Beyond Strain and Ligand Effects: Microstrain-Induced Enhancement

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Beyond Strain and Ligand Effects: Microstrain-Induced Enhancement of the Oxygen Reduction Reaction Kinetics on Various PtNi/C Nanostructures Raphael Chattot, Tristan Asset, Pierre Bordet, Jakub Drnec, Laetitia Dubau, and Frédéric Maillard ACS Catal., Just Accepted Manuscript • DOI: 10.1021/acscatal.6b02356 • Publication Date (Web): 21 Nov 2016 Downloaded from http://pubs.acs.org on November 22, 2016

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Beyond Strain and Ligand Effects: MicrostrainInduced Enhancement of the Oxygen Reduction Reaction Kinetics on Various PtNi/C Nanostructures Raphaël Chattot †, ‡, *, Tristan Asset †, ‡, Pierre Bordet ⊥, ς, Jakub Drnec ψ, Laetitia Dubau †, ‡ and Frédéric Maillard †, ‡,* † Univ. Grenoble Alpes, LEPMI, F-38000 Grenoble, France ‡ CNRS, LEPMI, F-38000 Grenoble, France ⊥ Université Grenoble Alpes, Institut Néel, F-38000 Grenoble, France

ς CNRS, Institut Néel, F-38000 Grenoble, France

ψ European Synchrotron Radiation Facility, ID 31 Beamline, BP 220, F-38043 Grenoble Cedex, France

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ABSTRACT

The electrical performance of a proton exchange membrane fuel cell (PEMFC) is limited by the slow oxygen reduction reaction (ORR) kinetics. Catalytic improvements for the ORR have been obtained on alloyed PtM/C or M-rich core@Pt-rich shell catalysts (where M is an early or late transition metal) compared to pure Pt/C, due to a combination of strain and ligand effects. However, the effect of the fine nanostructure on the ORR kinetics remains under investigated. In this study, nanometre-sized PtNi/C electrocatalysts with low Ni content (~15 at. %) but different nanostructures and different densities of grain boundary were synthesized: solid, hollow or “sea sponge” PtNi/C nanoalloys, and solid Ni-core@Pt-shell/C nanoparticles. These nanostructures were characterized by transmission and scanning-transmission electron microscopy (TEM and STEM, respectively), X-ray energy dispersive spectroscopy (X-EDS) Synchrotron wide-angle X-ray scattering (WAXS), atomic absorption spectroscopy (AAS) and electrochemical techniques. Their electrocatalytic activities for the ORR were determined, and structure-activity relationships established. The results showed that (i) the compression of the Pt lattice by ca. 15 at. % Ni provides mild ORR activity enhancement compared to pure Pt/C, (ii) highly defective PtNi/C nanostructures feature up to 9.3-fold enhancement of the ORR specific activity over a commercial Pt/C material with similar crystallite size, (iii) the enhancement in ORR catalytic activity can be ascribed to the presence of structural defects as shown by two independent parameters: the microstrain determined from WAXS, and the average COads electrooxidation potential (μ ) determined from COads stripping measurements. This work indicates that, at fixed Ni at. %, ORR activity can be tuned by nanostructuring and suggests that targeting structural disorder is a promising approach to improve the electrocatalytic properties of mono or bimetallic nanocatalysts.


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KEYWORDS: proton exchange membrane fuel cell; oxygen reduction reaction; structural defects; grain boundary; microstrain.

INTRODUCTION

The United Nations Climate Change Conference (COP 21 or CMP 11) held in Paris in December 2015 has expressed the willingness of 195 states to lower their greenhouse gas emissions (mostly carbon dioxide and methane), increase their energy efficiency and the ratio of renewable sources in their energy portfolio. The transportation sector is crucial for the development of a carbon-free landscape 1. Hence considerable efforts are devoted worldwide to develop “electromobility” as shown by the 16 billion US dollars investment in infrastructures, fiscal incentives and research and development in favour of electric vehicles realized by United States, United Kingdom, China, Japan, Germany and France between 2008 and 2014 2. Low-temperature fuel cells, in particular proton exchange membrane fuel cells (PEMFC) using hydrogen play a key role in this energetic transition. Indeed, the high energetic density of hydrogen (33 kWh kg-1) and the ability of PEMFC to be refuelled (instead of being recharged as it is the case for battery-based vehicles) make them particularly interesting for automotive application 3,4.

However, the high cost and the limited durability of PEMFC systems still hinder their widespread development

5,6

. The high cost of PEMFC electrodes is mainly due to the large

quantity of platinum (Pt) needed to activate the slow oxygen reduction reaction (ORR) kinetics at the cathode 7. The density functional theory (DFT) calculations made by Nørskov et al.

8

on Pt(111) have shown that the origin of the large ORR overpotential is the strong

adsorption of oxygenated ORR intermediates, which block protons and electrons transfer. 3 ACS Paragon Plus Environment

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Since the interactions between a metal surface and oxygenated species depend on the d-band centre (ɛd) and of its degree of filling, an efficient strategy to tune the electronic structure and reactivity of Pt atoms consists in mixing with M atoms due to strain

9–13

, and ligand

14,15

effects (M is an early or late transition metal). A 10-fold and a 90-fold enhancement of the ORR kinetics were obtained on Pt3Ni(111) single crystal surface compared to Pt(111) or commercial nanometre-sized Pt/C, respectively

16

. However, transposing this catalytic

enhancement on nanomaterials remains highly challenging.

Surface-segregated type architectures hold promises to meet the electrocatalytic activity and stability requirement of PEMFC cathodes. These nanostructures are characterized by an alloyed core surrounded by a Pt-rich shell, which prevents leaching of transition metal atoms in the PEMFC environment. Besides, they offer core composition, Pt shell thickness, particle size and shape as tuneable parameters for catalytic activity improvement. Using (electro)chemical dealloying, various surface segregated PtxNi1-x nanocatalysts were recently prepared,

17

among which dealloyed PtNi3 displayed the highest ORR activity and stability,

both in rotating disk electrode (RDE) experiments and PEMFC systems 18. Intensive studies (see

19

and references therein) revealed that: (i) the compression of the Pt lattice induced by

the M element controls the ORR activity 20, (ii) considerable surface strain is only possible via high M molar ratio (about 70 at. %) in the bulk of the nanomaterial, (iii) surface strain is controlled by the distribution of M up to 10 atomic near surface layers from the surface but M dissolve from the first two-three monolayers

