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Nano Energy

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Nanoalloy catalysts inside fuel cells: An atomic-level perspective on thefunctionality by combined in operando x-ray spectroscopy and totalscattering

Valeri Petkova,⁎, Yazan Maswadeha, Yinguang Zhaob, Aolin Lub, Hannah Cronkb,Fangfang Changb, Shiyao Shanb, Haval Kareemb, Jin Luob, Chuan-Jian Zhongb, Sarvjit Shastric,Peter Keneseic

a Department of Physics and Science of Advanced Materials Program, Central Michigan University, Mt. Pleasant, MI 48859, United StatesbDepartment of Chemistry, State University of New York at Binghamton, Binghamton, NY 13902, United Statesc X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, Argonne, IL 60439, United States

A R T I C L E I N F O

Keywords:Electrocatalysts for fuel cellsSynchrotron x-ray spectroscopy and totalscatteringIn operando structure-function studies

A B S T R A C T

We introduce an experimental approach for structural characterization of catalysts for fuel cells combiningsynchrotron x-ray spectroscopy and total scattering. The approach allows probing catalysts inside operating fuelcells with atomic-level precision (~ 0.02 Å) and element specificity (~ 2–3 at%) in both time (~ 1min) andspace (~ µm) resolved manner. The approach is demonstrated on exemplary Pd-Sn and Pt-Ni-Cu nanoalloycatalysts for the oxygen reduction reaction (ORR) deposited on the cathode of an operating proton exchangemembrane fuel cell. In operando x-ray data show that under operating conditions, the catalyst particles canundergo specific structural changes, ranging from sub-Å atomic fluctuations and sharp nanophase transitions to agradual strain relaxation and growth, which inflict significant losses in their ORR activity. Though triggeredelectrochemically, the changes are not driven solely by differences in the reduction potential and surface energyof the metallic species constituting the nanoalloys but also by the formation energy of competing nanophases,mismatch between the size of individual atomic species and their ability to interdiffuse fast in search of en-ergetically favorable configurations. Given their complexity, the changes are difficult to predict and so theresulting ORR losses remain difficult to limit. We show that in operando knowledge of the structural evolution ofnanoalloy catalysts helps create strategies for improving their activity and stability. In particular, we show thatshaping Pd-Sn nanoalloys rich in Pd as cubes reduces the interdiffusion of atoms at their surface and so makesthem better catalysts for ORR in fuel cells in comparison to other Pd-Sn nanoalloys. In addition, we demonstratethat the approach introduced here can provide knowledge of other major factors affecting the performance offuel cells such as operating temperature and the overall catalyst utilization, in particular the evolution of ele-mental and mass distribution of catalyst particles over the cells’ cathode. Last but not least, we discuss how inoperando x-ray spectroscopy and total x-ray scattering can bridge the knowledge gap between the widely used insitu SAXS, EXAFS and monocrystal surface XRD techniques for structural characterization of nanoalloy catalystsexplored for energy related applications.

1. Introduction

Fuel cells are a viable alternative to mankind's dependence on fossilfuels. Unfortunately, fuel cells are not on the market yet, mostly be-cause of the lack of efficient and affordable catalysts for the sluggishchemical reactions driving cells’ operation, such as the oxygen reduc-tion reaction (ORR). The reaction takes place at the cathode of virtuallyall fuel cells and, without loss of generality, can be expressed as O2

+4H+ +4e- →H2O, that is, oxygen molecules adsorbed and reducedat the cell's cathode react with protons supplied by the cell's anode toform water. Studies have shown that the ORR proceeds via a number ofsteps including, among others, a dissociative adsorption of molecularoxygen and removal of reaction intermediates such as atomic oxygenand hydroxyl (OH) groups. The current understanding is that an effi-cient catalyst for the ORR would bind oxygen molecules with sufficientstrength to facilitate the cleavage of O˭O bonds but weakly enough to

https://doi.org/10.1016/j.nanoen.2018.04.049Received 27 November 2017; Received in revised form 24 March 2018; Accepted 16 April 2018

⁎ Corresponding author.E-mail address: petko1vg@cmich.edu (V. Petkov).

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release the reaction intermediates and product when the reaction iscomplete [1–6].

Pure Pt is the best monometallic catalyst for the ORR, even though,according to theory, it binds oxygen somewhat too strongly by about0.2 eV [7]. Unfortunately, pure Pt is not suitable for large-scale appli-cations mostly because it is one of the world's scarcest metals. Binaryand ternary nanoalloys involving noble metals, such as Pt, Au and Pd,transition metals, such as Ni, Co, Fe and Cu, and post-transition metals,such as Sn and Pb, have proven rather affordable and more activecatalysts for the ORR as compared to pure Pt nanoparticles (NPs)[8–12]. Qualitatively, the improvement has been attributed to one ormore of the following factors: i) ligand/electronic effects arising fromcharge exchange between the noble and non-noble metal atoms, ii)strain effects arising from the difference between the size of noble andnon-noble metal atoms and iii) geometric effects where particularconfigurations of surface noble and non-noble metal atoms appearbeneficial to the ORR [11–14]. When used inside fuel cells though, thevery promising nanoalloy catalysts produced so far are often found lessefficient than expected. The reason is that the catalyst particles wouldundergo specific atomic-level changes that inflict significant losses intheir ORR activity, thereby limiting the cells’ performance. The problemtriggered off a number of studies on the structural transformations ofnanoalloy catalysts exposed to the harsh environment inside operatingfuel cells, including in situ synchrotron SAXS, EXAFS and monocrystalsurface XRD [5,9,12,15–19]. Despite extensive research, the drivingforce, scope and dynamics of the atomic-level changes of functioningnanoalloy catalysts remain not well understood and so the resultingORR activity losses remain difficult to limit. Here we introduce anothersynchrotron x-ray radiation based approach for studying catalysts in-side fuel cells that combines the proven advantages of energy dispersivex-ray spectroscopy (EDS) and high-energy total x-ray scattering incharacterizing nanosized materials at atomic level [20,21]. The greatcapabilities of the approach are summarized in Scheme 1.

