Record activity and stability of dealloyed bimetallic
catalysts for proton exchange membrane fuel cells†
Binghong Han,
ac
Christopher E. Carlton,
ac
Anusorn Kongkanand,*
d
Ratandeep S. Kukreja,
d
Brian R. Theobald,
e
Lin Gan,
f
Rachel O'Malley,
e
Peter Strasser,*
f
Frederick T. Wagner
d
and Yang Shao-Horn*
abc
We demonstrate the unprecedented proton exchange membrane fuel cell (PEMFC) performance durability
of a family of dealloyed Pt–Ni nanoparticle catalysts for the oxygen reduction reaction (ORR), exceeding
scientific and technological state-of-art activity and stability targets. We provide atomic-scale insight into
key factors controlling the stability of the cathode catalyst by studying the influence of particle size, the
dealloying protocol and post-acid-treatment annealing on nanoporosity and passivation of the alloy
nanoparticles. Scanning transmission electron microscopy coupled to energy dispersive spectroscopy
data revealed the compositional variations of Ni in the particle surface and core, which were combined
with an analysis of the particle morphology evolution during PEMFC voltage cycling; together, this
enabled the elucidation of alloy structure and compositions conducive to long-term PEMFC device
stability. We found that smaller size, less-oxidative acid treatment and annealing significantly reduced Ni
leaching and nanoporosity formation while encouraged surface passivation, all resulting in improved
stability and higher catalytic ORR activity. This study demonstrates a successful example of how a
translation of basic catalysis research into a real-life device technology may be done.
Broader context
Reduction of costly Pt usage in proton exchange membrane fuel cell electrodes is one of the major challenges towards the development and commercialization
of fuel cell vehicles. In order for fuel cell vehicles to be economically viable, the US Department of Energy (DOE) has set specic performance targets for initial
fuel cell catalyst activity and fuel cell durability. While a number of different low-Pt catalyst concepts have met the initial-activity target using preliminary and
idealized disk electrode screening tests, few to none have met the initial-activity requirements in a realistic fuel cell device. No catalyst material has ever met the
demanding fuel cell durability targets. Here we report previously unachieved catalyst activity and device durability in fuel cells measured under automotive
testing conditions. The performance durability of this new fuel cell catalyst meets and exceeds the official 2017 DOE targets. This step-change in fuel cell catalyst
performance was made possible through a uniquely effective DOE-funded team of academia and industrial catalyst and MEA developers. This enabled a swi
translation of fundamental atomic-scale insights on catalyst synthesis, activity and degradation of dealloyed core–shell nanoparticles into industrial catalyst
material scale-up coupled to unprecedented device performance.
1. Introduction
Proton exchange membrane fuel cells (PEMFCs) represent a
means of converting renewable chemical energy into electrical
energy with high theoretical efficiency and good enough power
density for it to be of great interest to automotive applications.
Due to the sluggishness of the oxygen reduction reaction (ORR) at
the cathode of PEMFCs, Pt-based catalysts are required to achieve
reasonable overpotentials at realistic current densities.
1
Although
Pt nanoparticles supported on high surface area carbon have
traditionally been used as the ORR catalyst, high Pt loadings
needed to achieve sufficient power densities for automotive
demands make them unsuitable for commercialization due to the
preciousness and limited supply of Pt.
1,2
Studies of bimetallic Pt–
M surfaces and nanoparticles, where M is a transition metal, have
demonstrated several times higher intrinsic ORR activity than
a
Department of Materials Science and Engineering, Massachusetts Institute of
b
Department of Mechanical Engineering, Massachusetts Institute of Technology,
Cambridge, Massachusetts 02139, USA
c
Electrochemical Energy Laboratory, Massachusetts Institute of Technology,
Cambridge, Massachusetts 02139, USA
d
Electrochemical Energy Research Lab, General Motors Powertrain, Honeoye Falls, NY
14472, USA. E-mail: anusorn.kongkanand@gm.com
e
Johnson Matthey Technology Centre, Blount's Court, Sonning Common, Reading, RG4
9NH, UK
f
The Electrochemical Energy, Catalysis and Materials Science Laboratory, Department
of Chemistry, Chemical Engineering Division, Technical University Berlin, Berlin
†Electronic supplementary information (ESI) available. See DOI:
10.1039/c4ee02144d
Cite this: Energy Environ. Sci.,2015,8,
258
Received 10th July 2014
Accepted 12th September 2014
DOI: 10.1039/c4ee02144d
www.rsc.org/ees
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pure Pt.
