This journal is ©The Royal Society of Chemistry 2019 Energy Environ. Sci.
Cite this: DOI: 10.1039/c9ee01185d
Real-time imaging of activation and
degradation of carbon supported octahedral
Pt–Nialloyfuelcellcatalystsatthe
nanoscale using in situ electrochemical
liquid cell STEM†
Vera Beermann,‡
a
Megan E. Holtz,‡
b
Elliot Padgett,
b
Jorge Ferreira de Araujo,
a
David A. Muller*
bc
and Peter Strasser *
a
Octahedrally shaped Pt–Ni alloy nanoparticles on carbon supports have demonstrated unprecedented
electrocatalytic activity for the oxygen reduction reaction (ORR), sparking interest as catalysts for low-
temperature fuel cell cathodes. However, deterioration of the octahedral shape that gives the catalyst its
superior activity currently prohibits the use of shaped catalysts in fuel cell devices, while the structural
dynamics of the overall catalyst degradation are largely unknown. We investigate the time-resolved
degradation pathways of such a Pt–Ni alloy catalyst supported on carbon during cycling and startup/
shutdown conditions using an in situ STEM electrochemical liquid cell, which allows us to track changes
happening over seconds. Thereby we can precisely correlate the applied electrochemical potential
with the microstructural response of the catalyst. We observe changes of the nanocatalysts’ structure,
monitor particle motion and coalescence at potentials that corrode carbon, and investigate the dissolu-
tion and redeposition processes of the nanocatalyst under working conditions. Carbon support motion,
particle motion, and particle coalescence were observed as the main microstructural responses to
potential cycling and holds in regimes where carbon corrosion happens. Catalyst motion happened
more severely during high potential holds and sudden potential changes than during cyclic potential
sweeps, despite carbon corrosion happening during both, as suggested by ex situ DEMS results. During
an extremely high potential excursion, the shaped nanoparticles became mobile on the carbon support
and agglomerated facet-to-facet within 10 seconds. These experiments suggest that startup/shutdown
potential treatments may cause catalyst coarsening on a much shorter time scale than full collapse of
the carbon support. Additionally, the varying degrees of attachment of particles on the carbon support
indicates that there is a distribution of interaction strengths, which in the future should be optimized for
shaped particles. We further track the dissolution of Ni nanoparticles and determine the dissolution rate
as a function of time for an individual nanoparticle – which occurs over the course of a few potential
cycles for each particle. This study provides new visual understanding of the fundamental structural
dynamics of nanocatalysts during fuel cell operation and highlights the need for better catalyst-support
anchoring and morphology for allowing these highly active shaped catalysts to become useful in PEM
fuel cell applications.
a
Electrochemical Energy, Catalysis and Material Science Laboratory, Department of Chemistry, Technical University Berlin, 10623 Berlin, Germany.
E-mail: pstrasse[email protected]
b
School of Applied and Engineering Physics, Cornell University, Ithaca, NY 14850, USA. E-mail: david.a.mul[email protected]
c
Kavli Institute at Cornell for Nanoscale Science, Cornell University, Ithaca, NY 14850, USA
†Electronic supplementary information (ESI) available: Experimental procedures, HAADF STEM images for beam damage comparison, ex situ STEM images, in situ
and ex situ electrochemical data, differential electrochemical mass spectral data, and HAADF STEM Movies S1–S3. See DOI: 10.1039/c9ee01185d
‡These authors contributed equally to this work.
Received 13th April 2019,
Accepted 20th May 2019
DOI: 10.1039/c9ee01185d
rsc.li/ees
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Broader context
In situ and operando electrochemical liquid cell scanning transmission electron microscopy (E-chem STEM) offers previously unavailable insight into the
evolution of the morphology and structure, as well as the composition and electronic structure, of nanoscale electrocatalysts under operating environments.
While conventional ex situ STEM or identical location STEM present before/after images, and environmental TEM is largely limited to low-pressure
environments, E-chem STEM is able to perform operando imaging of the physico-chemical transformations at electrified solid–liquid interfaces in real-
time. In this contribution, E-chem STEM is used to track the microstructural events during the degradation of carbon-supported octahedrally-shaped PtNi fuel
cell nanocatalysts during potential cycling and potential holds in real time, revealing carbon support motion, particle motion, and rapid particle coalescence
within seconds. This suggests that fuel cell startup/shutdown conditions can cause catalyst coarsening on a much shorter time scale than previously assumed
and call for better catalyst-support anchoring. Tracking the activation and degradation of PtNi alloy nanocatalysts provides a clear and visual understanding of
their fundamental structural dynamics during fuel cell operation.