For the sake of generality, we demonstrate the usefulness of theapproach by applying it on binary and ternary nanoalloy catalysts in-volving different noble, transition and post-transition metals. The na-noalloy catalysts probed here have been a subject of ex situ studies forORR activity and stability but never evaluated inside an operating fuelcell. In particular, we study three members of the family of Pd-Sn na-noalloys (with Pd ≥ 50 at%), which have been shown to exhibit bothhigh catalytic activity for the ORR and enhanced tolerance for ethanoland methanol [22]. Largely, the improved ORR activity of Pd-Sn na-noalloys is explained in terms of charge transfer from the less

electronegative Sn (1.96) to the more electronegative Pd (2.2.) atoms.The resulting charge-deficient Sn and negatively charged (electron-rich) Pd atoms in the nanoalloys would bind oxygen species strongerand weaker comparing to pure Sn and Pd surfaces, respectively, therebyaccelerating the ORR kinetics [23,24]. We also study a catalyst from thefamily of Pt-Ni-Cu nanoalloys (Pt< 50 at%). Here, the improvement inORR activity is attributed to charge transfer from the less electro-negative transition metal atoms to more electronegative Pt atoms (1.90for Ni and Cu vs 2.28 for Pt) and, furthermore, shortening of Pt-Ptbonding distances, i.e. compressive strain. The latter is due to the dif-ference in the elemental size of Pt (2.77 Å), Ni (2.50 Å) and Cu (2.56 Å)atoms. Both the charge transfer and compressive strain would lead to adownshift of the d-electron band of surface Pt atoms and, hence, di-minishing the “poisoning” effect of ORR intermediates. Besides, it isconsidered that the presence of two types of transition metal atoms inthe nanoalloys creates opportunities for manoeuvring of oxygen speciesover different surface sites, thereby enabling an extra bi-functionalmechanism for accelerating the ORR kinetics [9–11,25–27]. Note thataccording to Pauling's theory of chemical bonding [28], and as shownby independent experimental [29] and DFT studies [30], the length andstrength of bonds between atoms in alloys and the exchange of chargebetween atoms involved in the bonds are strongly coupled. Hence, herewe consider that the electronic and strain effects in Pd-Sn and Pt-Ni-Cunanoalloy catalysts largely act together at atomic level and refer tothem as bonding effects. Furthermore, we use the noble metal atom-involving bonding distances as a relative measure of the bonding effectsi.e. strength of metal-to-metal atom interactions, in the respective na-noalloys. Lastly, we validate the approach on standard pure Pt andPt3Ni catalysts.

2. Experimental

Nanoalloy Pd-Sn and Pt-Ni-Cu particles were synthesized followinga wet chemistry route described in the Supplementary information (SI).The as-synthesized NPs were deposited on carbon black and activatedfor catalytic applications by a post-synthesis treatment involvingheating at 260 °C in N2 for 30min followed by heating at 400 °C in15 vol% H2 for 120min. The overall chemical composition of activatedPd-Sn and Pt-Ni-Cu alloy NPs, hereafter referred to as fresh Pd-Sn andPt-Ni-Cu alloy NPs, was determined by inductively coupled plasmaatomic emission spectroscopy (ICP-AES) and confirmed by independentEDS as described in the SI. Fresh Pd-Sn alloy NPs appeared with anoverall chemical composition of Pd43Sn57, Pd78Sn22 and Pd95Sn5. The

Scheme 1. Mechanistic description of ORR atthe cathode of a PEMFC and in operando EDS &HE-XRD experiments on the PEMFC cathodecatalyst. Hydrogen (H2) molecules are deliv-ered to the PEMFC anode and split into protonsand electrons via a reaction known as the hy-drogen oxidation reaction (HOR). The re-sulting protons permeate through a membraneelectrode assembly (MEA) to the PEMFCcathode while electrons reach it via an externalcircuit, thus generating current output. Oxygen(O2) molecules reduced at the PEMFC cathodereact with the protons arriving at it, formingwater. The reaction, known as ORR, is six or-ders of magnitude slower than HOR, necessi-tating a very efficient catalyst, e.g. nanoalloyparticles, for speeding it up. Thanks to theirhigh-flux and energy, synchrotron x-rays areable to probe the PEMFC cathode catalyst withsub-Å precision, in both time (~ min) and

space (~ µm) resolved manner. Experimental data provide information for the evolution of ORR activity, chemical composition, nanophase content, geometricsurface area, 3D structure and bonding interactions between the constituent atoms of the PEMFC cathode catalysts as they function under actual operating conditions.The catalysts’ utilization, related to the evolving chemistry and mass distribution of the catalysts’ particles over the cathode side of MEA, is also assessed.

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chemical composition of fresh Pt-Ni-Cu alloy NPs turned out to bePt37Ni29Cu34.

X-ray photoelectron spectroscopy (XPS) measurements described inthe SI (see Figs. S18 and S19) confirmed that Pd-Sn particles are na-noalloys wherein the individual Pd and Sn atoms influence each other'selectronic structure, including transfer of charge from Sn to Pd atoms.XPS data for Pt37Ni29Cu34 alloy NPs indicated a downshift of the oc-cupied d-electron states of Pt atoms, likely rendering surface Pt sites inthe NPs less reactive in comparison to pure Pt surfaces.

The size and shape of fresh Pd-Sn and Pt-Ni-Cu alloy NPs were de-termined by Transmission Electron Microscopy (TEM) as described inthe SI. Exemplary TEM and high-resolution (HR)-TEM images areshown in Figs. S1 and S2, respectively. As can be seen in the Figures, theNPs are approximately 7.8 (± 1.5) nm in size and spherical in shape.In addition, as indicated by the clear lattice fringes appearing in therespective HR-TEM images, the NPs possess a good degree of crystal-linity. The size of fresh NPs was also determined from the full width atthe half maximum (FWHM) of strong Bragg-like peaks in their high-energy x-ray diffraction (HE-XRD) patterns. Within the limits of theexperimental accuracy, HE-XRD and TEM determined sizes of fresh Pd-Sn and Pt-Ni-Cu alloy NPs appeared about the same.

The catalytic activity of fresh Pd-Sn and Pt-Ni-Cu alloy NPs for theORR was determined by ex situ cyclic voltammetry (CV) and rotatingdisk electrode (RDE) experiments conducted on a standard 3-electrodecell. The experiments are described in the SI. Results are shown in Fig.S3. As can be seen in the Figure, fresh Pd78Sn22 alloy NPs are superiorto both fresh Pd43Sn57 and Pd95Sn5 alloy NPs in terms of mass activity(MA) for ORR. Upon undergoing 520 potential cycles though, Pd95Sn5alloy NPs show a significantly increased whereas Pd43Sn67 andPd78Sn22 alloy NPs show a greatly decreased ORR activity, respectively.Evidently, upon undergoing electrochemical cycling, Pd95Sn5 alloy NPsacquire a structural state promoting the ORR over their surface. Noprecise information for the changes of the phase type and content,chemical composition, size and bonding interaction between the con-stituent atoms of Pd-Sn alloy NPs occurring during the ex situ electro-chemical cycling though was possible to obtain.