3–17
This catalytic enhancement has been attributed to
optimizing the oxygen absorption energy of surface Pt through
strain
5,10,11,18–20
and ligand
6–8,11,21–23
effects in a Pt sandwich segre-
gation structure induced by annealing
3,4,7,24,25
or a Pt skeleton
structure induced by acid leaching of transition metal.
3,4,7,18,26,27
Although supported Pt–M nanoparticles have shown greater
stability than supported Pt nanoparticles,
5–9,17,27,28
considerable
losses are noted in the specic and mass activities of Pt–Mnano-
particle catalysts during fuel cell operation.
28–32
Theactivitylosscan
result from the dissolution and loss of Pt from the cathode,
33,34
reductionofstrainandligandeffects associated with the disso-
lution of transition metal,
28,32,35,36
re-deposition of dissolved Pt on
the particle surface,
13,28,33,35,36
and Erlebacher-like
37–41
dealloying.
Previous work has shown that the Pt–M catalyst stability is highly
dependent on the composition, size and morphologies of catalyst
nanoparticles.
4,16,26–28,35,42–47
Of signicance, previous work has
shown that having Ni or Co compositions greater than 50 at.% in
Pt
x
Ni or Pt
x
Co nanoparticles could lead to the formation of spongy
particles with nano-scale porosity due to severe metal leaching
26,28
while small particles below certain critical size could avoid porosity
formation.
16,42,43
However, the effects of catalyst preparation such
as the acid-leaching protocol and post-acid-leaching heat treat-
ment, on metal leaching and surface porosity formation are still
poorly understood. Although recent studies show that de-alloyed
Pt-based catalysts can exhibit remarkably high activities,
48–50
these
catalysts have been examined using rotating disk electrode (RDE)
measurements but not by membrane electrode assembly (MEA)
testing, which is critical to the development of efficient, durable
MEAs for PEMFCs.
In this work, we have systematically investigated the inu-
ence of precursor morphology, acid-leaching conditions and
post-acid-treatment annealing on nanoporosity and surface
passive layer formation of Pt–Ni nanoparticles, as well as the
consequent evolution of their compositional, electrochemical
and morphological properties during PEMFC operation. What
sets this work apart from previous studies
7,9,14,24,48–50
is the
combination of systematic changes in the catalyst preparation
conditions, MEA testing and durability data with detailed
transmission electron microscopy (TEM) characterization
before and aer PEMFC testing. This combination made a step-
change in fuel cell catalyst performance and enabled a swi
translation of fundamental atomic-scale insights on dealloyed
core–shell nanoparticles into industrial catalyst material scale-
up coupled to unprecedented device performance. Our results
show that smaller particle size, non-oxidative acid treatment
and post-acid-treatment annealing can reduce transition metal
leaching from catalyst nanoparticles, and suppress the nano-
porosity formation. This provides insights into the design of Pt–
M nanoparticle catalysts with enhanced stability and activity. As
a result, we demonstrate record PEMFC cathode activity and
durability exceeding scientic and technological targets.
51,52
2. Experimental
2.1 Catalyst and electrode preparation
Two PtNi
3
catalyst precursors (P1 and P2) supported on high-
surface-area carbon were developed using highly scalable
synthesis routes and provided by Johnson Matthey Fuel Cells
Ltd (UK), which showed very similar shapes and Pt : Ni ratios
but signicantly different particle size distributions (PSDs)
53
having P2 with smaller particles size (5 nm) with a narrower
PSD, compared to P1 with larger particles of 8 nm (Fig. S1†).