Introduction
Heightened interest in alternative renewable power sources
has increased technological and scientific focus on fuel cell
technologies. A large part of such research and development is
focused on novel cathode catalyst materials for the oxygen
reduction reaction (ORR) where efficiency losses have remained
high. New catalyst systems based on alloying Pt with transition
metals like Fe, Co, Ni, and Cu in unshaped alloy nanoparticles
have led to improved ORR activities.
1,2
At least an order of
magnitude improvement in catalytic ORR activity over these
conventional alloy nanoparticles has been reported for shape-
controlled octahedral Pt–Ni alloy particles, because they exclu-
sively expose highly active {111} Pt–Ni facets.
3–6
While unshaped
Pt–M alloy fuel cell catalysts are beginning to be deployed in
commercial applications,
7,8
shape-controlled particles still face
challenges in terms of stability, especially in the final MEA
(Membrane electrode assembly) device.
7,9
Shaped Pt–Ni nano-
particles have been observed to quickly lose their shape after
cycling, in part due to nickel dissolution.
10,11
The detailed
degradation processes of octahedral Pt–Ni particles on carbon
supports have remained elusive. Hence, better understanding
of their structural behavior and degradation is critically
required before these high-activity catalysts can be deployed
in commercial applications.
Many physical characterization methods have been used to
gain a better understanding of the morphology and composi-
tion of fuel cell catalyst materials before and after degradation.
Most of the work to date has relied on ex situ characterization
techniques, often involving scattering from X-ray, light or
electrons to describe the initial or post mortem material. For
fuel cell catalyst nanoparticles, transmission electron micro-
scopy (TEM) is a popular method to determine the particle
shape and distribution on the support material, as well as
elemental distribution and stability. Identical location TEM (IL
TEM) has been used extensively for some Pt-based nanoparticle
materials to track and study changes of identical particles or
catalyst parts before and after electrochemical treatment.
12–16
In addition to ex situ techniques, there has been a recent surge
of interest and capability in in situ and operando methods that
enable probing the material under working conditions, garnering
valuable understanding of material operation and degradation.
17–29
Lately, several groups reported in situ electrochemical TEM
investigations on fuel cell materials
30,31
and lithium ion battery
materials.
18,32–39
These experiments typically use liquid-cell
systems in conventional TEMs with SiN windows on chips
encapsulating a thin liquid layer. Using this powerful tool,
it is possible to perform conventional electrochemistry and
electrocatalysis while imaging the reactive particles of interest
in real time on the nanometer scale, obtaining operando
information about the nanocatalyst at work.
In this study, we investigate the degradation of carbon-
supported octahedral shaped Pt–Ni alloy nanoparticle catalysts
for advanced fuel cell cathodes. We use an in situ electro-
chemical liquid-cell and scanning transmission electron micro-
scopy (STEM) to track the nanoscale changes to the catalyst
under electrochemical conditions that arise or are applied at the
cathode. We monitor the translational, structural, and – thanks to
atomic number contrast in high angle annular dark field (HAADF)
STEM – compositional dynamics and evolution of individual
nanoparticles as well as ensembles of nanoparticles in real time
and with nanometer-scale resolution and support findings regard-
ing carbon corrosion with DEMS. This study gives new insight
into how initially shape-controlled nano-octahedra transform into
unshaped and partially agglomerated particle clusters, and pro-
vides visualization of how the degradation of the carbon support
affects the catalyst material.
Results and discussion
We investigated B8 nm octahedral Pt–Ni nanoparticles that
were supported on Vulcan carbon supports (Fig. 1a). These
supported Pt–Ni/C particles showed electrochemical ORR activ-
ity that was about 25greater than commercial Pt/C, and, from
a more practical perspective, were large enough to be imaged in
an in situ TEM liquid environment at low beam doses. They
exhibit an average composition of Pt
62
Ni
38
and a Pt enriched
surface, and were carefully washed with ethanol in order to
remove remaining surfactants from the synthesis.
40
In addition
to octahedral Pt–Ni nanoparticles, the catalyst also contained a
minor alloy phase consisting of 20–50 nm Ni-rich particles,
which enabled study of rapid electrochemical dissolution pro-
cesses of undesired alloy phases during catalyst activation.
These Ni-rich particles consist of a Pt-rich core which is
encased by a thick Ni shell resulting in an overall composition
of Pt
5
Ni
95
.
40
Fig. 1a shows the catalyst particles ex situ as
synthesized: the initial octahedral shape is evident from the
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faceting of the particles, which are homogeneously distributed
on the carbon support. After ex situ voltammetric stability tests
in perchloric acid, the octahedral particles lost their sharp
faceting and showed agglomeration (shown in Fig. 1b). The
particles display slightly concave edges and facets suggesting
that the facets are relatively Ni-rich and are therefore etched
more quickly during acid and electrochemical treatments,
leaving behind a Pt-rich skin. After treatment, we observe an
average composition of Pt
69
Ni
31
, reflecting the relative loss of Ni.