A PEMFC optimized for HE-XRD experiments was used to reveal theatomic-level changes and resulting losses in the ORR activity of Pd-Snand Pt-Ni-Cu alloy NPs functioning under actual operating conditions.The PEMFC is shown in Fig. S4. Membrane electrode assemblies (MEAs)for the PEMFC were prepared as described in the SI. Following a pro-tocol recommended by DOE [31], the PEMFC was first cycled between0 V and 1.2 V for about 1 h and then cycled between 0.6 V and 1.0 V foranother 5–6 h. The cycling was carried out at 80 °C. In total, the na-noalloys underwent about 1600 potential cycles as deposited on thecathode side of MEAs. During the cycling, high purity hydrogen andnitrogen gas were fed to the PEMFC anode and cathode compartments,respectively. Pure Pt NPs were used as the PEMFC anode catalyst. Thecurrent output of the PEMFC was non-stop recorded during the po-tential cycling. Exemplary polarization curves showing the currentoutput of the PEMFC resulting from the repetitive application of ex-ternal voltages are shown in Fig. 1c. The curves are typical for ac-celerated tests for activity and stability of nanoalloy catalysts for ORR(e.g. see Fig. 2 in Ref. [32]). Following the protocol of ex situ catalyticstudies, values of the MA of Pd-Sn and Pt-Ni-Cu alloy NPs for ORR wereobtained from the current output of the PEMFC at 0.9 V on the polar-ization curves. Rate of changes in the current (in %) with the potentialcycling are summarized in Fig. S5. Note that data in Fig. S5 are nor-malized against the PEMFC current/MA values for the respective fresh,i.e. not cycled NPs. Hence, hereafter, we refer to the data as effectiveMA activity of cycled Pd-Sn and Pt-Ni-Cu NPs for ORR. As can be seenin the Figure, and in line with the findings of ex situ catalytic studies,the ORR activity of cycled Pd43Sn67 and Pd78Sn22 alloy NPs steadilydecreases during the PEMFC operation. On the other hand, after aninitial decay, the ORR activity of Pd95Sn5 alloy NPs rebounds a bit andkeeps its relatively high value with the further potential cycling.

Atomic-level changes of Pd-Sn and Pt-Ni-Cu nanoalloy catalyststaking place during the PEMFC operation (potential cycling), includingchanges in their chemistry, geometric surface area, phase content and3D atomic structure, were probed by EDS and total x-ray scatteringcoupled to atomic pair distribution function (PDF) analysis.Experiments were done at the 1-ID-C beamline of the Advanced PhotonSource at the Argonne National Laboratory using x-rays with energy of100.2 keV (λ=0.1236 Å). Details of the experimental set-up are shownin Fig. S4. HE-XRD patterns were collected in intervals of 1minthroughout the PEMFC operation. Experimental in operando HE-XRDpatterns for cycled Pd-Sn and Pt-Ni-Co alloy NPs are shown in Fig. S6.As can be seen in the Figure, the patterns exhibit a few strong Bragg-likepeaks at low diffraction angles and almost no sharp features at highdiffraction angles, i.e. are rather diffuse in nature. The average size ofcycled NPs was determined from the FWHM of the strongest Bragg-likepeak in the respective HE-XRD patterns using the Scherrer formalism[33]. Results showed that the average size of Pd43Sn57, Pd78Sn22,Pd95Sn5 and Pt37Ni29Cu34 alloy NPs undergone 1600 potential cyclesinside the PEMFC increased from 7.8 (± 1.5) nm to about 9.5 (± 1.5)nm, 9.1 (± 1.5) nm, 8.8 (± 1.5) nm and 8.6 (± 1.5) nm, respec-tively. The so-called geometric surface area (GSA) of nanoalloy cata-lysts is inversely proportional to the size of nanoalloy particles. In turn,generally, the GSA is proportional to the so-called electrochemical ac-tive surface area (ECSA) of the catalysts, which is vital for achievingand maintaining high ORR activity [5,6,31,34,35]. The GSA of Pd-Snand Pt-Ni-Cu alloy NPs was calculated from the experimental values forthe NP size and found to diminish gradually during the PEMFC op-eration. The rate of change of the GSA, and so of ECSA, of Pd-Sn and Pt-Ni-Cu alloy NPs during the PEMFC operation is shown in Fig. S7. Therather diffuse nature of in-operando HE-XRD patterns though preventedapplying well-established procedures for in-depth structural char-acterization of bulk metals and alloys to the nanoalloy particles studiedhere. Therefore, the patterns were analyzed in terms of atomic pairdistribution functions (PDFs) shown in Fig. 2. Details of the derivationof atomic PDFs from HE-XRD patterns can be found in the SI. Here it isto be underlined that experimental total x-ray scattering data, i.e. datacomprising both Bragg-like and diffuse x-ray intensities, and so theirFourier counterparts, the atomic PDFs, reflect ensemble-averagedstructural features of all metallic alloy NPs sampled by the x-ray beamin a way traditional powder XRD patterns reflect ensemble-averagedstructural features of all polycrystalline particles sampled by the x-raybeam in those experiments. Using NP-ensemble averaged structuralfeatures to understand and explain NP-ensemble averaged functionalproperties (e.g. catalytic, magnetic, optical, etc.) puts the exploration ofatomic-level structure and properties of metallic alloy NPs on the samefooting [36]. EDS spectra for cycled Pd-Sn and Pt-Ni-Cu alloy NPs weretaken concurrently with the HE-XRD patterns, that is, in intervals of1min. Exemplary EDS spectra are shown in Fig. S8. From the spectra,changes in the overall chemical composition of the NPs during thePEMFC operation were determined following standard procedures [37].The changes are summarized in Fig. 1a, b and, in better detail, in Fig.S9.

