that oxygen is adsorbed onto the carbon and/or CO is formed on the Pt surface. During long-term operation, a reduction in catalyst layer thickness is observed during drive cycle operation, which can be due to the loss of carbon through carbon corrosion or possibly due to compaction; both effects likely lead to a loss of void volume. This reduction in thickness includes a sharp decrease in catalyst layer thickness within the first 100 hours of operation (30%), eventually reaching ~50% of its thickness after 1000 hours. Most of this reduction in electrode thickness does not appear to be directly due to carbon corrosion as there is little evidence for carbon corrosion from microscopic analysis, especially during the early stages of operation, where the thickness reduction is substantially more than what should be due to carbon corrosion. These results show that carbon corrosion occurs during normal potential operation, and is exacerbated by potential variations during operation. To minimize the carbon corrosion, the size of the steps in potential should be minimized. For a more stable catalyst, another requirement is for the Pt particles to be stabilized on graphitized carbon supports. Acknowledgments Funding for this work is from DOE EERE FCTO, Technology Development Manager Nancy Garland Figure 1

Catalyst support materials are an essential component of proton exchange membrane (PEM) fuel cells to decrease catalyst Pt loading. To be classified as a good catalyst support, a material must have high surface area, high electrical conductivity, be durable in harsh environments, and be economically feasible for large scale commercial production. In light of this, carbon is one such material that has widely been used in PEM fuel cells as a catalyst support. Various carbon materials of differing morphologies have been utilized as supports where high surface area is beneficial to greater Pt accessibility and higher performance, while graphitization and lower surface areas are beneficial to durability especially at high potentials. Here, we present a durability analysis of platinum cobalt catalyst supported on five structurally different carbon supports which are as follows: high surface area carbon TEC36E32 (29% Pt / 2.3% Co / 526.6 m2/g-Cat), high surface area graphitized carbon TEC36EA32 (29.3% Pt / 3.3% Co / 108.5 m2/g-Cat), Vulcan carbon TEC36V32 (29.1% Pt / 3.6% Co / 160 m2/g-Cat), graphitized Vulcan carbon TEC36VA32 (29.3% Pt / 3.4% Co / 70 m2/g-Cat), and acetylene black TEC36F32 (29.3% Pt / 3.4% Co / 523.6 m2/g-Cat). Each of the aforementioned carbon supported dealloyed-PtCo catalysts were spray coated onto one side of a DuPont® XL membrane at a loading of approximately 0.1 mgcat/cm2. The other side of the membrane, which served as the anode, was kept the same for all samples and spray coated in the same manner at the same loading with platinum supported on Vulcan carbon TEC10V20E (20% Pt). The membrane electrode assemblies (MEAs) were then incorporated into 5 cm2 differential cells with SGL Sigracet® Gas diffusion layers (GDLs). Subsequent to conditioning, catalyst performance was evaluated via the following characterizations: local oxygen transport resistance measurement, fuel cell polarization curves, sheet resistance measurement, impedance measurement in air and HelOx, mass activity measurement, and cyclic voltammetery measurements. DOE recommended square wave catalyst accelerated stress testing (AST) was utilized to degrade the catalysts with a recovery protocol and characterization performed after 15,000 and 30,000 cycles of the AST.1 Figure 1 illustrates the performance of TEC36V32 and TEC36E32 cathode catalyst MEAs where power densities of 0.643 W/cm2 and 0.716 W/cm2 were obtained respectively at 80oC, 100%RH and 150kPa pressure. In addition to this characterization, evolution of the resistances that arise from oxygen transport during the catalyst durability testing, both pressure dependent and pressure independent, will be presented for the various carbon supports. Further characterization of the catalysts after the testing will also be presented to provide insight on the influence of carbon support materials on catalyst performance with cell degradation. Acknowledgement: This research is supported by the U.S. Department of Energy Fuel Cell Technologies Office, through the Fuel Cell Performance and Durability (FC-PAD) Consortium (Fuel Cells Program Manager: Dimitrios Papageorgopoulos and Technical Development Manager: Greg Kleen). References: 1United States Department of Energy Fuel Cells Technologies Office Multi-Year Research, Development, and Demonstration Plan. https://www.energy.gov/sites/prod/files/2017/05/f34/fcto_myrdd_fuel_cells.pdf. Figure 1

This work examines the effect of various hydrocarbons on fuel processor light-off and reforming. Major hydrocarbon fuel constituents, such as aliphatic compounds, napthanes, and aromatics have been compared with the fuel processing performance of blended fuel components and reformulated gasoline to examine synergistic or detrimental effects the fuel components have in a real fuel blend. Short chained aliphatic hydrocarbons tend to have favorable light-off and reforming characteristics for catalytic autothermal reforming compared with longer-chained and aromatic components. Oxygenated hydrocarbons have lower light-off requirements than do pure hydrocarbons. Gas phase oxidation favors higher cetane # fuels, which tend to be longer chained hydrocarbons. Energy consumption during the start-up process shows a large fuel effect. Methanol and dimethylether (DME) show lower start-up energy demands for the fuel processor start-up than do high temperature reforming hydrocarbon fuels such as methane, gasoline and ethanol. Aromatics and longer chained hydrocarbons show a higher tendency for carbon formation, increasing the amount of carbon formed during the light-off phase while the addition of oxygenates tends to lower the carbon formed during the start-up process.