40
To better understand this degradation, we perform an in situ
study of the real time nanometer-scale evolution of the carbon
supported, octahedral Pt–Ni fuel cell catalysts.
The in situ TEM experiments were carried out in a Proto-
chips Poseidon holder and a flow cell chip equipped with a silicon
nitride window. The cross section and top view of the chip
are illustrated in Fig. 1c showing the carbon working electrode
and the Pt reference and counter electrodes.
18,29
The platinum
reference electrode was calibrated in 0.1 M perchloric acid using
the well-known characteristics of the hydrogen underpotential
deposition region of platinum-based materials, as shown in
Fig. 1d. With that, 0.0 V
RHE
was correlated to 0.8 V
Pt
.Allfurther
potentialsarereportedagainstthereversiblehydrogenelectrode
(RHE) based on this calibration to allow better comparability
to the literature. The cell had a liquid thickness of 300 nm,
as estimated by electron energy loss spectroscopy.
41
Prior to the in situ TEM electrochemical investigations,
we identified a suitable beam dose that did not visibly affect
the octahedral particles in the electrolyte for the duration of
the applied electrode potential. Even though the beam alone
may not influence the particles, the combination of beam and
electrochemical cycling may have an effect. To account for this,
we compared the final state (after electrode potential cycling) of
the particles that were imaged during the in situ experiment to
other particles that were not continuously imaged in the
electrochemical cell, to crosscheck for similar transformations.
We further compared particles that were on the electrode to
those that were not on the electrode to ensure that the effects
were electrochemical rather than chemical inside the liquid cell
(Fig. S1, ESI†). As a final check, we qualitatively compared the
data from in situ experiments to those of ex situ experiments.
Overall, we found that the electrochemical effects observed
were not driven by the electron beam, nor from the chemical
environment in the cell. However, the in situ experiments
appeared to be harsher on the particles due to the additional
effects of the electron beam.
We performed in situ electrochemical STEM investigations
using different electrochemical electrode potential cycling
protocols resembling those routinely applied to single fuel cells
to electrochemically activate and stress-test their catalysts,
42,43
cycle in standard operating ranges, and cycle under extreme
potentials to simulate startup/shutdown conditions – which
may occur in a PEMFC, reaching values of up to 1.6 V
RHE
, due to
oxygen and hydrogen present at the anode side.
First, the catalyst on the working electrode was cycled 20 times
inside a potential range between 0.0 and +1.0 V
RHE
with 100 mV s
1
to mimic an activation procedure. Fig. 2a shows the applied
potential profile with selected marked time points corres-
ponding to the STEM images shown in Fig. 2b–g. The selected
field of view displays a collection of octahedral nanoparticles
surrounding a larger Ni rich particle. During the electrochemical
treatment, there were no discernible changes in the octahedral
particle structure. While surface Ni may dissolve, the relatively
high Pt content in the alloy passivates the surface preventing
further dissolution. However, the large Ni-rich particle marked
by the arrow in Fig. 2b–g gradually dissolved according to
Ni
0
-Ni
2+
+2e
(E
0
= 0.26 V
44
) during the applied potential
cycling. After 10 cycles (Fig. 2c) the particle started to lose mass
as observed by first a change in the overall HAADF intensity and
eventually by a change in the particle diameter. The dissolution
process took place over several cycles, and after 15 cycles (Fig. 2f)
only a small fraction remained. Because the HAADF intensity
began to drop before the diameter of the particle shrank, and
because the intensity of the particle became modulated (Fig. 2c)
compared to its initial state (Fig. 2b), we expect the particle first
became less dense, possibly becoming sponge-like and porous,
before it disappeared completely.
60
In our corresponding ex situ
experiment discussed below, we also observed a Ni-rich particle
that had modulations in the ADF intensity (Fig. S2e, ESI†),
further suggesting that the particles do not dissolve radially
inward. The dissolved Ni-rich particle left behind an octahedral
particle, which may have been contained inside the Ni-rich
phase, evidenced by the bright contrast in the center (see
Fig. 2b). The dissolution of another Ni-rich particle was observed
Fig. 1 Preliminary characterization and in situ TEM chip design. (a) Initial
particles ex situ after synthesis, showing octahedral shape with strong
faceting in the {111} planes (see inset) and (b) ex situ after electrochemical
potential cycling for 40 cyclic voltammograms, where facets are curved
and particles are agglomerated (see inset). (c) Overview of electrochemical
cell setup, with a cross section of the liquid cell holder on the top and view
of the electrodes on the bottom. (d) Cyclic voltammogram of Pt–Ni
nanoparticles on the carbon working electrode inside the electrochemical
cell in 0.1 M HClO
4
with a sweep rate of 100 mV s
1
.