17,21

, (iv) nanoparticles exceeding 10-13 nm in

size are ineluctably subjected to nanoporosity formation during the dealloying process or during PEMFC operation due to incapacity of the Pt-rich shell to prevent the leaching of M atoms 22–25, (v) the loss in M atoms increases the Pt specific surface area but is insufficient to compensate for the ORR activity drop due to Pt-lattice relaxation

22

. Similar to alloys or 4

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segregated nanostructures, a fine tuning of the electrocatalytic activity of Pt atoms is possible on core@shell nanoparticles

26

. Parameters such as the nature and the structure of the core

material are critical to catalytic improvement 27. A maximal usage of Pt atoms is theoretically possible by decreasing the Pt shell to a monolayer but remains hardly feasible,

28

even when

using the under potential deposition of Cu as an intermediate deposition step 29. Indeed, the underlying metal core is incompletely covered with Pt atoms yielding corrosion and spontaneous evolution to hollow nanoparticles under acidic conditions

30–35

. Interestingly,

despite its pivotal fundamental and applied importance and because of the abundant variety of electrochemical protocols used worldwide,

36,37

a direct comparison of the ORR kinetics on

these different catalysts architectures has never been achieved, and is the focus of the present study.

Herein, we synthetized and characterized nanometre-sized PtNi/C electrocatalysts with similar crystallite size (~ 3 nm), Ni content (~15 at. %) but different nanostructures: solid, hollow or “sea sponge” PtNi/C nanoalloys, and solid Ni-core@Pt-shell/C nanoparticles. A 9.3-fold enhancement of the ORR kinetics was measured on the most active bimetallic catalyst. Using physical, chemical and electrochemical techniques, we correlate these changes in ORR electrocatalysis to two independent parameters: the microstrain determined from Synchrotron wide-angle X-ray scattering (WAXS), and the average COads electrooxidation potential (μ ) determined from COads stripping measurements. The results show that, at fixed Ni at. %., ORR activity can be tuned by nanostructuring

RESULTS AND DISCUSSION Different experimental procedures were used to synthesize Ni core@Pt shell, solid or hollow alloy, or sea-sponge PtNi/C nanoparticles. 5 ACS Paragon Plus Environment

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Synthesis and Characterization of Solid Alloy PtNi/C Nanoparticles The polyol synthesis is a method of choice to prepare PtNi/C alloy nanoparticles with controlled size and chemical composition

38,39

. In this method, the PtNi/C nanoparticles are

formed by the co-reduction of dissolved Pt and Ni salts in the presence of a high surface area carbon support (here Vulcan XC72) dispersed in ethylene glycol (EG). EG acts as solvent, stabilizer and reducing agent (the oxidation of EG into glycolic acid (GA) produces electrons, which are used to reduce the metal salt precursors). Controlling the pH at a value of 9-10 allows deprotonation of GA, formation of glycolate ions and CO molecules which prevent nanoparticle agglomeration

40

. Despite the stabilizing properties of glycolate ions, we

observed that the co-reduction of Ni atoms and Pt atoms resulted into aggregated PtNi/C nanoparticles, in agreement with former literature reports

41

. Using mono-sodium citrate as

surfactant in the initial EG solution, monodisperse PtNi/C particles could however be synthesized. Figure 1 displays TEM images and particle size distribution for the monodisperse or agglomerated PtNi/C nanomaterials synthesized in this study and the reference Pt/C material (provided by Tanaka Kikinzoku, TKK), denoted as PtNi/C, A-PtNi/C, and Pt/C TKK, respectively.


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Figure 1. TEM images, particle size distributions and graphical representations (insert) of (a) the reference Pt/C TKK, (b) isolated alloy PtNi/C (PtNi/C) and (c) agglomerated alloy PtNi/C (A-PtNi/C) nanoparticles. The particle size distributions for each catalyst were built by measuring the diameter of more than 200 isolated (i.e. not-agglomerated) nanoparticles.

Synthesis and Characterization of Ni@Pt Nanoparticles Various synthesis routes have been developed to produce core@shell nanostructures

42

.

However, most of them (including (electro)chemical dealloying, adsorbate-induced 7 ACS Paragon Plus Environment

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segregation and thermal induced segregation from a PtM alloy) produce M-rich core covered by a Pt-rich shell instead of two distinct phases. Herein, to produce pure Ni-core@Pt-shell nanoparticles, a sequential polyol method was used 43. Briefly, Ni nanoparticles supported on Vulcan XC72 were first prepared by the polyol method in the presence of mono-sodium citrate. The Ni/C nanoparticles were then used as template for the deposition of Pt via the polyol route, resulting in the formation of Ni-core@Pt-shell nanoparticles.

Figure 2. TEM images, particle size distributions and graphical representations (insert) of the (a) Ni/C template and (b) Ni-core@Pt-shell/C nanoparticles prepared via two sequential polyol syntheses. The particle size distributions of each catalyst were built by measuring the diameter of more than 200 isolated (i.e. not-agglomerated) nanoparticles. The determination of the core-shell nature of a bimetallic catalyst is highly challenging

44

.

However, we obtained convincing evidences that Ni-core@Pt-shell nanoparticles were synthesized. First, the increase in particle size from the initial Ni/C template to the final 8 ACS Paragon Plus Environment

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Ni@Pt/C nanoparticles (Figure 2) strongly suggests that Pt4+ ions were mainly chemically reduced on Ni/C. Second, X-EDS measurements revealed presence of Ni atoms in the final Ni@Pt/C nanoparticles ( Table 1). Table 1. Structural and chemical parameters for the electrocatalysts evaluated in this study. Atomic composition determined by X-EDS, Pt weight fraction determined by AAS, lattice parameter and crystallite size determined from Rietveld refinement of the Synchrotron WAXS patterns. X-EDS Electrocatalyst

At. (%)

AAS Comp Pt

Synchrotron WAXS Lattice parameter (Å)

Contraction vs. Pt/C

Crystallite size

(%)

( )

20.0*

3.918

0

1.3 ± 0.2

weight fraction (wt.%)

Pt/C TKK

Pt Ni

PtNi/C

Pt ± Ni ±

16.9 ± 0.2

3.885

0.8

1.2 ± 0.2

A-PtNi/C

Pt ± Ni ±

19.7 ± 0.1

3.862

1.4

4 ± 0.5

Ni@Pt/C

Pt ± Ni±

19.4 ± 0.2

3.915

0.1

1.8 ± 0.2

Hollow PtNi/C

Pt ± Ni±

22.6 ± 0.1

3.870

1.2

2.7 ± 0.4

Sea sponge PtNi/C

Pt ± Ni±

16.9 ± 0.2

3.820

2.5

4.5 ± 0.9

* The weight fraction of the reference Pt/C material was guaranteed by TKK.