Here PtNi
3
was chosen as the precursor for a number of reasons.
First, the initial 75 at.% Ni content is safely above the parting
limit of Pt–Ni solid solution alloys and thus ensures the exis-
tence of a critical dealloying potential at which Ni selectively
leaches from the precursor particles. Consistent with earlier
studies,
26,44
this dealloying process results in the formation of
high surface area core shell nanoparticle catalysts with compact
Pt skin. A series of different acid-leaching conditions and post-
dealloying treatments were applied to these precursors at GM,
as shown in Fig. 1. The precursors were either treated in (1) 1 M
nitric acid at 70 C for 24 h in air (marked as P1-NA and P2-NA)
or (2) 0.5 M sulfuric acid at 80 C for 24 h in nitrogen (P2-SA).
The resultant catalysts were then thoroughly washed with DI
water and dried in nitrogen at room temperature. Aer the
dealloying, some P2-SA catalysts were thermally annealed at 400
C for 4 h in 5% H
2
and 95% N
2
,
54
marked as P2-SA-AN. For
comparison, 46 wt% Pt supported on Vulcan carbon from TKK
(Tanaka Kikinzoku K.K., Japan) was used.
Catalyst-coated membranes (CCMs) of 50 cm
2
were used in
fuel cell testing, which were prepared using the decal transfer
method, with a loading of 0.1 mg
Pt
cm
2
in the cathode and 0.05
mg
Pt
cm
2
in the anode. Naon® D2020 ionomer, with an
equivalent weight of 950 g equiv.
1
, was utilized with an ion-
omer to carbon weight ratio of 0.95 and 0.6 in the cathode and
anode, respectively. Catalysts derived from P1 and P2 precur-
sors were studied in the cathode while 20 wt% Pt/C was used as
the anode catalyst. Naon® NRE211 (25 mm thick) was used as
the membrane. The gas diffusion medium used was carbon
ber paper backing (200 mm) coated with a microporous layer
(MPL, 30 mm). The detailed procedure for the fabrication of
electrodes and MEA can be found in previous work.
1
2.2 Electrochemical measurements
The MEAs were conditioned in a fuel cell for about 7 h followed
by a series of electrochemical measurements to give a
Fig. 1 Schematic showing different catalysts studied in this paper. The
precursors were either treated in (1) 1 M nitric acid at 70 C for 24 h in
air (P1-NA and P2-NA) or (2) 0.5 M sulfuric acid at 80 C for 24 h in
nitrogen (P2-SA). After the dealloying, some P2-SA catalysts were
thermally annealed at 400 C for 4 h in 5% H
2
and 95% N
2
, marked as
P2-SA-AN.
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conditioned MEA aer a total run time of 17 h. To accelerate the
degradation of the cathode catalyst, some MEAs were then
subjected to voltage cycling tests which were performed in 200/
50 standard cubic cm per min of H
2
/N
2
at 80 C, 100% RH and
101 kPa
abs
for up to 30 000 (30k) cycles at a scan rate of 50 mV
s
1
between 0.6 and 1.0 V vs. RHE (30k-cycled). Averages from at
least 4 MEAs with standard deviations are reported.
At each aging stage, electrochemical properties such as mass
and area-specic ORR activities, and Pt electrochemically active
surface area (ECSA) were measured. The mass and specic
activities of oxygen reduction reaction at the cathode were
calculated at 0.9 V vs. the reversible hydrogen electrode (RHE),
which was obtained from the H
2
/O
2
polarization curve collected
at 80 C, 100% relative humidity (RH) and 150 kPa
abs
aer high
frequency resistance-correction.
1
Polarization curves in H
2
/air
were obtained at 80 C, 100% RH, and 170 kPa
abs
. Here we
assume that the Pt loading did not change during electro-
chemical measurements, and the initial Pt loading in the MEA
was used in the mass activity calculation. The ECSAs were
obtained from the H adsorption/desorption (HAD) peaks using
cyclic voltammetry.