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between Fig. 2b (where it is a fractional particle already) and
Fig. 2c in the central lower part of the frame. After the first
20 cycles, there were still Ni-rich particles remaining in areas
outside of the region imaged in Fig. 2, and after an additional
20 cycles (total of 40 cycles), the remaining Ni-rich particles
elsewhere on the electrode also disappeared. The real-time
imaging of the catalyst during electrochemical cycling with the
animated potential profile corresponding to the data in Fig. 2b–g
is shown in Movie S1 of the ESI.†
The opportunity to image the Ni dissolution process in real
time allows us to estimate a dissolution rate for an individual
Ni particle. Annular dark field (ADF) intensities of the Ni
particle were obtained by integrating over a region containing
the Ni particle, subtracting off the background intensity from a
neighboring region to account for liquid thickness variation,
and subtracting off the average intensity of the last 5 frames,
when the Ni particle had fully disappeared. In Fig. 2h, we first
see a gradual decrease in ADF intensity, which corresponds in
the image to the particle becoming less dense. Then, the
particle dissolves rapidly, decreasing in size drastically between
cycle 9 and 13. While we expect Ni dissolution above 0.26 V
RHE
,
we do not observe periodic changes in the dissolution rate with
the potential in each cycle. Although we expect Ni dissolution to
be a thermodynamically driven process, the sudden dissolution
may happen due to an increase of surface area as it becomes
spongy, with exposure of a fresh, unpassivated surface. From
the background subtracted ADF intensity of the Ni-rich particle,
we can calculate its dissolution rate because the ADF intensity
is proportional to the mass present. We assume that the initial
ADF intensity corresponds to the number of nickel atoms in a
solid, spherical shell of nickel that has the inner and outer
diameters as measured in the ADF image. The derivative of the
ADF intensity which is scaled to number of Ni atoms present
then gives the dissolution rate in atoms per s. This dissolu-
tion rate, plotted in Fig. 2i, reaches values of around 10 000 to
30000 atoms per second during cycles 9 through 13. Assuming
a perfectly round Ni particle with a diameter of 10 nm, the
number of initial particle surface atoms would be around 4500 Ni
atoms. Thus, a dissolution rate of 1000 atoms per s corresponds
to one monolayer every 4.5 seconds (for this 10 nm diameter
particle). Assuming it is etching along the 111 direction of
Ni which has a 2 Å spacing, that corresponds to an etching rate
of about 27 nm min
1
– whereas typical etch rates for bulk
Ni-materials are on the order of 10 nm min
1
.
45
So, for a
nanoparticle system where the geometry, chemistry and local
potential can be quite different, we believe this is a reasonable
etch rate for Ni.
After the peak in dissolution, the remaining amount of
Ni-rich particles slowly dissolve away – perhaps due to low
surface area or Pt enrichment as Ni is selectively removed. This
is the first quantitative observation of the nanometer-scale
reaction dynamics of a selective Ni dissolution process during
the disappearance of a Ni-rich alloy particle.
To compare the activation processes during the in situ experi-
ments and corresponding ex situ treatments, ex situ experiments
in a conventional three-electrode cell setup were carried out using
Fig. 2 HAADF STEM in situ imaging of the catalyst structure during electrochemical potential cycling between 0.0 and +1.0 V
RHE
in 0.1 M HClO
4
for
20 CV with 100 mV s
1
sweep rate. (a) Potential profile over time with marked points corresponding to the images in (b–g). A Ni-rich particle marked by
the arrow disappears during cycling, first becoming less dense, then spongy, and finally dissolving completely. (h) ADF intensities of the Ni-rich particles
over time during potential cycling and (i) the resulting Ni dissolution rate.
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identical protocols (Fig. S2, ESI†). The observed cyclic voltam-
metry current in the in situ cell may differ from an ex situ
experiment for several reasons, including the small area of the
working electrode, the electrolyte that has not been degassed,
and the different diffusion geometry in the thin encapsulated
cell. Nevertheless, some trends in the in situ cyclic voltammo-
grams are noteworthy as they show the same processes as
observed in the STEM images (Fig. S2b, ESI†). With increasing
cycle number, the current at higher potentials due to Ni dissolu-
tion trails off, and the redox waves inside the H
upd
region become
sharper, which is consistent with generating a cleaner, more
Pt-rich, and more facetted surface due to Ni dissolution and Pt
diffusion.
11,40,46,47
Unlike in situ, we noticed several residual large
Ni-rich nanoparticles after 40 ex situ cycles (Fig. S2d and e, ESI†),
some of which appeared to have experienced partial dissolution.
Thus, we conclude that the in situ conditions were more corrosive
than the ex situ conditions, possibly due to the confined liquid
cell environment, electron beam effects, and the lower geometric
Pt loading.