Considering the last step of the polyol synthesis (acidification of the solvent for deposition of the nanoparticles on the carbon support, see Experimental Section), this is clear evidence of complete coverage of the Ni cores with Pt. Note however that typical reflections of the small Ni core could not be detected by WAXS with Synchrotron radiation or pair distribution function (PDF) (see Figure S1 in the Supporting Information, where the distance typical of 9 ACS Paragon Plus Environment

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Ni-Ni pairs at r = 2.5 Å was not detected). This was attributed to the fact that the scattering amplitude of X-rays increases with the product of the atomic numbers of the atoms constituting the pair. Hence, at a given Pt:Ni stoichiometry, the intensity of the Pt-Pt pairs is 7 times higher than that of the Ni-Ni pairs. Obviously, this trend is magnified by the low Ni content targeted in this study (~15 at.%). Lastly, despite their larger crystallite and particle sizes, the Ni@Pt/C nanoparticles exhibited a higher Pt specific surface area compared to Pt/C TKK (see Structure – ORR Catalytic Activity Relationships Section): this is convincing evidence of the better utilization of Pt atoms in the Ni-core@Pt-shell nanostructure.

Because the chemical composition and the size of a core@shell particle are linked by the shell thickness, we could assess the number of deposited Pt monolayers on the template Ni cores. By combining the mean particle size established from a statistical analysis of TEM images, the chemical composition provided by X-EDS measurements, and the geometric model introduced by Montejano-Carrizales, 45,46 the Pt shell was estimated to be ca. 0.9 ± 0.2 nm thick (ca. 4.5 ± 1 monolayers, see Figure S2 in the Supporting Information).

Synthesis and Characterization of Porous Hollow PtNi/C Nanoparticles Porous hollow PtNi/C nanoparticles were synthetized via a one-pot synthesis first reported by Bae et al. 47. In this synthesis, Ni-rich/C nanoparticles are first formed in the reactor, and then act as a sacrificial template for the deposition of Pt2+ ions via the combined actions of galvanic displacement and chemical reduction. The resulting gradient in chemical potential in the Ni-rich core@Pt-rich shell and the difference in oxophilicity between Pt and Ni atoms then provide two strong driving forces for the evolution from a core@shell to an alloy nanostructure. However, due the faster interdiffusion coefficient of Ni atoms compared to that of Pt atoms in the PtNi alloy, a flux of vacancies enters the nanomaterial causing the 10 ACS Paragon Plus Environment

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formation of nanometre-sized voids (a phenomenon which is known as the Kirkendall effect 48–53

). During the last step of the synthesis, any residual Ni atom exposed to electrolyte is

removed via acid leaching in 1 M H2SO4 54. The final hollow PtNi/C nanoparticles are highly polycristalline, as shown by high-resolution transmission electron microscopy (HR-TEM) images of Figure 3b.

Figure 3. (a) High angle annular dark field scanning transmission electron microscopy (HAADF-STEM) image and X-EDS elemental maps, (b) TEM, particle size distribution and (c) graphical representation of hollow PtNi/C electrocatalysts. The particle size distribution was built by measuring the maximum Feret diameter of the outer metal shell and of the inner core on more than 200 isolated (i.e. not-agglomerated) nanoparticles.

Synthesis and Characterization of Porous ‘Sea Sponge’ PtNi/C Nanoparticles To produce ‘sea sponge’ PtNi/C nanostructures, the polyol process described previously was slightly modified: a Pt2+ rather than a Pt4+ precursor was chosen, and the initial Ni:Pt 11 ACS Paragon Plus Environment

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precursor ratio was fixed as 3:1. As shown by Figure 4b, the resulting “sea sponge” nanoparticles feature open porosity (see also scanning electron microscopy images in Figure S3 and more STEM images coupled with X-EDS elemental maps in Figure S4 in the Supporting Information). The open porous structure seen in Figure 4b is believed to result from the massive dissolution of Ni atoms during the acidic post-treatment (see Experimental Section).

Figure 4. (a) STEM images and X-EDS elemental map, (b) TEM, (c) particle size distribution and graphical representation of the “sea sponge” PtNi/C electrocatalyst synthesized in this study. The particle size distribution was built by measuring the diameter of more than 200 isolated (i.e. not-agglomerated) nanoparticles.

Structure and Chemical Composition of the Different Electrocatalysts The chemical composition of the synthesized bimetallic catalysts was determined by XEDS and the Pt weight fraction by atomic absorption spectroscopy (AAS). As shown by 12 ACS Paragon Plus Environment

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Table 1, all catalysts feature a Ni content ca. 10-15 at. % and a Pt weight fraction close to 20 % in agreement with our initial objectives. The lattice parameter and the size of the PtNi/C nanocrystallites were obtained by Rietveld refinement of the WAXS patterns (the parameters used for the analysis, and an example of refinement are provided Table S1 and Figure S5 respectively in the Supporting Information).

The Synchrotron WAXS patterns measured on all catalysts are presented in Figure 5. The shift of the Bragg peaks to higher 2θ values compared to Pt/C TKK is clear evidence of Pt lattice contraction for the bimetallic catalysts. However, the degree of contraction of the Pt lattice (lattice strain) strongly depends on the fine arrangement of Pt and Ni atoms, being small for core@shell nanostructure and large for open porous nanostructures, as shown in Figure S6 in the Supporting Information.

Figure 5. Synchrotron wide angle X-ray scattering patterns of the synthetized bimetallic catalysts and of the reference Pt/C TKK material. 13 ACS Paragon Plus Environment

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The measured crystallite sizes consistently agree with the particle sizes or the Pt-rich shell thicknesses estimated by TEM or STEM on the Pt/C TKK, PtNi/C and Ni@Pt/C nanomaterials: 2 ± 0.5 nm, 2 ± 0.5 nm, 0.8 ± 0.2 nm, respectively. The marked difference between crystallite and particle sizes for the hollow and “sea sponge” PtNi/C nanostructures (11.4 ± 2.0 nm and , 61 ± 21 nm, respectively) illustrates the fact that these nanomaterials are composed of individual nanocrystallites interconnected with each other via grain boundaries.