55
Although it has been shown in single-
crystal electrode studies that the hydrogen adsorption charge
per surface area (the conversion constant used in the HAD
method) would decrease with increasing Co or Ni content in Pt–
Co or Pt–Ni alloys due to a modied electronic structure,
56–58
meanwhile the electrochemical CO stripping might give a
different ECSA than those measured by HAD,
48
both methods
gave comparable results for all of our MEA studies with CO
giving a 5–20% higher value than HAD does.
57
Therefore, only
ECSAs measured by H adsorption/desorption were reported in
this work, ignoring the inuence of Ni content change on the
HAD conversion constant. The evolution of ECSA before and
aer 30k cycles is shown in Fig. S2 in the ESI.†
2.3 Electron probe micro-analysis
Cross-sectioned MEA samples were studied by electron probe
micro-analysis (EPMA) using an SX100 electron probe micro-
analyzer from CAMECA Instruments, Inc. Small pieces of the
sample, approximately 4 10 mm, were cut from the center of
the MEAs and mounted in brominated epoxy under vacuum.
Aer allowing the epoxy to harden, the samples were polished to
a mirror nish manually using standard polishing papers fol-
lowed by lapping using 1-micron diamond suspension. A thin
layer (3 nm) of Au–Pd alloy was then sputter coated immedi-
ately aer the nal lapping to provide surface electron
conductivity needed for EMPA. The cross-sectioned MEAs were
analyzed with an electron probe of 15 kV and 10 nA. An area of
75 75 mm covering the anode MPL, anode, membrane,
cathode and cathode MPL, was scanned with a step size of 0.3
mm at 10 ms per step to generate elemental maps of Pt and Ni.
Wavelength dispersive spectrometers were used to collect X-rays
of Pt-Maand Ni-Kalines using the thallium acid phthalate
crystal and large lithium uoride crystal, respectively. At least
three such maps were generated from different locations on
each sample. X-ray intensities from pure Pt and Ni were deter-
mined to calculate k-ratios from X-ray intensities in band
proles constructed from X-ray maps, thus allowing conversion
to weight/atomic percent concentration proles. The concen-
tration proles thus generated were used for quantication and
distribution analyses of Pt and Ni atomic percentages. Induc-
tively coupled plasma-atomic emission spectrometry on the
non-active-area membranes was also performed to quantify the
amount of Pt and Ni that migrated away from the active area.
2.4 Transmission electron microscopy
TEM samples of cathode catalysts in the MEA were prepared by
using a razor to remove the carbon paper diffusion layer, care-
fully cleaning the MPL offof the membrane and scraping the
membrane directly into a scintillation vial. The cathode scrap-
ings were sonicated briey in ethanol (Sigma-Aldrich, r 99.5%),
and then applied to TEM grids (Electron Microscopy Science,
lacy carbon/copper lm 200 mesh). All TEM studies were per-
formed on a JEOL 2010F microscope equipped with an ultra-
high resolution polepiece, having a point resolution of 0.19 nm.
For each sample, over 100 catalyst nanoparticles under bright-
eld-diffraction-contrast imaging were used to determine the
PSDs. The bright eld diffraction contrast images were formed
by aligning a 20 mm aperture with the transmitted beam, and
captured by a CCD camera. The digital images were analyzed
using Gatan Digital Micrograph v2.01 (Gatan Inc.) and ImageJ
v1.44p (National Institute of Health, USA). High-resolution TEM
(HRTEM) images were obtained without an objective aperture
and were analyzed using a Gatan Digital Micrograph. High
angle annular dark eld (HAADF) scanning TEM (STEM) images
were formed by aligning the Ronchigram with optic axis and a
50 mm condenser aperture with the center of the Ronchigram.