Our observations show that the typical electrochemical
activation comprising cyclic voltammetry in liquid does not
harm the shape or distribution of the Pt–Ni octahedra, validating
the suitability of these commonly used activation procedures.
Dissolution of nickel in the Pt–Ni octahedra was not possible to
determine at the dose-limited resolution because the ADF inten-
sity is largely dominated by the large Pt signal and insensitive to
small changes due to Ni dissolution. While the Pt–Ni octahedra
are stable, the undesired Ni-rich clusters dissolve within minutes
of the activation protocol. In all, we successfully imaged the
activation dynamics of a shaped Pt alloy fuel cell catalyst by
electrochemical dealloying and selective corrosion in real time.
Next, we studied the impact of sequential sets of potential
cycles separated by periods with constant applied electrode
potential, a frequently used test cycle motif for automotive
or stationary PEM fuel cells. The potential versus time profile
is given in Fig. 3a, again with marked time points for the
snapshots shown in Fig. 3b–h, with two zoomed in smaller
regions to track individual particles. An initial potential hold at
Fig. 3 HAADF STEM in situ imaging of the impact on the catalyst structure of an electrochemical sequence consisting of electrochemical potential
cycling between 0.0 to +1.2 V
RHE
and 0.0 to +1.4 V
RHE
for 10 CV with 100 mV s
1
and holding on different upper potentials in 0.1 M HClO
4
. (a) Potential
profile over time with marked points for shown images. (b–i) Images taken at the marked potential and cycle number, with cutouts of small regions from
other areas of the movie sequence to track individual particles (left) and stringy Pt growth (right). The field of view of the larger cutout (left) is 50 nm and
44 nm for the smaller (right). For example, the green arrow marks two Pt–Ni particles that slowly move together, while both growing larger. On the right
cutout, we see stringy particles growing mostly during potential cycling, and moving during holding.
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+0.8 V
RHE
is followed by cyclic voltammetry with an upper
potential of +1.2 V
RHE
with 100 mV s
1
sweep rate. Then, there
is a subsequent fixed potential hold at +1.2 V
RHE
, followed
by cyclic voltammetry to an upper potential of +1.4 V
RHE
at
100 mV s
1
. This is finished with another fixed potential hold
at +1.4 V
RHE
. For this experiment, we imaged a region of the
Pt–Ni nanoparticles that was previously not immersed in the
electrolyte, and did not show evidence of previous cycling. This
was possible because the liquid cell was only partly full or had a
small bubble over the working electrode, and as the liquid
flowed over the course of the experiment, the liquid front
gradually moved from one side to the other to cover the entire
electrode region.
Still images in Fig. 3b–i are taken from the in situ STEM data
that are presented in Movie S2 (ESI†), which also incorporates
the animated potential profile. During the hold at +0.8 V
RHE
carbon remained stable and we saw minimal changes in the
catalyst structure (Fig. 3c). When cycling to +1.2 V
RHE
, a slight
movement of the carbon support was observed, visible by the
motion of whole particle agglomerates. We also observed
the nucleation and growth of stringy, new Pt-rich deposits, as
seen in the smallest cutouts in Fig. 3b–i, which is likely
due to chemical metal redeposition as we will later discuss
(Fig. 3d, f and g, yellow arrow).
During the hold at +1.2 V
RHE
the redeposited Pt/Ni abruptly
moves, as if it was not firmly attached and became dislodged
when held at elevated potential (Fig. 3e, orange circle). When
the Pt-rich redeposits appear to collide with other parts of the
sample or working electrode, their motion slows or stops. At the
same time, the former octahedral Pt–Ni nanoparticles also
seem mobile (see the smaller cutouts going from Fig. 3d to e)
and started to grow slowly in size (Fig. 3f, orange cycle), as
expected from both electrochemical and beam-induced rede-
position. Upon cycling to +1.4 V
RHE
, the redeposited Pt/Ni again
becomes mobile and swings about, while carbon-supported
Pt–Ni particles also move notably (Fig. 3f and g). Additional
stringy Pt-rich deposits form. Finally, holding the potential at
+1.4 V
RHE
again causes abrupt motion of the redeposited Pt
(Fig. 3h, orange cycle), while carbon corrosion (C + 4H
2
O-
CO
2
+4H
+
+4e
(E
0
= 0.21 V)
44
) appears to occur rapidly enough
to cause sustained motions of the carbon-supported Pt–Ni
particles (Fig. 3i) and the Pt–Ni particle growth continues.
Several phenomena were observed in Movie S2 (ESI†)–
including growth and motion of stringy Pt-rich deposits, Pt–Ni
nanoparticle growth, and catalyst structure changes, which are
likely due to carbon corrosion – which we will discuss in the
following paragraphs.