3. Results and discussion

As data in Fig. 1a show, on average, the chemical composition ofPd43Sn57, Pd78Sn22 and Pd95Sn5 alloy NPs undergone 1600 potentialcycles inside the PEMFC has changed to Pd74Sn26, Pd65Sn35 andPd81Sn19, respectively. Evidently, due to the highly corrosive conditionsat the PEMFC cathode, metallic species have leached from the NPsduring the PEMFC operation. Based on the differences in the reductionpotential (−0.14 V for Sn vs 0.92 V for Pd) and surface energy (0.79 J/m2 for Sn vs 2.15 J/m2 for Pd) of Sn and Pd, one may expect that Snspecies in cycled Pd-Sn NPs would tend to segregate towards theirsurface and, furthermore, leach from the NPs in large numbers. Ex-perimental data though show that it is always the majority species that

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has leached from the respective Pd-Sn NPs. Furthermore, as data inFig. 2b show, the chemical composition of Pt37Ni29Cu34 NPs undergone1600 potential cycles inside the PEMFC has changed to Pt52Ni17Cu31.Here, likely due to their rather low reduction potential (−0.25 V for Nivs 0.34 V for Cu vs 1.19 V for Pt), minority Ni species have primarilyleached from the cycled Pt-Ni-Cu NPs. The observed significant changesin the chemical composition of cycled Pd-Sn and Pt-Ni-Cu alloy NPsprompted us to conduct both qualitative and quantitative (nano)phaseanalysis of the NPs. The analysis involved a “search and match” pro-cedure based on rigorously comparing atomic PDFs derived from modelnanophases with the respective in operando PDF data sets. A modelreproducing the experimental data in very good detail, i.e. minimizingthe residual between the model computed and experimental PDF dataset, was considered a truthful representation of the nanophase type andcontent of the respective NPs. Model nanophases were based on crystalstructures adopted by bulk Pd-Sn and Pt-transition metal alloys (tran-sition metal=Ni, Cu). For reference, bulk Pt-(Ni/Cu) and particularly

Pd-Sn alloys exhibit a number of intermetallic phases over a broadrange of alloys’ compositions (see Figs. S10, S11 and S12). Fragments ofthe phase types explored in the nanophase analysis are shown in Fig.S13. Exemplary results of the nanophase analysis of fresh and cycledPd-Sn and Pt-Ni-Cu alloy NPs are shown in Figs. 3c, d and 4, S14 andS15. More details of the nanophase analysis can be found in the SI.

As can be seen in Fig. 3d, fresh Pd43Sn57 NPs appear as an orthor-hombic nanoalloy. Soon after the PEMFC stars operating, that is withinthe first hour of PEMFC operation, Pd43Sn57 NPs suffer big loss in Snspecies (see Fig. 1a) and segregate into a mixture of a dominant hex-agonal- (87 vol%) and secondary cubic-type (13 vol%) nanophases.After five more hours of potential cycling, the NPs remain phase seg-regated into a mixture of a hexagonal (79 vol%) and cubic-type (21 vol%) nanophases. As can be seen in Fig. 4, fresh Pd78Sn22, Pd95Sn5 andPt37Ni29Cu34 NPs appear as cubic-type nanoalloys. Within the first hourof potential cycling though, Pd78Sn22 NPs suffer significant loss in Pdspecies (see Fig. 1a) and segregate into a mixture of a dominant

Fig. 1. Evolution of the chemical composition and nanophase type of (a) Pd-Sn and (b) Pt-Ni-Cu NPs during the electrochemical cycling. Phase diagrams of therespective bulk counterparts are shown on the right. (c) Selected polarization curves representing the PEMFC current output recorded during the electrochemicalcycling of Pd-Sn and Pt-Ni-Cu NPs in the range from 0.6 V to 1.0 V. The curves tend to collapse, i.e. the PEMFC current output diminishes, during the PEMFCoperation (see the thin black arrows on both sides of the curves). Following the protocol of ex situ catalytic studies, the current output of the PEMFC at 0.9 V on thepolarization curves (follow the thick orange arrows) is used as a measure of the effective MA of Pd-Sn and Pt-Ni-Cu NPs for ORR.

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orthorhombic (82 vol%) and secondary cubic-type (12 vol%) nano-phases. Upon undergoing 1600 potential cycles, the NPs end up as amixture of an orthorhombic (74 vol%) and cubic-type (26 vol%) nano-phases. Likewise, within the first hour of PEMFC operation, cycledPt37Ni29Cu34 NPs suffer loss in both Cu and Ni metal species (seeFig. 1b) and segregate into a mixture of a tetragonal (44 vol%) and

cubic-type (56 vol%) nanophases. During five more hours of potentialcycling, the NPs suffer losses mostly in Ni species and remain phasesegregated. On the other hand, regardless having suffered significantloss in Pd species (see Fig. 1a), cycled Pd95Sn5 NPs retain their cubic-type phase state throughout the PEMFC operation. Notably, as data inFigs. 1, 3 and 4 show, both the structure type of fresh Pd-Sn and Pt-Ni-

Fig. 2. In operando atomic PDFs for Pd-Sn and Pt-Ni-Cu NPs vs time of electrochemical cycling. PDFs are obtained from the HE-XRD patterns shown in Fig. S6.

Fig. 3. (a) 2D scedasticity map of the in operando atomic PDFs for Pd43Sn57 alloy NPs. The map shows that the NPs undergo significant structural changes during thePEMFC operation. Among others, the changes include a relatively sharp nanophase transition (thick bright spots) and continuous readjustment of atomic-pairdistances/strain relaxation (thin vertical lines). (b) Color map of the first peak in the in operando atomic PDFs for Pd43Sn57 NPs. The map shows that the structuralchanges of the NPs also include highly dynamic sub-Å atomic fluctuations likely arising from re-packing of nearby Pd and Sn atoms (see the sequence of bright anddark spots). (c) Exemplary phase analysis for fresh and cycled Pd43Sn57 NPs. The time of cycling is given for each data set. The experimental PDF data (symbols)cannot be reproduced by a model (red line) featuring a single phase fcc-type nanoalloy. As shown in (d), fresh Pd43Sn57 NPs are indeed an orthorhombic-typenanoalloy whereas cycled Pd43Sn57 NPs segregate into a mixture of a majority hexagonal and minority cubic nanophases. Note that per EDS data, the composition ofNPs cycled for about 6 h has changes from Pd43Sn57 to Pd74Sn26 due to leaching of Sn species.

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Cu alloy NPs and the trajectory of structural phase transformations theyexperience during the PEMFC operation mirror the phase diagram ofthe respective bulk alloys.