An electron beam probe size of 0.5 nm and a camera length of
40 cm were used to minimize the effect of diffraction contrast
on the HAADF images. Parallel-beam and STEM energy
dispersive spectroscopy (EDS) data were collected from indi-
vidual catalyst particles and analyzed using the Inca (Oxford
Instruments) soware. The Ni and Pt atomic compositions were
determined from the Ni-K series and the Pt-L series using the
INCA soware as described in previous work.
35
For parallel-
beam EDS measurements, three different spots with a diameter
of 200 nm (totally containing over 100 particles) were used to
collect the average atomic percentage of Ni for each sample. The
uncertainty in the EDS-determined atomic percentage is 10%
of the measured value.
3. Results and discussion
3.1 Pt
x
Ni PEM single cell performance activity and durability
PEMFC polarization curves of cathodes with P1-NA or P2-NA are
compared with that of a commercial Pt/C catalyst (from TKK) in
Fig. 2a. Despite the fact that P1-NA and P2-NA cathodes
(0.1 mg
Pt
cm
geo.
2
) had a Pt loading which is four times lower
than that of the commercial Pt/C (0.4 mg
Pt
cm
geo.
2
), P1-NA and
P2-NA showed comparable voltage proles as a function of
geometric current density. Aer 30k cycles (0.6–1.0 V vs. RHE) of
accelerated voltage cycle testing, both P1-NA and P2-NA showed
noticeable voltage loss (dashed lines), with smaller voltage loss
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for the P2-NA cathode compared to P1-NA aer cycling, which
indicates lower mass activity loss. It is interesting to note that
the voltage loss at high current densities can be greatly miti-
gated by lowering the upper voltage limit to which the voltage
cycling was performed from 1.0 V to 0.925 V despite more
extensive cycles (Fig. S3†). The smaller loss may be attributed to
decreases in Pt surface area loss, mass activity loss, Ni and Pt
dissolution, and carbon corrosion.
59
The specic ORR activity of
the P1-NA cathode was reduced by 1/3 aer 10k cycles
(Fig. 2b), which was further reduced from 10k to 30k cycles. In
contrast, the specic activities of P2-NA and the other P2-
derived cathodes were found to increase with PEMFC cycling,
where activities greater than 1.7 mA cm
pt
2
were obtained aer
30k cycles. Moreover, the mass activity of the P1-NA catalyst
reduced considerably during PEMFC cycling (Fig. 2c), where 1/
2 and 2/3 of the activity found in the conditioned MEA was lost
aer 10k and 30k cycles, respectively. On the other hand, the P2-
NA cathode as well as the other P2-derived catalysts showed
much smaller loss in mass activity than P1-NA during cycling.
Of signicance, all P2-derived catalysts surpassed the US
Department of Energy 2017 PEMFC cathode catalyst target
(initial activity of conditioned MEA >0.44 A mg
Pt
1
and <40%
activity loss aer 30k voltage cycling),
51
which has not yet been
reported for any catalysts and sets the record for MEA catalyst
mass activity and stability, as shown in Fig. 2d.
Previous studies have shown that the reduced specic
activity loss of Pt-alloy catalysts can be attributed to transition
metal dissolution
28,32,35,36
while mass activity loss can result
from loss in the specic activity and electrochemically active
surface area due to Pt dissolution,
33,60
coarsening
34,60
and coa-
lescence of Pt nanoparticles.
34,35
We will examine and discuss
differences in the Ni atomic percentage change in the MEA,
particle size, morphologies and Ni atomic percentage distribu-
tion within individual particles of P1-NA and P2-derived cata-
lysts before and aer PEMFC cycling in order to explain their
markedly different stability in PEMFCs.
3.2 Catalyst composition changes inside the MEA cathode
aer PEMFC testing
Examining the average Ni atomic percentage of catalyst particles
in CCMs using EMPA, we found that the P1-NA catalyst had much
less Ni than P2-derived catalysts in the pristine CCMs, as shown in
Fig. 3a. In addition, sulfuric acid appeared to remove less Ni from
the P2-derived catalysts than nitric acid. PEMFC conditioning
testing further reduced the Ni atomic percentage in these cata-
lysts, where the largest loss of Ni was found for P1-NA. Moreover,
PEMFC cycling led to signicant Ni loss in all catalysts, where P2-
SA-AN was found to have the highest Ni percentage (15 at.%)
while P1-NA had the least Ni percentage (5at.%)aer 30k cycles.