First, we will discuss the stringy metal redeposition which
is highlighted in the smallest (right) cutout of Fig. 3b–i. We
believe that the stringy redeposition is primarily Pt being
reduced and redeposited, because the high contrast in the
STEM images is consistent with a Pt or Pt-rich composite.
Additionally, Ni has a lower standard potential than Pt
(0.26 V
RHE
for Ni vs. 1.18 V
RHE
for Pt), indicating that Pt will
be preferentially reduced while Ni will more likely be dissolved.
From the literature we also believe that Ni is unlikely to become
reduced again on the particle surface.
47
There are two potential
mechanisms for the observed stringy redeposition: (1) chemical
redeposition, which is driven by a reducing chemical environment
such as one generated by the electron beam,
48
or (2) electro-
chemical reduction, which occurs at low, reducing potentials only
at locations which have electrical contact with the working
electrode. We observe that these Pt-rich deposits form faster
during cyclic voltammetry than during potential holds. There
may be two reasonable explanations for this: (1) electrochemi-
cally assisted deposition might occur during the sweep to
low potentials, or (2) the Pt is chemically deposited, and this
happens faster during the cyclic voltammetry because of
increased Pt dissolution during the cyclic voltammetry, so there
is more Pt in solution to be chemically deposited (Pt
0
2Pt
2+
+
2e
(E
0
= 1.18 V)).
44
If the effect were purely electrochemical
deposition at low potentials, we would expect the time at low
potential to determine the amount of redeposition. However,
if the redeposition is chemical, it will be faster at sweeps to
higher potential since more Pt will be dissolved into the system,
which will then be available for redeposition. Indeed, we
observe that the deposition appears to be faster at the cyclic
voltammetry with upper potentials of 1.4 V
RHE
compared to
1.2 V
RHE
, indicating that the redeposition is likely in part a
chemical process. We thus conclude that the stringy deposits
are chemically redeposited Pt that is driven by the reducing
effects of the electron beam. These are similar to the formation
of so-called pure Pt deposits (referred to as ‘‘Pt bands’’) reduced
by dissolved hydrogen inside fuel cell membranes.
49
While most
of the chemically redeposited Pt grows in the polymer matrix of
the membrane, some hydrogen may make it to the cathode and
deposit stringy Pt in the open pore spaces of the cathode. We see
that the chemically deposited stringy Pt may be quite mobile in the
pores in the cathode material during cycling.
We observed the most severe and sudden changes to the
catalyst structure precisely at the transitions from the voltam-
metric cycling to the potentiostatic holds with chronoampero-
metric monitoring of the current density. The lightly attached,
stringy Pt-rich deposits become loose, very mobile and detached
from the catalyst support when carbon corrosion at high poten-
tials starts to occur. They appear to move until they collide
with another feature in the catalyst structure. Furthermore, the
particle motion and coalescence are correlated with the applied
electrode potential during both cyclings and holds, corroborating
the detrimental effect of the anodic upper potentials. Thus, our
in situ STEM studies evidenced how strongly platinum and nickel
oxidation and dissolution accelerated with increasing upper
turning potentials as predicted by the mean-field Butler–Volmer
relation.
50–52
In addition to the stringy Pt/Ni formation, we also observe
that the initial Pt–Ni octahedral particles grow over the course
of the treatment. During the potential holds at high electrode
potentials, catalyst particles continued to grow and lose their
shape. This may be due to redeposition by the electron beam.
Carbon corrosion is expected to become significant at
potentials of +1.1 V
RHE
and higher (C + 4H
2
O-CO
2
+4H
+
+
4e
(E
0
=0.21V)
44
).
53–55
At these high potentials (around 1.2 V
RHE
),
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weobservetwotypesofPtNinanoparticlemotion(MovieS2,ESI†).
One is that we see carbon support motion and crumpling, which
may be an effect of carbon corrosion, where the catalyst nano-
particles in one region of the carbon support appear to move
together as the carbon bends. The second effect is that the
corrosion appears to weaken the attachment of particles on the
carbon support, causing additional particle migration, coales-
cence and agglomeration. Both of these effects were more
severe at higher potentials. We found that holding at higher
potentials as opposed to potential cycling intensified and
accelerated particle catalyst degradation. High potentials facili-
tate distinct oxidation of all catalyst components in contrast to
cycling, where conditions are temporarily less corrosive at lower
potentials. With the help of DEMS, we note that carbon corro-
sion should happen both during the cyclic sweeps to high
potential, and during the potential holds (see Fig. S5 with
corresponding text, ESI†). Even though corrosion is happening
throughout, the high potential hold is more detrimental to the
overall catalyst morphology.