Alloying Pt with transition metals, such as Co and Ni, enhances the catalytic activity for the oxygen reduction reaction (ORR) at cathodes of polymer electrolyte fuel cells (PEFCs). The alloying process necessarily involves a high-temperature thermal annealing step that grows the size of catalyst nanoparticles above 4-5 nm. Particles of this size are known to be more stable under cyclic potentials as they are less prone to Pt dissolution and coalescence than finer scale 2-3 nm Pt particles. However, transition metals are unstable at operating fuel cell potentials and can leach out of the catalyst in the cathode and into the ionomer, resulting in loss of ORR activity. Three formulations of Pt-Co alloy catalysts supported on high surface area carbon were subjected to accelerated stress tests (ASTs) in H2-N2 environment. The AST protocol consisted of repeated square waves with 0.6 V lower potential limit (LPL), 0.95 V upper potential limit (UPL), and 3-s holds at LPL and UPL. Changes in cell performance were monitored by measuring polarization in H2-air, high frequency resistance (HFR), electrochemical surface area (ECSA), impedance in air and helox (79% He, 21% O2), and mass activity. These measurements were made at beginning of life (BOL) and after 1000 (1k), 20,000 (20k) and 30,000 (30k) cycles. The three catalysts had a different initial Co content (34%, 20%, and 15%) and electrode microstructure. At BOL, transmission electron microscopy (TEM) analysis of the high-Co content (H) catalyst electrode, 0.2 mg/cm2 Pt loading, showed a “spongy” porous (hollow) catalyst nanoparticle morphology with ~6.7 nm median Pt-Co particle size after conditioning. The BOL medium-Co content (M) catalyst electrode, 0.1 mg/cm2 Pt loading, showed a crystalline fully alloyed catalyst nanparticle morphology with ~4.4 nm particle size after conditioning. The BOL low-Co content (L) catalyst electrode, 0.1 mg/cm2 Pt loading, also showed catalysts with a fully alloyed crystalline morphology with ~4.5 nm median particle size. Post-mortem TEM analysis of aged electrode specimens showed different amounts of Co dissolution from the nanoparticles. There is an apparent direct relationship between the initial alloy Co content the amount of Co dissolved after potential cycles. WAXS analysis confirms significant Co loss from the nanoparticles and also indicates that the Pt-Pt d-spacing in Pt-Co alloy approaches the values typical of pure Pt particles. The effect of Co dissolution on the loss in specific activity for ORR is summarized in Fig. 1a. The data suggest that the specific activity depends not only on the initial Co content but also on the catalyst nanoparticle morphology and size. Even with 27-50% Co loss, the specific activity of Pt-Co alloy remains higher than 1000 mA/cm2-Pt, and exceeds the activity (650 mA/cm2-Pt) of pure Pt of similar particle size. The measured specific activities are consistent with a kinetic model and catalyst-specific choice of kinetic parameters. The modeled ORR kinetic losses depend on specific activity, ECSA, catalyst loading, and kinetic parameters (electrode morphology). As shown in Fig. 1b, the H-catalyst electrode shows the smallest kinetic overpotentials (ηc s) due to high Pt loading, but has the fastest increase in ηc s with ageing because of low Co retention in the electrode. The L-catalyst electrode shows higher ηc s due to lower Pt loading but has a smaller increase in ηc s with ageing because of high Co retention. For this catalyst, increase in is mostly due to loss of ECSA with ageing. The M-catalyst electrode shows similar increase in ηc s with ageing as the L-catalyst electrode, but this increase is due to the combined effects of Co dissolution and ECSA loss. The modeled increase in mass transfer overpotentials (ηm) with ageing correlates with Pt loading and ECSA loss and depends on the initial electrode morphology. There is no clear evidence of a relationship between Co dissolution (i.e., effect of Co2+ on O2 permeability in ionomer film) and the resulting increase in ηm. Acknowledgments The authors wish to acknowledge the financial support of the Fuel Cell Technologies Office and Fuel Cell Component R&D Team Lead, Dr. Dimitrios Papageorgopoulus, and Technology Development Manager, Dr. Nancy Garland. The authors also wish to acknowledge General Motors, IRD and Umicore for supplying materials used in this study. The submitted manuscript has been created by UChicago Argonne, LLC, Operator of Argonne National Laboratory (“Argonne”). Argonne, a U.S. Department of Energy Office of Science laboratory, is operated under Contract No. DE-AC02-06CH11357. Figure 1