These various Ni leaching results are different from a previous Pt–
Fig. 2 (a) Fuel cell polarization curves of MEAs with P1-NA and P2-NA cathodes at different aging stages (conditioned and 30k-cycled), and of
conditioned Pt as a reference. The catalyst loadings were 0.1 and 0.4 mg
Pt
cm
2
for the dealloyed PtNi and commercial Pt catalysts, respectively.
Testing conditions of the anode/cathode were H
2
/air, 100/100% RH, 170/170 kPa
abs
back pressure, and stoichiometric coefficient of 1.5/2,
respectively. (b) Specific activity and (c) mass activity of the catalysts made using P1 and P2. The activities were measured at 0.9 V (vs. RHE) from
the H
2
/O
2
polarization curve collected at 80 C, 100% relative humidity (RH) and 150 kPa
abs
after high frequency resistance-correction, where
dark grey, light grey and white represent conditioned, 10k-cycled and 30k-cycled MEA, respectively. From (a) to (c) the error bars represent the
standard deviations of activity measurements among at least 4 MEA samples. (d) The end-of-life (EOL, 30k-cycled MEA) vs. beginning-of-life
(BOL, conditioned MEA) mass activity for catalysts from this work and from previous work on MEAs: 3 M NSTF PtCoMn, NSTF Pt
3
Ni
7
,
61
Pt
3
Co/CNC
and Pt
2
Ni/CNC,
62
Supplier Pt Alloy and pure Pt.
52,63
The solid and dash horizontal lines in (c) and the red star and the green square regions
represent the DOE 2017 PEMFC cathode catalyst targets for BOL mass activity (0.44 A mg
Pt
1
) and EOL mass activity (0.26 A mg
Pt
1
),
51
respectively.
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Co study, where the remaining Co at.% reached a constant value
aer HClO
4
treatment irrespective of the initial composition,
implying that the different particle size, acid-treatment and
annealing conditions have a greater inuence than initial Pt
content on metal leaching.
64
As shown in Fig. 3b, the mass
activities of P1- and P2-based catalysts before and aer cycling
were found to largely correlate with the Ni atomic percentage in
the cathode. The P2-derived catalysts showed the maximum mass
activities with 30 at.% of Ni, which is consistent with previous
studies.
65
However, the decrease in mass activities aer 30k cycles
was largely governed by the decrease in Pt surface area (Pt surface
area decreased from 45 to 25 m
2
g
Pt
1
for the conditioned and
30k-cycled samples regardless of the catalysts, as shown in
Fig. S2†). In contrast, the specic activities (the activity per surface
area of Pt) of these catalysts were found largely independent of Ni
atomic percentage in the MEA, as shown in Fig. 3c, which high-
lights the importance of the near surface composition on the
catalytic activity.
26
Here it is important to note that the estimation
of ECSA using the HAD method may be affected by the leaching of
Ni,
56
and, therefore, Fig. 3c highlights the trends of activity change
aer cycling. Moreover, synchrotron X-ray absorption studies of
these catalysts showed that the composition of the subsurface Ni
could greatly affect the distribution of different adsorbed oxygen
species and hence the ORR activity.
66
3.3 Correlating particle size, morphology and composition
of dealloyed nanocatalysts in the CCM
P1-NA catalysts were found to have larger particle sizes (7 nm)
with a broader particle distribution (s3 nm) than P2-NA (5 nm
1 nm) from TEM analysis, as shown in Fig. 3d. Here s
represents the standard deviation of PSD of more than 100
particles for each sample. At the same time, P2-NA, P2-SA and
P2-SA-AN showed similar particle sizes of 5 nm, implying that
the size effect was not the dominating factor among these
particles compared with acid-treatment and annealing effects.