In addition to the lattice strain (that is the degree of contraction of the lattice parameter deduced from the shift in the position of the WAXS peaks by comparison to Pt/C), Rietveld analysis was also used to extract the microstrain (that is the spread of the lattice parameter values around the mean)

55

. Microstrain, also referred as local lattice strain,

56

is caused by

the deviation of the atoms from their ideal positions due to the presence of structural defects (stacking faults, twins, grain boundaries, and or dislocation arrays) 57. Microstrain produces a broadening of the WAXS peaks without any change in the peaks position. Since both crystallite size and microstrain contribute to broadening of the WAXS peaks, analytical methods are generally employed to separate the two contributions. Note that the superiority of the Rietveld refinement method over others (such as the frequently employed WilliamsonHall analysis) has recently been evidenced by Maniammal et al. 55 for ~3 nm crystallites.

The values of microstrain (denoted as ‘raw microstrain’) are presented in Figure 6 and Table 2. In general, higher values are found for alloyed and polycrystalline/aggregated nanoparticles (sea sponge PtNi/C, hollow PtNi/C and A-PtNi/C) compared to monodisperse and unalloyed nanoparticles (Pt/C and Ni@Pt/C). One exception is the monodispersed PtNi/C catalyst, which exhibits a high raw microstrain value. However, this result is not surprising if 14 ACS Paragon Plus Environment

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one considers that microstrain depends on the crystallite size

58,59

(the microstrain increases

with decreasing the crystallite size) and that the monodispersed PtNi/C catalyst features very small nanocrystallites (1.2 nm vs. 2.9 nm in average for other samples).

Figure 6. Values of raw microstrain (left scale) and microstrain corrected from surface effect (“Corrected Microstrain”, see right scale) for the catalysts evaluated in this study. The size dependency of the raw microstrain is problematic for catalytic aspects. Indeed, the raw microstrain (a bulk property) overestimates the atomic disorder caused by a structural defect present at the surface of a small crystallite simply because of the high atomic surface/volume ratio. On the contrary, the disorder caused by a structural defect present at the surface of a large crystallite is compensated by a larger ordered core. To better take into account the size dependency of the raw microstrain, a corrected microstrain was obtained using the dispersion (that is the ratio between surface and bulk atoms) as:

! "#$ %#& '%) )

*+, -./01230+.4 '%%) 5

(1) 15


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Where D is the dispersion or the “Surface Atoms Ratio”, as defined by MontejanoCarrizales et al. in their model 45,46:

6 ) 100 x

89 8

(2)

And N and Ns are the total number of atoms and the number of surface atoms, respectively:

: ) 10

;<

+ 5" +

"+1

:s ) 10" + 2

(3)

(4)

“m” corresponds to the number of atomic layers composing the crystallite in the Montejano-Carrizales model. The values of “m” are linked to the crystallite size “d” size by the relation:

")

@

(5)

√0BC

With “rPt” the covalent radius of a Pt atom (rPt = 0.135 nm)

Table 2. Structural parameters (as-determined and corrected microstrain values, dispersion calculated from the mean crystallite size) determined by Rietveld refinement of the WAXS patterns on the different PtNi nanocatalysts and the reference Pt/C TKK material.


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Microstrain

Dispersion

(%%)

(%)

Corrected Microstrain (%)

Pt/C TKK

0

65 ± 8

0

PtNi/C

160.0

68 ± 8

2.4 ± 0.1

A-PtNi/C

90.0

30 ± 4

3 ± 0.2

Ni@Pt/C

28.7

54 ± 6

0.5 ± 0.1

Hollow PtNi/C

136

40 ± 6

3.4 ± 0.3

Sea sponge PtNi/C

233.3

28 ± 6

8.7 ± 1

Electrocatalyst

As shown in Figure 6 and Table 2, even if the overall microstrain values decrease after correction, the only change resulting from this correction occurs for isolated and agglomerated PtNi/C nanoparticles. Hence, corrected microstrain values better reflect the fact that the A-PtNi/C catalyst is composed of individual PtNi/C nanocrystallites necking via grain boundaries. Note also that the values of the corrected microstrain estimated on nanostructures exhibiting pure Pt phases - Ni@Pt/C and Pt/C TKK - are by far smaller than those of alloyed nanostructures. The absence of microstrain for the reference Pt/C TKK is particularly striking. However, we stress that similar values were obtained on a Pt/C catalyst manufactured by ETEK (raw microstrain: 17 %%; dispersion: 59 %; corrected microstrain: 0.29 %). A contrario, the nanostructures composed of aggregated crystallites featured the higher microstrain values (A-PtNi/C < hollow PtNi/C < sea sponge PtNi/C). This is believed to reflect the high density of structural defects (in particular of grain boundaries) found in these nanostructures.

Structure – ORR Catalytic Activity Relationships We then investigated the surface reactivity and the ORR activity of the different nanostructures. Prior to any catalytic activity measurement, all synthesized and reference nanocatalysts were introduced at E = 0.1 V vs. the reversible hydrogen electrode (RHE) in the 17 ACS Paragon Plus Environment

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electrochemical cell and “activated” by a procedure derived from the recommendations of Shinozaki et al.

37

(50 potential cycles at 0.5 V s-1 between 0.05 and 1.23 V vs. RHE in Ar-

saturated 0.1 M HClO4, in order to obtain a stable surface state). The first and the last “activation” cycles in the potential region 0.05 < E < 0.50 V vs. RHE measured on the different catalysts are shown in Figure 7 (the cycles in the potential region 0.05 < E < 0.50 V vs. RHE can be found in Figure S7 in the Supporting Information).

A strong dependence of the charge density associated with the desorption of underpotentially deposited H (Hupd) on the Ni content in the surface and near-surface region was found, in agreement with literature

60,61

. The presence of sharped Hupd peaks before and after

electrochemical activation for Pt/C TKK and Ni@Pt/C confirms that they feature a pure Pt surface (“Pt to Pt”). In opposite, sharp Hupd peaks developed on PtNi/C and A-PtNi/C during electrochemical activation, which indicates that Ni atoms were leached out from the surface and near-surface region during this procedure resulting in the formation of a dealloyed surface (“PtNi to Pt”). Finally, on the hollow and “sea sponge” PtNi/C nanoparticles, a small bump at E = 0.3 V vs. RHE suggested the presence of (100) type facets (Figure 7e and Figure 7f): this feature nicely agrees with their highly defective nanostructure, namely the presence of grain boundaries at the surface of these highly disordered nanocatalysts (“unconventional Pt”).