The performance and durability of proton exchange membrane fuel cells (PEMFCs) is of great interest for vehicular power applications. In traditional PEMFCs, Pt catalysts are used in both the anode and the cathode. However, efforts to improve PEMFCs often exposes the membrane electrode assembly of PEMFCs with other metallic cations. For example, reducing the cost of PEMFCs and in particular of Pt-based catalysts can be accomplished with Pt-Ni or Pt-Co alloying catalysts [1-3]. Additionally, Cerium (Ce) is incorporated in membranes in order to improve the chemical durability of the membrane [4]. While these materials effectively improve performance and durability, they have shown to leach out contaminants such as Ce3+, Co2+, and Ni2+ions into the ionomer [5-7]. Such leaching can reduce the efficacy of the additives, decrease ionomer conductivity and reduce cell performance. In this work, we have studied the impact of metallic cationic contaminants (Ce, Co and Ni) in the catalyst layer ionomer on the oxygen reduction reaction (ORR) of a Pt microelectrode in various conditions. A well-established microelectrode mini-cell set-up was used for the electrochemical measurements of the ORR activity at the Pt-ionomer interface with and without the presence of cationic contaminants in concentrations observed during fuel cell operation. Nafion was doped with cation contaminant and subsequently deposited on Pt micro-electrodes. X-ray fluorescence (XRF) was used to verify that the contaminant concentrations in the thin films was equivalent to experimentally determined concentrations in the CL ionomer [4]. Results obtained using electrochemical and other characterization techniques will be presented to further understand the role of the contaminants in the performance of Pt electrodes for the ORR. These techniques include laser profilometry to measure the thin film thickness, conductivity cells and impedance measurements to understand the impact of contaminants on the ionomer resistance. Results indicate that in presence of Ni, we observe a decrease in ionomer conductivity (Nafion® 211), from 102 mS cm-1 in proton form to 35 mS cm-1 with 5.2 mgNi cm-3 at 80°C, 100% RH. This performance correlates with a decrease in ORR performance, most notably lower limiting currents 1.1x10-1 mA cm2 in proton form vs. 6.3x10-2 mA cm-2 with 7 mgNi cm-3and lower ORR on-set potentials at 25°C, 100% RH. References 1. Kang, Y., et al., Multimetallic core/interlayer/shell nanostructures as advanced electrocatalysts. Nano letters, 2014. 14(11): p. 6361-6367. 2. Cui, C., et al., Compositional segregation in shaped Pt alloy nanoparticles and their structural behaviour during electrocatalysis. Nat Mater, 2013. 12(8): p. 765-771. 3. Chi, M., et al., Surface faceting and elemental diffusion behaviour at atomic scale for alloy nanoparticles during in situ annealing. Nature Communications, 2015. 6: p. 8925. 4. Coms, F.D., The chemistry of fuel cell membrane chemical degradation. ECS Transactions, 2008. 16(2): p. 235-255. 5. Baker, A.M., et al., Cerium migration during PEM fuel cell accelerated stress testing. Journal of The Electrochemical Society, 2016. 163(9): p. F1023-F1031. 6. Antolini, E., J.R.C. Salgado, and E.R. Gonzalez, The stability of Pt–M (M = first row transition metal) alloy catalysts and its effect on the activity in low temperature fuel cells: A literature review and tests on a Pt–Co catalyst. Journal of Power Sources, 2006. 160(2): p. 957-968. 7. Mani, P., R. Srivastava, and P. Strasser, Dealloyed binary PtM3 (M = Cu, Co, Ni) and ternary PtNi3M (M = Cu, Co, Fe, Cr) electrocatalysts for the oxygen reduction reaction: Performance in polymer electrolyte membrane fuel cells. Journal of Power Sources, 2011. 196(2): p. 666-673. LA-UR-17-23298