Major atomic-level changes in metallic nanoalloy catalysts func-tioning under real operating conditions, in particular structural nano-phase transitions, are seen as big changes in the number, position andintensity of major peaks in the concurrently obtained in operandoatomic PDFs (e. g. see the PDF data sets for Pd43Sn57 NPs shown inFig. 3c and d). Hence, the changes are easy to recognize by con-ventionally inspecting the PDFs (e.g. see the color maps of PDF peaksshown in Fig. 3b and S20). Subtle but important atomic-level mod-ifications of functioning nanoalloy catalysts such as, for example, localstrain relaxation, may be overshadowed by strong features in the re-spective PDFs and so are easy to miss. Therefore, we analyzed the inoperando atomic PDFs for cycled Pd-Sn and Pt-Ni-Cu NPs in terms of thestatistical scedasticity formalism. As discussed in the SI, changes instatistical properties of a time series of successive in operando PDFs,such as scedasticity, can reveal subtle atomic-level modifications of thestudied NPs. 2D scedasticity map computed from the in operando PDFsfor Pd43Sn5 NPs is shown in Fig. 3a. 2D scedasticity maps for Pd78Sn22,Pd95Sn5 and Pt37Ni29Cu34 NPs are shown Fig. S20(first row). The mapsconfirm that cycled Pd-Sn and Pt-Ni-Cu alloy NPs experience majorstructural changes soon after the PEMFC starts operating. Besides, themaps reveal that the characteristic atomic pair distances and so localstrain in the NPs gradually evolve during the PEMFC operation. To bemore precise, Pd-involving bonding distances in fresh Pd95Sn5, Pd78Sn22and Pd43Sn57 alloy NPs appear expanded by about 1.5%, 1.9% and4.6%, respectively, as compared to the bonding distance in bulk Pd(2.755 Å). Bonding distances in Pd95Sn5 and Pd78Sn22 NPs undergone1600 cycles appear further expanded, thereby increasing the localtensile strain in the NPs to 1.8% and 2.5%, respectively. Bonding dis-tances in Pd43Sn57 NPs undergone 1600 cycles though shrink, dimin-ishing the local tensile strain in the NPs to about 2.9%. When comparedto the bonding distance in bulk Pt (2.775 Å), Pt-involving bondingdistances in fresh Pt37Ni29Cu34 NPs appear compressed by 5.3%. Uponundergoing 1600 potential cycles, the local compressive stain in theNPs diminishes to 3.7%. Markedly, as the oscillations of the intensity ofthe first peak in the respective in operando PDFs show (see the colormaps in Fig. 3b, S16 and S20), the structural changes of cycled Pd-Snand Pt-Ni-Cu NPs also involve frequent sub-Å atomic fluctuations likelyrelated to a re-packing of nearby atoms in the NPs.

The observed continuous changes in the chemical composition andaverage size of cycled Pd-Sn and Pt-Ni-Cu alloy NPs signal the presenceof an effective mass transport within the cathode side of MEA. Themagnitude of the mass transport was evaluated by measuring andcomparing the elemental and mass distribution (loading) of catalystparticles over the PEMFC cathode for three characteristic time periodsof the PEMFC operation. Measurements were done by scanning a4.5 mm by 4.5mm area of the MEAs using an x-ray beam with a foot-print of about 40 µm by 20 µm. Maps of the elemental and mass dis-tribution of cycled Pd43Sn57 alloy NPs over the PEMFC cathode ob-tained after the PEMFS has operated for about one, two and a half, andfive hours are shown in Fig. 5. Maps indicative of the time evolution ofthe elemental and mass distribution of Pd78Sn22, Pd95Sn5 andPt37Ni29Cu34 alloy NPs over the PEMFC cathode are shown in Figs. S21,S22 and S23, respectively. Inspection of the maps reveals that, duringthe PEMFC operation, catalyst particles accumulate in some and vir-tually disappear from other mm-size areas of the PEMFC cathode.Though to a smaller extent, the average chemical composition of theNPs is also seen to experience mm-size fluctuations during the PEMFCoperation. Experimental atomic PDFs from spots of the MEAs whereatthe loading and chemical composition of the cycled NPs exhibit largefluctuations (follow the slanted arrows in Fig. 5, S21, S22 and S23) aresummarized in Figs. S14 and S15. The PDFs are well reproduced bymodels featuring the nanophases characteristic for the cycled NPs as awhole (compare with the respective data sets shown in Figs. 3 and 4).Results indicate that, largely, it is the loading and average chemicalcomposition and not the phase type of nanoalloy particles that exhibitsignificant fluctuations over the cathode side of MEA during the PEMFCoperation. Altogether, data in Figs. 2–5, S21, S22 and S23 show that,under realistic operating conditions, Pd-Sn and Pt-Ni-Cu alloy NPs un-dergo significant atomic-level changes, including dissolution of metallicspecies, sub-Å atomic fluctuations, sharp structural transitions leadingto a partial or complete nanophase segregation and local strain re-laxation. Besides, the NPs exchange dissolved atomic species over mm-range distances and, on average, increase in size by 2–3 nm. Notably,several of the changes occur simultaneously but at significantly dif-ferent rates.

To quantify the relationship between the atomic-level changes andloss in ORR activity experienced by Pd-Sn and Pt-Ni-Cu alloy NPsduring the PEMFC operation, we summarized experimental data for thephase content, GSA and ORR activity of the NPs in Fig. 6. Experimental

Fig. 4. Nanophase analysis of fresh and cycled Pd78Sn22, Pd95Sn5 and Pt37Ni29Cu34 alloy NPs. The time of cycling and outcomes of nanophase analysis, including thepercentage of individual nanophases in the NPs, are given for each data set. Experimental PDFs (symbols) are matched with computed PDFs (red line) derived frommodels based on face-centered-cubic (FCC), orthorhombic (Orth) and tetragonal (Tetr) type nanophases. Fragments of the nanophases are given in Fig. S13. Notethat, per EDS data, the composition of cycled NPs changes (follow the black arrows in the upper part of the plots) due to leaching of metallic species during thePEMFC operation.