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Figure 7. Comparison of the Hupd desorption region before (1st scan) and after (50th scan) electrochemical activation for (a) the reference Pt/C TKK, (b) PtNi/C, (c) Ni@Pt/C, (d) APtNi/C, (e) hollow PtNi/C, and (f) “sea-sponge” PtNi/C (f) catalysts. Ar-saturated 0.1 M HClO4; v = 0.500 V s−1; T = 298 ± 1 K; no rotation of the electrode; Pt loading: 20 µgPt cmI FGH . COads stripping measurements were performed to gain further insights into the fine nanostructure of the synthesized nanomaterials. Indeed, the COads electrooxidation is a highly 19 ACS Paragon Plus Environment

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structure sensitive reaction and many particle properties such as surface composition

62–64

,

particle size 65,66, particle agglomeration 66,67 and crystallographic orientation of the facets 68,69 influence the number, the position and the shape of the COads stripping peaks. The COads stripping voltammograms recorded on the synthesized materials feature two different contributions (Figure 8c). In agreement with former literature findings,

65–67,70–72

the first

peak located between 0.40 < E < 0.73 V vs. RHE was ascribed to COads electrooxidation on individual nanocrystallites interconnected via grain boundaries, and the second peak between located 0.7 < E < 1.0 V vs. RHE was ascribed to COads stripping on isolated (i.e. non agglomerated) nanocrystallites. This is unequivocally shown by Figure 8c: the COads stripping voltammograms are dominated by the high potential peak on monodispersed catalysts (Pt/C TKK, PtNi/C and Ni@Pt/C), and by the low potential peak on highly defective nanocatalysts (hollow and on ‘sea sponge’ PtNi/C). The catalyst A-PtNi/C, which contains both isolated and agglomerated nanocrystallites, represents an intermediate case.


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Figure 8. (a) Base and (c) background-subtracted COads stripping voltammograms measured on the nanomaterials evaluated in this work, (b) Tafel plots obtained from the steady-state I−E curves measured at ω = 1600 rpm after correction for Ohmic losses and diffusion in solution and (d) ORR specific activity measured at E = 0.95 V vs. RHE for the electrocatalysts evaluated in this work. (a) and (c) were conducted in Ar-saturated 0.1 M HClO4 at v = 0.020 V s−1 without rotation of the electrode. (b) and (d) were conducted in O2-saturated 0.1 M HClO4 at a potential sweep rate v = 0.005 V s−1. Other conditions: T = 298 ± 1 K, Pt loading = 3.92 µg, ω = 1600 rpm. The error bars are the standard deviation of at least three independent measurements.


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It is striking that both COads stripping voltammograms (electrochemical measurement) and the corrected microstrain values derived from WAXS patterns (physical measurement) both suggest that the density of structural defects (such as grain boundaries) increased in the order: Pt/C TKK, PtNi/C, Ni@Pt/C < A-PtNi/C < hollow PtNi/C < sea sponge PtNi/C. To extract quantitative information, the COads electrooxidation voltammograms were statistically analysed. The potential weight in COads stripping voltammograms J'K ) in V I was defined as:

M . J'K )!K ) 1

(6)

Yielding:

J'K) )

N'O)

(7)

P MQ.R N'O)@O

Where S'K) is the measured current density (µA cmI TU ) at a given electrode potential E (here comprised between 0.4 – 1.0 V). Therefore, the average COads oxidation potential μ (in V) corresponds to the first moment of the potential weight:

μ ) M . K. J'K )!K

(8)

The values of μ determined for each nanostructure are displayed in Table 3. Despite the approximations inherent to this statistical parameter (which was applied for both unimodal and bimodal COads oxidation current distribution), the relation between COads electrooxidation 22 ACS Paragon Plus Environment

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kinetics and the grain boundaries content

65–67

resulted into low μ values on highly

defective nanomaterials (hollow and sea-sponge PtNi/C nanoparticles). On the contrary, higher values of μ were found for pure Pt phases (Pt/C TKK and Ni-core@Pt-shell). Table 3. Pt specific surface area (SPt,CO), average COads electrooxidation potential (μVW ) and iR + mass transport corrected ORR specific activity (SA0.95) and mass activity (MA0.95) measured at E = 0.95 V vs. RHE on the PtNi/C nanostructures and the reference Pt/C TKK, and corresponding enhancement factor (E.F.). The ORR activity was determined in O2saturated 0.1 M HClO4 at a potential sweep rate v = 0.005 V s−1. Other conditions: T = 298 ± 1 K, Pt loading = 3.92 µg, ω = 1600 rpm.

SPt,CO Electrocatalyst

μ]^ \ vs. RHE

SA0.95

E. F.

MA0.95

E. F.

IX 'μ_ ` Z[ )

SA

'_ Y I\ Z[ )

MA

' X Y I\ Z[ )

(V)

Pt/C TKK

78 ± 5

0.82

27 ± 4

1

21 ± 5

1

PtNi/C

104 ± 12

0.81

24 ± 1

0.9

25 ± 2

1.2

A-PtNi/C

39 ± 1

0.76

76 ± 1

2.8

30 ± 1

1.4

Ni@Pt/C

95 ± 6

0.82

36 ± 2

1.3

34 ± 2

1.6

Hollow PtNi/C

41 ± 4

0.72

151 ± 6

5.6

62 ± 8

3

Sea sponge PtNi/C

46 ± 4

0.66

250 ± 32

9.3

122 ± 19

5.8

The ORR activity of the various nanostructures was investigated according to the electrochemical protocol defined by Kocha et al. 37,73–77. The Tafel plots and the ORR specific activity determined at E = 0.95 V vs. RHE are presented in Figure 8, and the ORR mass activity in Table 3 (for the sake of comparison with other studies, the specific and the mass activity for the ORR were also calculated at E = 0.9 V vs. RHE (SA0.90 and MA0.90, respectively) and are displayed in Table S2). A slight improvement of the ORR specific 23 ACS Paragon Plus Environment

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activity compared to pure Pt/C TKK was found for the bimetallic nanostructures composed of isolated nanoparticles (PtNi/C and Ni@Pt/C). This slight improvement can be rationalized based on the mild Pt lattice compression, and the similar Pt-rich surface for Pt/C TKK, PtNi/C and Ni@Pt/C electrocatalysts (Figure 7b, Figure 7c, and Figure 7d). The ORR specific activity was significantly enhanced on nanostructures composed of aggregated PtNi/C nanocrystallites: a 2.8, 5.8 and 9.3-fold enhancement of the ORR specific activity was monitored on A-PtNi/C, hollow PtNi/C and sea sponge PtNi/C, respectively compared to the reference Pt/C TKK material. This is clear evidence that, at fixed Ni at. %, structural defects (such as grain boundaries) significantly enhance the ORR kinetics.