The cost and durability of Polymer Electrolyte Membrane Fuel Cells (PEMFCs) are the two major barriers to the commercialization of these systems for stationary and transportation power applications.1 Given the stringent performance requirements that PEMFCs have to meet for automotive applications, Pt-alloy catalysts are being considered as cathode catalysts. While PtCox alloy catalysts exhibit better mass activity than Pt/C catalysts and can meet the beginning of life performance requirements, their durability is not proven. Several studies have indicated that the base metal leaches out of these catalysts over time and any benefits due to alloying is lost over the lifetime of the Membrane Electrode Assembly (MEA).2 The U.S. DOE Fuel Cell Tech Team has recommended various ASTs for PEMFC components.3,4 Los Alamos National Laboratory (LANL) has previously reported on the durability of several Pt/C catalysts during both simulated drive cycle testing and accelerated stress tests (ASTs).5 A catalyst AST that involved applying a square wave with 3 sec at 0.6V and 3 sec at 0.95V was determined to be an appropriate AST for evaluating the durability of Pt electrocatalysts.6In this study the performance of PtCo/C cathode catalyst based MEAs subjected to this AST were evaluated. The durability of PtCo/C catalysts was found to be comparable or better than that of the Pt/C based catalysts. Moreover the activity of the PtCo/C catalyst based MEA was also enhanced with respect to the Pt/C MEA. Figure 1 illustrates the evolution of the electrochemical surface area (ECSA) over 30,000 cycles of the square wave AST. The PtCo/C based MEA exhibited little change in ECSA over the course of the test. The ECSA initially increased due to insufficient conditioning and then decreasing due to catalyst particle size increase. In contrast, the Pt/C catalyst showed significant increase in ECSA over the course of the test. Finally the starting ECSA of the PtCo/C catalyst based MEA was significantly lower than that of the Pt/C MEA, while the end of test (EOT) ECSA was almost identical. These results are consistent with the small (≈ 2nm) starting particle size of the Pt and relatively large (≈ 5nm) starting particle size of the PtCo catalyst. The performance of the two MEAs before and after the durability AST is presented in Figure 2. It should be noted that the Pt/C based MEA had a loading of 0.15mg.Pt/cm2 whereas the PtCo/C MEA had an initial loading of 0. 21mg.Pt/cm2. The beginning of test (BOT) performance of the PtCo/C based MEA was significantly better in the kinetic region due to both the alloying effect and higher Pt loading. However, the BOT performance of the Pt/C MEA was better in the mass transport region due to extra mass transport limitations observed in alloy catalysts. The End of test (EOT) performance of the PtCo/C based MEA is only slightly higher than that of the EOT performance of the Pt/C based MEA, consistent with the fact that most of the Co has leached out of the PtCo catalyst. In this presentation the durability implications of using PtCo alloy catalysts will be discussed in detail with results presented from TEM, SEM, XRF, XRD and electrochemical testing. References 1. R. Borup, et al., Chemical Reviews, V. 107, No. 10, 3904-3951 (2007). 2. S.C. Ball, S.L. Hudson, B.R.C. Theobald and D. Thompsett, “PtCo, a Durable Catalyst for Automotive Proton Electrolyte Membrane Fuel Cells?” ECS Trans., 11 (1), 1267 (2007). 3. DOE Cell Component AST and polarization curve Protocols for PEM Fuel Cells (Electrocatalysts, Supports, Membranes and MEAs), Revised December 16, 2010. 4. N. L. Garland, T.G. Benjamin, J. P. Kopasz, ECS Trans., V. 11 No. 1, 923 (2007). 5. R. Mukundan, G. James, J. Davey, D. Langlois, D. Torraco, W. Yoon, A. Z. Weber, and R. Borup, “Accelerated Testing Validation” ECS Trans., 41 (1), 613 (2011) 6. R. Borup et al. DOE Hydrogen and Fuel Cell Program, FY15 Annual Progress Report. https://www.hydrogen.energy.gov/pdfs/progress15/v_e_1_borup_2015.pdf Acknowledgements The authors wish to acknowledge the financial support of the Fuel Cell Technologies Program and Technology Development Managers: Dimitrios Papageorgopoulos and Nancy Garland. The authors also wish to acknowledge SGL Carbon for the GDLs used in this study. Figure 1