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data for the evolution of atomic-level strain in the NPs, as obtainedfrom changes of the respective bonding distances, are also shown in theFigure. As can be seen in the Figure, the ORR activity (effective MA) ofboth Pd-Sn and Pt-Ni-Cu NPs drops fast soon after the PEMFC startsoperating and, furthermore, the magnitude of the drop is different forthe different NPs. The drop occurs when the NPs undergo a sharp phasetransition and/or the atomic-level strain in the NPs changes fast. It canbe rationalized as follows: Based on the difference in the elemental sizeof Sn (~ 3.2 Å) and Pd (2.75 Å) [38,39], it may be expected that Sn andPd atoms in Pd-Sn alloy NPs would be under significant “compressive”and “tensile” stress, respectively. It follows then that diminishing thenumber of majority atoms in the respective NPs, e.g. through dissolu-tion, would be energetically favorable because it would reduce theexcessive stress on the minority atoms. Furthermore, Pd and Sn atomsremaining in the nanoalloys would tend to restructure in an en-ergetically favorable configuration. The restructuring would be fastbecause the interdiffusion of atomic species in nanoalloys is severalorders of magnitude faster than that in bulk alloys [40,41]. Besides, itmay trigger phase transitions into structure types characteristic of thebulk counterparts of the nanoalloys (see the phase diagram for Pd-Snsystem shown Fig. 1a). The local coordination of atoms and bondingdistances in the NPs would change accordingly. Concurrent nanophasesegregation, changes in the local atomic coordination and bondingdistances would disrupt irreversibly cooperative interactions betweenunlike atomic species in cycled Pd-Sn alloy NPs that are beneficial tothe ORR kinetics, in particular the strain and electronic effects dis-cussed above. Hence, the catalytic activity of the NPs for ORR woulddrop fast, as indeed observed here. The initial trajectory of changes inthe atomic structure and ORR activity of cycled Pt-Ni-Cu alloy NPs alsoappears to be strongly influenced by factors related to minimization ofatomic-level stresses and intra-particle diffusion because the NPs areseen to segregate into a mixture of an fcc and not so densely packedtetragonal nanophases (see PDF fits in Fig. 4), both occurring with therespective bulk alloys (see the phase diagram for Pt-TM system shownin Fig. 1b). The gradual diminishing of the ORR activity of both Pd-Snand Pt-Ni-Cu alloy NPs with the further PEMFC operation coincides

with incremental structural changes of the NPs such as a continuouslocal strain relaxation, variation in the nanophase content and ex-change of leached-off metallic species between individual NPs leadingto a substantial increase in their size. The concurrent effective masstransport over the PEMFC cathode and likely coalescing of individualNPs can be added to the list. As discussed in work of others[4,6,12,25,42], largely, such structural changes would gradually di-minish the ECSA and so overall catalytic activity of the NPs for ORR,which is exactly what our experimental data show. Altogether, data inFig. 6 indicate the presence of an instant link between the structuralchanges and ORR activity losses experienced by nanoalloy catalystsfunctioning inside a PEMFC.

To validate our findings, we studied standard (E-tek) Pt NPs strictlyfollowing the experimental protocol described above. Results aresummarized in Fig. S25. As can be seen in Figure, Pt NPs cycled forabout 5 h suffer a moderate loss in ORR activity, which may be at-tributed to the combined effect of a local strain relaxation and gradualdecrease in their GSA/ECSA. Mass transport over the PEMFC cathode isalso moderate in comparison to that observed with Pt-Ni-Cu alloy NPs(compare data in Fig. S23 and S25). Results are in line with the findingsof previous studies on the ORR activity and stability of pure Pt NPs[5,12,31,43], thus attesting to the reliability of present in operando EDSand total x-rays scattering experiments. We also studied standard Pt3Nialloy NPs, which are known to exhibit high ORR activity, this timechanging the experimental protocol. In particular, we cycled the NPs for1 h at room temperature (~ 25 °C) and then for another hour at 80 °C.Results are summarized in Fig. S24 and S26. As can be seen in theFigures, the NPs undergo significant atomic-level changes at roomtemperature, including loss in Ni species and a concurrent sharp ex-pansion of bonding distances. For reasons discussed above, it may beexpected that the ORR activity of the NPs would drop, which is what weobserve. An increase in the operating temperature of the PEMFC by55 °C results in an immediate surge in the current output of the PEMFC(by about 30%). Largely, the surge may be attributed to an improve-ment of the MEA permittivity [31]. The surge though is likely to betemporary because the catalyst particles keep growing in size and re-

Fig. 5. (first row) Mass distribution (loading) of Pd43Sn57 nanoalloy particles over a 4.5 mm×4.5mm area of the MEA obtained after the PEMFC has operated for70min, 152min and 300min. The distribution is normalized against the average catalyst loading for that area. (second row) Actual percentage of Pd over the sameMEA area for the same time periods of the PEMFC operation. Black broken lines connecting black rectangles on top of the plots emphasize the continuously changingaverage percentage of Pd during the PEMFC operation. Slanted arrows mark MEA's spots whereat the catalyst loading and chemical composition exhibit largefluctuations about the respective average values. Atomic PDFs for the spots are given in Fig. S14. Data are obtained by EDS as described in the SI.

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distributing over the PEMFC cathode. The effect of permanently in-creased thermal displacements of atoms in the catalyst particles on theirORR activity and, hence, performance of the PEMFCs in terms of cur-rent output deserves more detailed structure studies [44]. Nevertheless,results of the experiment involving a Pt3Ni nanoalloy catalyst for ORRattest to the sensitivity of in operando EDS and total x-rays scattering tospecifics of the fuel cells’ operation. Note that, given the strict tem-perature, gas flow rate and duty-cycling control exercised during thestudies of Pd-Sn and Pt-Ni-Cu alloy NPs, trends exhibited by the datasets presented in Fig. 6 appear highly reliable.

Among others, the in operando experiments indicated that the abilityof atomic species in nanoalloy particles to interdiffuse fast facilitatesstructural transformations leading to nanophase segregation and abruptchanges in bonding effects, thus inflicting losses in the nanoalloy's ac-tivity for ORR. Here we explored the possibility to bring the inter-diffusion under certain control and so reduce the related to it ORRlosses by bringing the particle's composition, structure type and mor-phology in accord with each other. In particular, a comparison betweenin operando (see Fig. S5) and in house (see Fig. S3) catalytic data in-dicated that Pd-Sn alloy NPs with a relatively low Sn content (< 25 at%) and cubic-type atomic structure are potentially excellent catalystsfor the ORR taking place inside PEMFCs. Studies of others have also

indicated that the optimal surface concentration of the non-noble metalcomponent of binary nanoalloy catalysts for ORR that is capable ofproviding oxygen species to the NM component, (NM=Pt and Pd),would be in the range of 20–30 at% [45]. Furthermore, the ORR ac-tivity of Pd surfaces is known to increase in the order Pd(110)fcc < Pd(111)fcc « Pd(100)fcc [46]. Lastly, it has been demonstrated by experi-ment that facets generic to structures of metallic atoms tend to stabilizethe structures and, hence, their physicochemical properties to a greatextent [28,47]. Thus, hinted by the results of present study and work ofothers, we synthesized Sn-poor nanoalloy particles terminated by(100)fcc facets alone, hereafter referred to as Pd-Sn nanocubes. Detailsof the synthesis of Pd-Sn nanocubes can be found in the SI. Re-presentative TEM and HR-TEM images of the nanocubes are shown inFig. 7c and d, respectively. As can be seen in the Figures, the nanocubesappeared with an average size of approximately 7.5 (± 1.5) nm andgood crystallinity. Per ICP-AES, the average chemical composition ofthe nanocubes turned out to be Pd96Sn4. The nanocubes were subjectedto an accelerated test for ORR activity and stability following the pro-tocol described above. Analysis of the experimental in operando atomicPDFs showed that, regardless having suffered loss in Pd species (~ 15 at% from the initial value; see Fig. S9), the nanocubes maintain theircubic-type atomic structure and bonding distances during the potential