To confirm our hypothesis, the ORR specific activity was plotted as a function of two parameters which are reflective of the fraction of structural defects: the corrected microstrain and the average COads oxidation potential, μ (note that plotting the ORR specific activity as a function of the raw microstrain produced similar result, see Figure S8 of the Supporting Information). The nearly linear structure – ORR activity relationships displayed in Figure 9 unambiguously shows that, at low Ni at. %, ORR activities can be tuned by nanostructuring.


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Figure 9. Relationships between the ORR specific activity measured at E = 0.95 V vs. RHE and the density of structural defects for the different nanostructures evaluated in this study. The density of structural defects was estimated by two independent parameters: (a) the average COads oxidation potential and (b) the corrected microstrain. The ORR specific activities were extracted from ORR voltammograms conducted in O2-saturated 0.1 M HClO4 at a potential sweep rate v = 0.005 V s−1 after correction from Ohmic losses and diffusion in solution. Other conditions: T = 298 ± 1 K, Pt loading = 3.92 µg, ω = 1600 rpm.

The results of this study confirm the conclusions of References

35,78

where the beneficial

effect of structural defects on the ORR activity was first observed on hollow PtNi/C nanoparticles. However, they represent a significant step forward since the concept of microstrain-induced enhancement of the ORR activity is, for the first time, quantitatively extended to six different catalysts featuring different atomic arrangement (isolated or aggregated PtNi/C nanoparticles, solid or hollow alloy PtNi/C nanoparticles, Ni-core@Ptshell nanoparticles). This catalytic trend is in line with the recent ab initio calculations of Calle-Vallejo et al.

79

. Using density functional theory (DFT) and experiments on single

crystal surfaces, the authors evidenced that catalytic sites with generalized coordination numbers greater than 7.5 feature superior ORR activity (the calculated optimum being 8.3). Hence, the results presented in this study provide evidence that the approach of Calle-Vallejo et al. is transposable on nanocatalysts featuring different nanostructure such as core@shell or solid/hollow nanoalloys. They show that, by using grain boundaries as a proxy to implement catalytic sites with reduced atomic density and thus higher generalized coordination numbers, the ORR specific activity of nanocatalysts can be improved by a factor of more than 9 while keeping the PtNi stoichiometry identical. Considering the above arguments, the results also suggest that the remarkable ORR activity measured on highly polycristalline nanomaterials, 25 ACS Paragon Plus Environment

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such as PtM nanowires

80–82

or PtM aerogels

83,84

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(which display a COads electrooxidation

peak comprised between 0.4-0.73 V vs. RHE in comparable experimental conditions of the present work) may be due to a combination of structural, strain and ligand effects. Hence, implementing structural disorder in nanomaterials appears to be a promising route to enhance the ORR activity of bimetallic electrocatalysts.

CONCLUSION In this study, Pt-Ni nanoparticles with similar crystallite size and chemical composition (Pt:Ni stoichiometry ~ 85:15) but various nanostructures were synthetized and characterized. By combining STEM-X-EDS elemental maps, Synchrotron WAXS and electrochemical characterizations, two independent structural parameters (corrected microstrain and average COads oxidation potential, μ ) were extracted and used to bridge the ORR activity and the fraction of structural defects of a given catalytic nanomaterial (in particular its fraction of grain boundaries). The results indicate that ORR activity can be tuned by nanostructuring, and suggests that targeting structural disorder is a promising approach to improve the electrocatalytic properties of mono or bimetallic nanocatalysts.

MATERIALS AND METHODS Reference Electrocatalyst A Pt/Vulcan XC72 catalyst with a Pt weight fraction (wt. %) of 20% (TEC10V20E) was purchased from TKK and used as reference material without any treatment.

Synthesis of Monodisperse Solid PtNi/C Nanoparticles To synthesize monodisperse solid PtNi/C nanoparticles, 53 mg of H2PtCl6.6H2O (99.9 % metals basis, Alfa Aesar) and 24 mg NiCl2.6H20 metal salts (Putratronic® 99.9995 %, Alfa 26 ACS Paragon Plus Environment

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Aesar) and 118mg of monosodium citrate (≥99 %, Roth) were first dissolved in a vial containing a 20 mL mixture of deionized water and ethylene glycol (EG) (Rotipuran® ≥ 99.9 %, Roth) (vol. ratio 1:1). Then, 74 mg of carbon black (Vulcan XC72, Cabot) was dispersed by sonication in a separated vial containing also 20 mL of 1:1 EG:water mixture. Then, the contents of each vial was mixed in 20 mL of pure EG, leading to a 2:1 EG:water volumic ratio. The pH of the obtained mixture was adjusted to 10 using a 0.5 M NaOH (Suprapur® 99.99 %, Merck) solution (diluted in 1:1 EG:water mixture). The resulting suspension was afterwards kept under vigorous stirring for 1 hour under argon atmosphere before being refluxed at T = 433 K for 3 hours. The solution was allowed to cool down to room temperature under air atmosphere for 12 hours with continuous stirring

85

. The pH of the

mixture was then adjusted to 3 using a 0.5 M H2SO4 (Suprapur® 96 %, Merck) aqueous solution and left for 24 extra hours. Finally the solution was filtered and copiously washed with deionized water before being dried at T = 383 K for 1 hour.