Polymer electrolyte membrane (PEM) fuel cell performance and materials degradation, particularly associated with the cathode catalyst layer (CCL), can be directly attributed to the structure and chemistry of individual material components, as well as their uniformity/homogeneity within a CCL. The individual material constituents used to form the CCL within the membrane electrode assemblies (MEAs), which include the electrocatalyst, catalyst support, and ionomer films, and especially the critical interfaces that are formed between these various constituents, are critically importance in controlling fuel cell perfomance. Understanding the specific microstructural characteristics of the individual materials within the CCL, and how the materials interact to “form” the CCL, is important for identifying materials optimization parameters that can significantly enhance performance and durability. Materials in several states/conditions, e.g., prior to incorporation in the CCL (as-synthesized), after MEA preparation (CCL), and after fuel cell testing, are beig evaluated and quantified using a combination of advanced electron microscopy methods, which are used to interrogate the materials constituents and interfacial structures and chemistries from the μm- to the Å-level. The as-processed (prior to and following incorporation into a CCL) microstructural evidence is directly correlated with observations of materials-specific degradation mechanisms that contribute to fuel cell performance loss, and are then used to identify potential processing variables (materials-based mitigation strategies) to improve the microstructure and compositional homogeneity within the electrode structure, and enhance MEA durability and stability during fuel cell operation. Research efforts at Oak Ridge National Laboratory are focused on the high-resolution microstructural and microchemical characterization of MEAs fabricated using different electocatalysts (typically Pt-based) and catalyst loadings, carbon-based support materials, and ionomer solutions, as well as the same MEAs subjected to accelerated stress tests (ASTs) designed to degrade specific MEA components. While a significant microscopy effort has been aimed towards understanding catalyst degradation (e.g., coarsening, de-alloying), recently, high-resolution analytical microscopy methods have been used to directly image/map the distribution and chemistry of the ionomer films/layers within CCLs, results of which are being combined with high-resolution imaging and 3-D tomography data of powder materials and MEAs, to provide unprecedented insight into the MEA architecture and interfaces (ionomer/support, ionomer/catalyst, catalyst/support, ionomer/pore). This presentation will focus on understanding materials distributions within CCLs as a function of materials used and ink/MEA processing variables, e.g., initial ionomer and/or ink chemistry, electrocatalyst (type, content, and dispersion), and the type of carbon support used. Additionally, the stability of the ionomer films, electrocatalysts, and support structures in CCLs following ASTs designed specifically for either catalyst degradation or carbon corrosion, will be evaluated. ________________________________________________________________________ Research sponsored by (1) the Fuel Cell Technologies Office, Office of Energy Efficiency and Renewable Energy, U.S. Department of Energy and (2) Oak Ridge National Laboratory’s Center for Nanophase Materials Sciences (CNMS), which is a U.S. Department of Energy, Office of Science User Facility.

One of the major degradation involves the electrocatalyst, including the corrosion of the carbons used as catalyst supports which leads to changes in the catalyst layer structure. These changes in the catalyst layer structure lead to reduced kinetics and to mass transport limitations. The catalyst layer degradation is due to a combination of degradation mechanisms, including localized effects and depends upon operating conditions