Fig. 6. Color maps of the phase content of (a) Pd43Sn67, (b) Pd78Sn22, (c) Pd95Sn5 and (d) Pt37Ni29Cu34 alloy NPs vs time of electrochemical cycling. Percentage of theindividual nanophases are to be read from the vertical axis on the right. The rate of change of the atomic-level strain in the NPs is also given (arrow and symbols ingreen). For reference, the atomic-level tensile strain in fresh Pd43Sn67, Pd78Sn22 and Pd95Sn5 alloy NPs is 4.6%, 1.9%, 1.5%, respectively. The atomic-level com-pressive strain in fresh Pt37Ni29Cu34 alloy NPs is 5.3%. As discussed in the text, changes in the atomic-level strain reflect the evolution of bonding distances/effects inthe respective NPs. Also given are the retention of the initial GSA (arrow and symbols in red) and effective MA (arrow and symbols in blue) of the NPs for ORR. Theformer is evaluated from changes in the average NP’s size whereas the latter reflects changes in the current output of the PEMFC (e.g. see Fig. 2c). Note that per EDSdata, the composition of cycled NPs changes due to leaching of atomic species during the PEMFC operation (follow the horizontal black arrows on top of the plots).Vertical arrow in (c) marks the instant of PEMFC operation when the effective MA of Pd95Sn5 alloy NPs partially recovers from the initial loses and then hardlychanges with further cycling.

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cycling (see Fig. 7a). The nanocubes do not suffer significant loss inORR activity either (see Fig. S5). On the contrary, under operatingconditions, they adopt transient structural states that exhibit somewhatimproved ORR activity (see the slow oscillations in the MA, GSA andstrain values summarized in Fig. 7b). We argue that Pd-Sn nanocubespreserve their structural integrity and so ORR activity because the in-terdiffusion of atoms near their facets is reduced to hopping of atomsbetween largely equivalent positions. The lack of significant fluctua-tions in both the position and intensity of the first peak in the in oper-ando PDFs for the cycled nanocubes also indicates that they do notexperience irreversible structural changes during the PEMFC operation(see Fig. S16). A comparison between the higher-r part of the experi-mental atomic PDFs for cycled Pd96Sn4 nanocubes and sphericalPd95Sn5 NPs indicates that the surface of the former indeed remainsmore ordered at atomic level as compared to that of the latter (see Fig.

S17). This is important because uneven surfaces of nanoalloys involvingPd are known to bind ORR intermediates stronger than smooth Pdsurfaces and so underperform as catalysts for ORR [4–6,18,43]. Cur-rently, we are optimizing the ORR activity of Pd-Sn nanocubes by fine-tuning details in their synthesis and post-synthesis treatment.

It may be difficult to decide ‘a priori’ on the relative importance offactors such as reduction potential, surface energy and size of distinctatomic species, propensity to form stable nanostructures through full-scale phase transitions and/or continuous local re-alloying, particlegrowth and mass redistribution over the cells’ cathode in determiningthe type and scope of atomic-level changes that nanoalloy catalystswould indeed undergo under operating conditions [39]. The experi-mental approach introduced here may be very helpful in this respectdue to its ability to identify the major factors behind the observed ORRactivity losses without evoking model assumptions and, hence, enable

Fig. 7. (a) Exemplary nanophase analysis of fresh and cycled Pd-Sn nanocubes. Note that per EDS data, the composition of nanocubes changes (follow the horizontalarrow on top of the plot) during the PEMFC operation. (b) Rate of change of the atomic-level strain in the nanocubes during the PEMFC operation (symbols and arrowin green). For reference, the tensile atomic-level strain in the fresh nanocubes is + 1.9%. As discussed in the text, changes in the atomic-level strain reflect theevolution of bonding distances/effects in the NPs. Also given are the retention of the initial GSA (symbols and arrow in red) and the effective MA (symbols and arrowin blue) of the nanocubes for ORR. The former is evaluated from changes in the average NP’s size whereas the latter reflects changes in the current output of thePEMFC at 0.9 V on the respective polarization curves. Vertical broken lines emphasize the concurrent changes in the bonding distances, GSA and MA of thenanocubes during the PEMFC operation. The oscillatory behavior of bonding distances indicates that the nanocubes adopt some transient structural states of an fcc-type that improve their ORR activity temporarily. (c) TEM image of fresh nanocubes. The nanocubes are with an average size of approximately 7.5 (± 1.5) nm. (d)Representative HR-TEM image for a fresh nanocube. The nanocube has a very regular surface terminated with (100)fcc-type facets (see the slanted lines in red). Theinterplanar (200)fcc distance in the nanocubes is 1.98 Å. It is slightly longer than the (200)fcc distances of 1.945 Å in bulk Pd, confirming that the nanocubes are Pd-Snnanoalloys.

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the development of a rational strategy for reducing the latter throughtaking control over the former. Besides, it can bridge the knowledge gapbetween other widely used in situ techniques for atomic-level char-acterization of nanoalloy catalysts such as in situ SAXS, EXAFS andmonocrystal surface XRD. In particular, in situ SAXS provides valuableinformation for the evolution of the size and morphology of catalystparticles during potential cycling [1,5,6,48]. Data for the evolution ofthe particle's structure type obtained by PDF analysis of x-ray scatteringdata can be used to clarify the findings of SAXS experiments, as ex-emplified in Ref. [19]. In situ EXAFS provides valuable information forthe local coordination and valence electronic structure of atoms innanoalloy catalysts [5,6,15,17]. Structural characteristics of cycledcatalyst particles obtained by EDS and total x-ray scattering experi-ments, in particular chemical composition, nanophase type and degreeof local atomic disorder, can help improve the accuracy of EXAFS fittingmodels [49,50]. In situ monocrystal surface XRD provides fundamentalknowledge of catalyst surfaces [12,18]. The knowledge can be extendedto facets of actual catalyst particles and proven in practice, as demon-strated above.