Synthesis of Agglomerated Solid PtNi/C Nanoparticles Agglomerated solid PtNi/C nanoparticles were obtained by following a similar protocol as the synthesis of monodisperse PtNi/C nanoparticles but without the addition of monosodium citrate. First, 53 mg of H2PtCl6.6H2O (99.9 % metals basis, Alfa Aesar ) and 4.3 mg of NiCl2.6H20 metal salts (Putratronic® 99.9995 %, Alfa Aesar) were dissolved in a vial containing a 20 mL mixture of deionized water and ethylene glycol (EG) (Rotipuran® ≥ 99.9 %, Roth) (volume ratio 1:1). 79 mg of carbon black (Vulcan XC72, Cabot) targeting 20 wt. % Pt loading was dispersed by sonication in a separated vial containing also 20 mL of 1:1 EG:water mixture. Then, the contents of each vial was mixed in 20 mL of pure EG, leading to a 2:1 EG:water volume ratio. The pH of the obtained mixture was adjusted to 10 using a 0.5 M NaOH (Suprapur® 99.99 %, Merck) solution (diluted in 1:1 EG:water mixture). The 27 ACS Paragon Plus Environment

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resulting suspension was afterwards kept under vigorous stirring for 1 hour under argon atmosphere before being refluxed at T = 433 K for 3 hours. The solution was allowed to cool down to room temperature under air atmosphere for 12 hours with continuous stirring 85. The pH of the mixture was then adjusted to 3 using a 0.5 M H2SO4 (Suprapur® 96 %, Merck) aqueous solution and left for 24 extra hours. Finally the solution was filtered and copiously washed with deionized water before being dried at T = 383 K for 1 hour.

Synthesis of Ni/C Nanoparticles Ni/C nanoparticles were synthetized according to a protocol adapted from Ref.

43

. Briefly,

236 mg of sodium citrate (≥ 99 %, Roth), 171 mg of carbon Vulcan XC72 (Cabot) and 174 mg of NiCl2.6H20 metal salts (99.9 % metals basis, Alfa Aesar) were mixed to 20 mL ethylene glycol (EG) (Rotipuran® ≥ 99.9 %, Roth). A 5 wt. % solution of KOH (99.99 % trace metals basis, Sigma Aldrich) in EG was then added dropwise to reach a 9-10 pH. The solution was then heated at T = 433 K for 20 hours in air atmosphere. After natural cooling, the final product was collected by filtration and dried at T = 383 K for 45 minutes. The protocol targeted a Ni weight fraction of 20 %, but AAS and TEM analysis revealed that the resulting powder featured only 7 wt. %. Ni.

Synthesis of Ni@Pt/C Nanoparticles Ni@Pt/C nanoparticles were synthetized according to a protocol adapted from Ref.

43

.

Briefly, 1.66 mL of a 20 mg mL-1 H2PtCl6.6H2O (99.9 % metals basis, Alfa Aesar ) aqueous solution was mixed to 15 mL EG (Rotipuran® ≥ 99.9 %, Roth), a 5 wt. % solution of KOH (99.99 % trace metals basis, Sigma Aldrich) in EG was then added dropwise to reach a pH ~9-10. Then, 50 mg of 7 wt. % Ni/C nanoparticles synthetized according to the previously described protocol were added to the mixture. The solution was heated at T = 433 K for 6 28 ACS Paragon Plus Environment

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hours in open atmosphere. After natural cooling, the final product was collected by filtration and dried at T = 383 K for 45 minutes. The nominal and experimentally measured weight fractions in Ni were 5.6 and 1 wt. %, which suggested a loss in Ni content during the deposition of Pt most likely due to the combination of chemical reduction and galvanic displacement.

Synthesis of Hollow PtNi/C Nanoparticles 154 mg of Pt(NH3)4Cl2 (Alfa Aesar, Specpure) and 180 mg of NiCl2 (Fluka, > 98.0 %) were first mixed with 300 mg Vulcan XC72 (Cabot), 10 mL of ethanol and 140 mL of de-ionized water (Millipore). An aqueous solution of NaBH4 (Aldrich 99.99 % - 5.5 mmol, 0.22 M) was then added at a rate of 5 ml min-1 and stirred for 1 h under magnetic stirring at room temperature (T = 293 ± 2 K). The resulting mixture was filtered, thoroughly washed by deionized water and dried for 45 minutes at T = 383 K. The catalysts powder was then acidtreated for t = 22 h in a stirred 1 M H2SO4 (Suprapur® 96 %, Merck) solution at T = 293 K and then filtered, washed and dried again following the same procedure.

Synthesis of Sea Sponge PtNi/C Nanoparticles Sea sponge PtNi/C nanoparticles were synthetized via an original polyol method described here for the first time. 120 mg of Vulcan XC72 (Cabot), 54mg of Pt(NH3)4Cl2 (Premion 99.995% metals basis, Alfa Aesar) and 37 mg NiCl2.6H20 metal salts (99.9 % metals basis, Alfa Aesar) were mixed in 20 mL ethylene glycol. After setting the pH to 9-10 by dropwise addition of a 5 wt. % solution of KOH in EG (Rotipuran® ≥ 99.9 %, Roth), the solution was heated in reflux at T = 473 K for 20 hours. The final product was allowed to cool in open atmosphere and was then collected by filtration and dried at T = 383 K for 45 minutes. The catalysts powder was then acid-treated for t = 22 h in a stirred 1 M H2SO4 (Suprapur® 96 %, 29 ACS Paragon Plus Environment

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Merck) solution at T = 293 K and then filtered, washed and dried again following the same procedure.

Wide Angle X-ray Scattering Measurements The wide-angle X-ray scattering measurements were performed at the ID31 beamline of European Synchrotron Radiation facility (ESRF) in Grenoble, France. The high energy X-ray radiation (61 keV or 0.20 Å) was focused on the powder sample contained in the 1 mm diameter Kapton capillary, and the scattered signal was collected by a Dectris Pilatus CdTe 2M detector positioned 300 mm behind the sample. The size of the beam at the sample position was 4 × 30 µm (vertical × horizontal). The energy, detector distance and tilts were calibrated using a standard CeO2 powder and the 2D diffraction patterns were reduced to the presented 1D curves using the pyFAI software package 86.

Rietveld Refinements Rietveld refinements were carried using the Fullprof software 87 for 2θ between 3.5 and 30°, using the Fm-3m structure of Pt metal. The instrumental resolution function was determined by the refinement of the CeO2 standard sample. After several trials, the Thomson-CoxHastings profile function was adopted with possibility for uniaxial anisotropic broadening from size origin

88

. The background was described by an interpolated set of points with

refinable intensities (this procedure allowed a more accurate description of the background leading to improved diffraction peak profiles). Final refinements included the background points, a scale factor, an overall atomic displacement parameter (a.d.p.), the cubic cell parameter and profile three parameters controlling isotropic strain/size and uniaxial size broadening effects.