Upon comparison of ex situ and in situ conditions, the ex situ
conditions again were less harmful to the catalyst structure
than the in situ ones (see Fig. S3, ESI†). After ex situ cycling up
to +1.2 V
RHE
the octahedral particle shape was still clearly
discernible, while edges and the corners were degraded after
cycling to +1.4 V
RHE
. While the general trends were consistent,
the impact of the applied electrode potential on the shape,
particle distribution and carbon corrosion was evidently less
pronounced under the ex situ conditions, which is reasonable
due to the absence of electron beam driven processes.
Finally, we imaged the structural evolution of the Pt–Ni fuel
cell catalysts under conditions simulating startup/shutdown and
air or fuel starvation, which often cause uncontrolled potential
steps in cathodic or anodic directions.
56–58
To achieve that, we first
applied 10 potential cycles to an elevated upper potential of
+1.2 V
RHE
after which the electrode potential experienced a potential
step to above +1.4 V
RHE
. The Pt–Ni nanoparticles in this experiment
had only undergone the activation profile corresponding to Fig. 2.
During the 10 potential cycles to +1.2 V
RHE
the octahedral
shape of most particles appears largely unaffected, but a few
experienced coalescence with close neighboring particles (Fig. 4b–d,
pink arrow) or small motions on the support (Fig. 4b–d, yellow
arrow). Considering that the upper potential lies outside the window
where carbon is kinetically stable, the motion and coalescence may
be due to carbon support corrosion.
The final anodic potential step dramatically affected the
global catalyst structure (Fig. 4e–g) and would have catastrophic
Fig. 4 HAADF STEM in situ imaging of the catalyst structure during electrochemical potential cycling between 0.0 and +1.2 V
RHE
for 20 CV with 100 mV s
1
sweep rate, followed by a step into a high potential. (a) Potential profile over time with marked points corresponding to the images in (b–g). Some changes to
catalyst particles are noted during cycling in (b–d) – for example, coalescence as indicated by the pink arrow and particle motion as indicated by the yellow
arrow (e–g). After cycling, going to a high potential, we see dramatic coalescence, where the Pt–Ni nanoparticles agglomerate into wires. Insets in (a–c) show
enlarged fractions of particles aligning on their facets.
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consequences for the fuel cell performance. The particles became
highly mobile on the carbon support,asshownintheimagesin
Fig.4andMovieS3(ESI†), with the majority of particles colliding
on the timescale of seconds into an agglomerate with a long,
branching geometry. Often, the particles appear to line up in
preferred orientations – often with their edges flush – forming in
straight lines or with regular angles. This suggests that they are
aligning along their {111} crystal facets before fusing together
(see insets Fig. 4a–c, and cartoon schematic in Fig. S3h, ESI†). We
observe similar agglomeration in neighboring regions on the
electrode outside the field of view, as shown in Fig. S3 (ESI†),
showingthatthisoccurredindependentofbeameffects.Aftera
catastrophic, abrupt agglomeration, the catalyst particles continued
to move on the carbon surface as carbon corrosion continues.
The corresponding ex situ images acquired after the same
electrochemical treatment confirm the observed trends (see
Fig. S4, ESI†). Although individual particles still appeared to
be octahedrally shaped they agglomerated at their crystal
facets. From the in situ magnification in Fig. 4, it is difficult
to unambiguously identify the remaining particles as octa-
hedral or unshaped, due to the low beam dose which is required
to avoid radiation damage. All data of Fig. 3 can be inspected in
Movie S3 (ESI†).
We should additionally note that the Pt–Ni particles were
dispersed onto the carbon support as entire particle ensembles,
that is, after their synthesis, rather than grown directly onto it,
which likely resulted in weaker interactions between particles
and the support compared to Pt/carbon catalysts synthesized
with impregnation methods, where a Pt molecular precursor is
reduced on the carbon in a dispersed state. Previous in situ
TEM investigations of unshaped Pt–Co nanoparticles grown
onto the support by impregnation methods found less particle
migration despite extreme carbon corrosion at high potentials.
59
Thisindicatesthatthisfailuremodemaybeuniquetothese
highly active shaped nanoparticles, as it was not observed under
similar conditions in past in situ studies. Because these high
potential conditions may happen unintentionally in practical fuel
cell devices, the ideal catalyst should also be robust under these
conditions. Recent ex situ studies have shown that the choice of
carbon support can dramatically reduce the rate of particle
agglomeration in membrane electrode assembly devices,
61,62
either by providing a stronger anchoring or by physically con-
straining particles in carbon pores to prevent collisions. This
experiment provides a dramatic illustration of catalyst agglo-
meration, which points out that an important avenue of research
will be in reducing the high mobility of shaped particles on the
carbon support to improve their overall stability. Options for
addressing this issue may be improving the chemical anchoring
points that bond particles to the carbon support, or selecting
carbon geometries which either improve this contact area or
constrain the particle motion. Further investigations will be
required to investigate how the presence of ionomer may alter
the effects of carbon corrosion and particle adhesion – as this
experiment was done in a liquid without ionomer or membrane.