4. Conclusions

In summary, combined in operando x-ray spectroscopy and totalscattering experiments conducted here evidence that, under operatingconditions, nanoalloy catalysts for the ORR taking place at the cathodeof fuel cells, in particular PEMFCs, can undergo structural changes witha largely different dynamics and/or scope. The changes can be identi-fied and quantified with atomic-level precision (~ 0.02 Å) and elementspecificity (~ 2–3 at%), and linked to the resulting ORR activity lossesin both time (~ 1min) and space (~ µm) resolved manner. This ex-perimental approach expands on the ex situ techniques currently in useby providing unique information for the evolution of the chemicalcomposition, geometric surface area, phase content, 3D structure andstrength of interactions between the constituent atoms of nanoalloycatalysts functioning under realistic operating conditions, as ex-emplified in Scheme 1. The overall catalyst utilization, related to theelemental and mass distribution of catalyst particles over the PEMFCcathode, and specifics of the PEMFC's operation such as, for example,operation temperature, can also be assessed. The knowledge can pro-vide valuable feedback on the success of the ongoing effort to engineerbetter MEAs in particular, and optimize the performance of PEMFCs ingeneral. Ultimately, combined with other widely used in situ techniquesfor structural characterization of functioning catalysts such as in situSAXS, EXAFS and monocrystal surface XRD, in operando x-ray spec-troscopy and total scattering can help transform the promise of fuel celltechnologies into an affordable reality.

Acknowledgments

This work was supported by DOE-BES Grant DE-SC0006877. Also, itused resources of the Advanced Photon Source at the Argonne NationalLaboratory provided by the DOE Office of Science under Contract No.DE-AC02-06CH11357. Thanks are due to Dr. B. Prasai for the help withHE-XRD experiments.

Notes

The authors declare no competing financial interest.

Appendix A. Supporting information

Supplementary data associated with this article can be found in theonline version at http://dx.doi.org/10.1016/j.nanoen.2018.04.049.

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Valeri Petkov is a Professor in Physics at Central MichiganUniversity. He has over 30 years of research experience inx-ray diffraction studies of materials with intrinsic struc-tural disorder, and more than 180 papers published in peer-reviewed journals.

Yazan Maswadeh is a Ph.D. candidate in the PhysicsDepartment at Central Michigan University. He is workingwith Prof. Petkov's group since 2015 and have contributedto a number of papers published in peer-reviewed journals.His current research interests include structural character-ization of nanosized materials using x-ray diffraction andspectroscopy, atomic structure-function relationship, andinvestigating electrocatalysts functioning inside fuel cells.

Yinguang Zhao received his MA degree in Chemistry fromthe State University of New York at Binghamton in 2017.His research project at Binghamton University involvedexperimental measurements of the phase structures andcatalytic activities of nanoalloy catalysts for fuel cell reac-tions.

Aolin Lu received her Ph.D. degree in Materials Physics andChemistry from Xiamen University, China in 2017. She wasa visiting Ph.D. student in Prof. Zhong's group at the Stateuniversity of New York at Binghamton. Currently, she is aPostdoctoral Fellow at Xiamen University. She has co-au-thored several papers in peer-reviewed journals. Her re-search interests include nanomaterials, electrochemistry,energy conversion and environmental catalysis.

Hannan Cronk received her Ph.D in Chemistry from theState University of New York at Binghamton in Dec 2017under the supervision of Prof. Zhong. She started recentlyher academic career at North Central State College. She hasco-authored several publications in peer-reviewed journals.Her research interests include design and synthesis of multi-metallic nanoscale materials for clean energy production,conversion and storage applications.

Fangfang Chang received her Ph.D. degree in Chemistryfrom Hunan University in 2017. Her research interests in-clude nanoscale catalyst synthesis, characterization andapplication in clean energy conversion technologies such asfuel cells. Part of her Ph.D. work was done at the StateUniversity of New York at Binghamton under the super-vision of Prof. Zhong. She started recently her academiccareer at Henan Normal University. She has co-authored anumber of peer-reviewed papers on supported nanocata-lysts for electrochemical reactions.

Shiyao Shan received his Ph.D in Chemistry at the StateUniversity of New York at Binghamton in 2016 under thesupervision of Prof. Zhong. He is currently a postdoctoralscientist. He has co-authored many peer-reviewed papersand has two patent applications pending. His research in-terests include design, synthesis, characterization and ap-plication of nanoscale materials in cutting-edge technolo-gical areas such as clean energy production, conversion andstorage, emission control, and chemical/biological sensing.

Haval Kareem is a Ph.D. candidate in the ChemistryDepartment at the State University of New York atBinghamton under the supervision of Prof. Zhong. He re-ceived his bachelor's degree of Chemistry in BaghdadUniversity in 2002. His current research interests includedesign, synthesis and characterization of robust, active, andlow-cost nanoparticle catalysts for reactions related to cleanenergy conversion technologies.

Jin Luo is a senior scientist at the State University of NewYork at Binghamton and an active entrepreneur at NSCTechnology and FlexSurface. His research interests involveinterfacial electrochemistry, catalysis, sensors, and na-noscale chemistry. Currently he has focused on developingadvanced nanomaterials for applications in fuel cells, bat-teries, chemical sensors, and biomedical devices. He has co-authored many papers in peer-reviewed journals and is aco-inventor of a number of US Patents.

Chuan-Jian Zhong is a Professor of Chemistry andMaterials Science at the State University of New York atBinghamton. He has many years of research experience ininterdisciplinary areas of materials and analytical chem-istry, focusing on design of functional nanomaterials forsustainable energy production, conversion and storage, andsensors for chemical and biological detection. He has pub-lished more than 280 peer-reviewed research articles, andis an inventor of about 20 US patents.

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Sarvjit Shastri (B.S., Caltech; Ph.D., Cornell, 1994) is aphysicist at the Advanced Photon Source at ArgonneNational Laboratory, primarily developing high-energy(50–150 keV) x-ray optics and scattering techniques. He hasalso worked on optics simulations for short-pulsed x-raysources such as SASE- and seeded-XFELS.

Peter Kenesei is a physicist (M.Sc. Eotvos University, Ph.D.Eotvos University, 2010) at the Materials Physics andEngineering Group at the Advanced Photon Source atArgonne National Laboratory. He works on the develop-ment of high-energy x-ray diffraction microscopy and mi-crotomography techniques. He has co-authored 80 pub-lications in peer-reviewed journals. His current researchinterests include elastic and plastic deformation, damageand fatigue of complex material systems (composites,multilayers, biomaterials, batteries, fuel cells), high preci-sion metrology, and computational materials science.

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