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Pair Distribution Function Analysis The PDF G(r) is based on the Fourier transform to direct space of total scattering powder diffraction data. It yields the probability of finding two atoms separated by a distance r in a sample 89 and thus provides a multiscale information about the atomic arrangement. Here, the PDF is defined as 90:

k

a ') ) 4cde') − e g ) Mk lmn idj'i ) − 1g$#&i !i h lop

(9)

where ρ(r) is the microscopic pair density, ρ0 the average pair density, S(Q) the structure function and Q = (4π.sinθ)/λ. Due to the subtraction of the average pair density, the G(r) function oscillates around 0, reflecting the excess presence or absence of interatomic distances at the corresponding r value. At distances above the particle size or coherent structural domain size, the microscopic pair density becomes equal to the average pair density and the PDF vanishes. PDFs were obtained from the powder patterns using the PDFGetX3 software 91 up to Qmax = 18 Å-1. The background scattering was subtracted using the pattern from the carbon support, while the instrumental resolution parameters were determined from the refinement of the CeO2 sample. The PDFs were refined using the PDFGUI software 92. Final refinement cycles included the scale factor, an overall a.d.p., the cubic cell parameter and the nanoparticle diameter, or isotropic coherent structural domain size 93.

TEM Measurements and STEM X-EDS Mapping The electrocatalysts were examined with a JEOL 2010 TEM operated at 200 kV with a point to point resolution of 0.19 nm. The X-EDS elemental maps were acquired using a JEOL 2100F microscope operated at 200 kV and equipped with a retractable large angle Silicon 31 ACS Paragon Plus Environment

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Drift Detector (SDD) Centurio detector. The X-EDS spectra were recorded on individual nanoparticles by scanning the beam in a square region adjusted to the particle size. The quantitative analyses were performed on Pt L and Ni K lines using the K-factor provide by the JEOL software.

Electrochemical Measurements All the glassware accessories used in this study were first cleaned by soaking in a H2SO4:H2O2 mixture for at least 12 hours and thoroughly washing with ultrapure water. The 1 M H2SO4 solution used for acid-leaching was prepared with Milli-Q water (Millipore, 18.2 MΩ cm, total organic compounds < 3 ppb) and H2SO4 96 wt. % (Suprapur, Merck). The electrochemical measurements were conducted using an Autolab PGSTAT302N in a custom-made four-electrode electrochemical cell thermostated at T = 298 ± 1 K. Fresh electrolyte solution (0.1 M HClO4) was daily prepared with Milli-Q water (Millipore, 18.2 MΩ cm, total organic compounds < 3 ppb) and HClO4 96 wt. % (Suprapur, Merck). The counter-electrode was a glassy carbon plate and the reference electrode - a commercial reversible hydrogen electrode (Hydroflex, Gaskatel GmbH) connected to the cell via a Luggin capillary. A Pt wire connected to the reference electrode was used to filter the high frequency electrical noise. To prepare the working electrodes, a suspension containing 10 mg Pt/C from the fresh or the heat-treated hollow PtNi/C nanoparticles, 54 µL of 5 wt. % Nafion® solution (Electrochem. Inc.), 1446 µL of isopropanol and 3.6 mL (18.2 MΩ cm) of deionized water (MQ-grade, Millipore) was made. After sonication for 15 minutes, 10 µL of the suspension was pipetted onto a glassy carbon disk, and sintered for 5 minutes at T = 383 K to ensure evaporation of the Nafion® solvents yielding a loading of ca 20 µgPt cm-2geo. Prior to any electrochemical experiment, the working electrode was immersed into the deaerated 32 ACS Paragon Plus Environment

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electrolyte at E = 0.40 V vs. RHE (Ar >99.999 %, Messer). Cyclic voltammograms were recorded in Ar-saturated electrolyte between 0.05 V and 1.23 V vs. RHE with a potential sweep rate of 0.050 or 0.020 V s-1. The real surface area was estimated using COad stripping coulometry assuming that the electrooxidation of a COads monolayer requires 420 µC per cm2 of Pt. The CO saturation coverage was established by bubbling CO for 6 min and purging with Ar for 34 min, while keeping the electrode potential at E = 0.1 V vs. RHE. The electrocatalytic activity for the ORR was measured in O2-saturated 0.1 M HClO4 solution (20 minutes of purging by oxygen > 99.99 %, Messer) by linearly sweeping the potential from 0.20 to 1.05 V vs. RHE at a potential sweep rate of 5 mV s-1 and at different rotational speeds (400, 900, 1600, and 2500 rpm). The ORR specific/mass activity was determined by normalizing the current measured at E = 0.95 V vs. RHE, after correction from the oxygen diffusion in the solution and the Ohmic drop, to the real surface area/mass of deposited Pt determined by COad stripping voltammetry, respectively.

AUTHOR INFORMATION Corresponding Authors * E-mail for R.C.: [email protected]. * E-mail for F.M.: [email protected] Author Contributions All authors performed experimental work, discussed the results, drew conclusions and have given approval to the final version of the manuscript.

CONFLICT OF INTEREST The authors declare no competing financial interest. 33 ACS Paragon Plus Environment

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ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge via the Internet at http://pubs.acs.org. Results of the PDF refinement between r =1 and 20 Å and relation between the chemical composition and the diameter of the Ni-core@Pt-shell/C nanoparticles; scanning transmission electron microscopy, scanning electron microscopy images; X-EDS elemental maps of the “sea sponge” PtNi/C catalyst; relation between the Ni composition estimated by X-EDS and the lattice parameter calculated from Rietveld refinement of the WAXS patterns as well as results of the Rietveld refinements and PDF analysis; cyclic voltammograms for the different nanostructures; ORR activity of the various catalysts at E = 0.90 V vs. RHE and specific activity vs. raw microstrain plot for the various nanostructures.

ACKNOWLEDGEMENTS This work was performed within the framework of the Centre of Excellence of Multifunctional Architectured Materials ‘CEMAM’ n° ANR-10-LABX-44-01. The authors acknowledge financial support from University Grenoble-Alpes through the AGIR program (grant number LL1492017G), and from the French National Research Agency through the HOLLOW project (grant number ANR-14-CE05-0003-01).

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Targeting structural disorder is a promising approach to improve the ORR activity of PtNi/C electrocatalysts


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Beyond Strain and Ligand Effects: Microstrain-Induced Enhancement

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