The EDS data collected from at least 3 different locations
(200 nm-diameter circular areas) totally containing over 100
nanoparticles showed that P1-NA contained less Ni (19 6.9
at.%) than P2-NA (25 3 at.%), which is in good agreement with
the average Ni atomic percentages obtained from EPMA results
in Fig. 3a. In addition, while both P1-NA and P2-NA consist of
both solid and porous particles, P1-NA exhibits a higher degree
of nanoporosity than P2-NA, which has been observed in
different Pt–M nanoparticles and different acid-leaching
conditions in previous studies as “spongy particles”.
16,28,35,67
The
formation of these percolated nanoparticles is typically
explained by dealloying with a critical size, which is dened as
the lower limitation of particle size for nanoporosity formation
under given leaching conditions and particle composition.
35,37,38
In P1-NA, the porosity could be found from particles as large as
20 nm and as small as 6 nm (see the top right inset of Fig. 4a),
indicating a critical size of 6 nm, which was smaller than the
critical size reported by previous work
16,42,43
ranging from 10–
30 nm. These ndings further conrm that P1-NA experienced
more extensive Ni leaching than P2-NA, which may result from
the presence of larger particles. The atomic origin of the
porosity formation in alloy particles above the critical size is
associated with the competition between the relative rates of
metal dissolution and Pt surface diffusion, which favors
Fig. 3 (a) EPMA-determined Ni atomic percentages, where black, dark grey, and white represent pristine CCM, conditioned and 30k-cycled
MEA, respectively. (b) and (c) The average Ni atomic percentages determined by EPMA plotted against the mass and specific activities of the
MEAs, respectively (dark grey: conditioned; white: 30k-cycled). From (a) to (c), the activity error bars represent the standard deviation of activity
measurements among at least 4 MEA samples, while the Ni-content error bars represent the standard deviation of EMPA measurements among
at least 3 different locations. (d) TEM-determined mean particle size from different MEAs, where black, dark grey and white represent pristine
CCM, conditioned MEA and 30k-cycled MEA, respectively. For each sample over 100 particles were measured to get the average size and the
error bars represent the standard deviation.
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Rayleigh surface instabilities resulting in bulk dealloying.
37–41
Large particles that are above the Rayleigh critical length will
have an unstable surface and become more vulnerable to acid
leaching. Assuming this mechanism, we would expect the crit-
ical particle size decrease as the acid leaching conditions
become more oxidative, because this would increase the Ni
dissolution rate. The observation of these small porous particles
in this study can be attributed to the fact that we used aerated
hot oxidative acid (HNO
3
), which is more oxidative compared to
room temperature non-oxidative acid (HClO
4
or H
2
SO
4
)in
previous studies.
16,42,43
The lower nanoporosity in P2-NA might
be attributed to most P2-NA nanoparticles having an average
size of 5 nm, below the critical size for nanoporosity formation
provided that 6 nm was the critical size.
We also found that leaching in H
2
SO
4
in N
2
, which is
substantially milder in oxidation than the nitric acid, could
facilitate the development of a Pt-rich skin near the
surface
3,4,16,28,35
and reduce the nickel leaching. P2-SA in the
CCM showed a much higher Ni content than P2-NA in the CCM
(Fig. 4c and e). In addition, P2-SA had nanoparticles with dense
core–shell structures with Pt-rich surface layers in the HAADF
STEM image for all observed particles, which is supported by
EDS analysis, as shown in Fig. 4f. Such Pt-rich surface layers
were also observed in P2-SA-AN particles (Fig. 4g and h). In
contrast, such Pt-rich shells near the particle surface were not as
noticeable in the P2-NA.