Our findings provide a real-time visual demonstration of the
catastrophic effects of uncontrolled fuel cell cathode potential
excursions to values of startup/shutdown. Our results further
underline the critical importance of a strict continuous upper
electrode potential control. Two distinct mechanisms could
contribute to the rapid coalescence at high potential: first,
the instantaneous formation of Pt and Ni surface oxides
induced by the abrupt anodic potential step may have lowered
the particle attachment and caused enhanced mobility; more
likely, however, supported by DEMS experiments, is the mecha-
nism involving sudden corrosion and removal of the carbon
support leaving Pt–Ni particles unanchored and causing
strongly enhanced particle movement by surface and bulk
diffusion until the detached particles have found neighboring
particles to agglomerate with (Fig. S5, ESI†).
Conclusions
We have presented STEM imaging of fuel cell catalyst activation
and degradation processes in an in situ electrochemical cell
supported by DEMS measurements. We investigated high activity
octahedral Pt–Ni nanoparticle fuel cell catalysts.
40
Our in situ
studies have revealed new insights into remaining key issues of
low temperature fuel cell catalysts including a more detailed
understanding of the degradation of octahedral PtNi alloy
catalysts such as carbon support corrosion, selective dissolu-
tion of non-noble metals, catalyst particle shape degradation,
and particle coarsening by coalescence and Pt redeposition.
The following issues were addressed: (1) degradation processes
were imaged on a time scale which allowed us to track changes
happening within a few seconds (e.g., rapid and sudden changes
when stepping into theconstantpotential and subsequent
stationary agglomerated state of the particles). (2) Agglomeration
along octahedral {111} facets duringanodicpotentialstepsasa
result of lined up particles during potential cycling. (3) Major
corrosion of the carbon support was happening during anodic
stepping, not during potential cycling. It is not possible to resolve
these differences in ex situ experiments. (4) By tracking the
dissolution of individual Ni nanoparticles we were able to deter-
mine atomic dissolution rates as a function of time and potential
cycle numbers which could provide guidelines for an optimized
voltammetric activation protocol of fuel cell catalysts.
In more detail: during catalyst activation, we observed the
nanometer-scale reaction dynamics of a selective Ni dissolution
process, observing that the Ni-rich particles become spongy
before fully dissolving. The Ni dissolution process does not take
place constantly but rather promptly after some electrochemical
cycles. We observe that the octahedral Pt–Ni alloy catalyst
remained morphologically stable during moderate potential
cycling up to +1.0 V
RHE
, while cycling to and holding at 1.2 V
RHE
and 1.4 V
RHE
caused increasingly severe coarsening. During
cyclic voltammetry to high potentials, we observed the electron
beam reduction of stringy deposits similar to Pt bands caused
by hydrogen cross-over in membrane electrode assemblies.
At high potential holds, Pt redeposition quickly obscures the
octahedral shape. Changes between potential cycling and
holds cause the most severe changes in the catalyst structure.
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Additionally, carbon corrosion was observed to increase particle
migration and coalescence, with Pt–Ni nanoparticles appearing to
typically coalesce on their {111} facets.
This study dramatically visualizes the dynamics of fuel cell
catalyst activation and degradation at the nanometer scale.
From these results we develop a better understanding of detri-
mental nanoscale effects which occur under different fuel cell
conditions and illustrate the urgent need for (1) more corrosion
stable support materials, (2) more oxidation stable alloy con-
figurations, and (3) the careful control of electrochemical
reaction conditions.
Author contributions
All authors conceived and designed the experiments. V. B.
carried out the chemical synthesis and the ex situ electro-
chemical experiments. M. E. H. performed the ex situ STEM
experiments and in situ image and movie processing. M. E. H.
and V. B. carried out the in situ STEM experiments. J. F. A.
planned, performed and evaluated the DEMS experiments.
All authors discussed the results, drew conclusions and parti-
cipated in writing the manuscript.
Conflicts of interest
There are no conflicts to declare.
Acknowledgements
Financial support was given by Deutsche Forschungsgemeinschaft
(DFG) grant STR 596/5-1 (‘‘Shaped Pt bimetallics’’). TEM Research
at Cornell was supported by the US Department of Energy
(DE-SC0019445). Elliot Padgett acknowledges support from an
NSF Graduate Research Fellowship (DGE-1650441). This work
made use of the Electron Microscopy Facilities at the Cornell
Center for Materials Research Shared Facilities which are
supported through the NSF MRSEC program (DMR-1719875).
We thank John Grazul and Mariena Silvestry-Ramos for help
with the electron microscopes.
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