3.4 Particle size, morphological and compositional changes
upon PEMFC conditioning and cycling
A small increase of 1 nm in the average particle size was found
for P1-NA and P2-derived catalysts aer PEMFC conditioning
and cycling (except for P2-SA whose size rst decreased aer
conditioning and then increased aer cycling), as shown in
Fig. 4 and 5 and S2–S4,†but the small size changes were still
within experimental uncertainty (the standard deviation of
particle sizes is 1 to 3 nm). Area EDS analysis revealed
signicant Ni loss from P1-NA and little Ni remained in catalyst
particles aer PEMFC conditioning and cycling (Fig. S4a†and
5a), which is in good agreement with EPMA results discussed
previously. P2-NA was found to exhibit less Ni loss from area
EDS analysis during PEMFC conditioning and aer 30k cycles.
Remarkably, P2-SA and P2-SA-AN retained considerable Ni aer
conditioning and even aer 30k cycling, with atomic percent-
ages of 24 7% and 31 5%, respectively (Fig. 5e and g).
Pronounced particle morphology changes were observed in
P1-NA upon conditioning, where the big spongy nanoparticles
shown in Fig. 4a disappeared and were replaced by donut-like
nanoparticles shown in Fig. S4a.†Similar spongy-particle
transformation has been found in previous work,
16,28,35,67
which
might be caused by the dissolution of highly porous particles
and Pt redeposition to form Pt-enriched-shell particles. As the
intensity in the HAADF images is proportional to the product of
the thickness and Z
2
(Zbeing the atomic number), the EDS data
in Fig. S4b†demonstrate that the center and edge have similar
Pt : Ni ratios, the lower brightness near the particle center
shown in Fig. S4b†can be attributed to the formation of donut-
shaped particles. In contrast, combined HAADF analysis of
conditioned (Fig. S4d†) and 30k-cycled P2-NA particles (Fig. 5d)
and accompanied point EDS analysis of Pt and Ni atomic
percentages (having a lower Pt concentration in the center)
revealed that the core thickness of cycled P2-NA could be greater
than particle surface regions, and thus the contrast shown in
the HAADF image (Fig. S4d†) could be attributed to Pt-rich
shells rather than a donut-like shape. The formation of Pt-rich
shells can result from bulk dealloying,
37,38
where leaching of
transition metal is conned to particle surface regions.
Fig. 4 Bright-field diffraction contrast TEM images with PSDs, Ni
atomic percentages determined from parallel-beam EDS (left) and
HAADF STEM images with point-mode EDS (right) from catalyst
particles taken from pristine CCMs. (a) and (b) P1-NA, (c) and (d) P2-NA,
(e) and (f) P2-SA, (g) and (h) P2-SA-AN. Here srepresents the standard
deviation of PSD from at least 100 measurements and the standard
deviation of EDS results among 3 different locations. The insert on the
top right corner is the HRTEM image of a nanoparticle with a diameter
of 5 nm that clearly shows porosity from CCM.
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P1-NA exhibited much worse durability in PEMFC cycling
than P2-NA, which can be attributed to extensive porosity found
in P1-NA. First, P2-NA maintained a higher Ni content which is
twice that of P1-NA both aer conditioning and aer 30k cycles
(Fig. S4a and S4c,†Fig. 5a and c). Second, P1-NA showed only
1/3 of its original mass activity and 1/2 of its original specic
activity aer 30k cycles while P2-NA showed almost no change
in mass and specic activity considering the measurement
uncertainty (Fig. 2b and c). The much reduced activity of P1-NA
in PEMFCs could be attributed to a greater Ni loss near the
surface regions, which reduced the surface electronic structure
modications of Pt atoms induced by transition metal
alloying.
5–8,10,11,19–23
The development of a Pt-rich passivation
layer is expected to improve the resistance to base metal
leaching and activity loss, which might be another reason for
the higher durability of P2-NA than P1-NA.
Point EDS analysis revealed that P2-SA-AN nanoparticles
retained considerably more Ni near surface regions than P2-SA
aer conditioning and 30k cycles (Fig. S4e and S4g,†Fig. 5e and
g), where no noticeable changes in the Ni atomic percentages
near surface regions aer conditioning and 30k cycles were
observed. The increase in stability with the thermal annealing
observed in MEAs in our study is in good agreement with
previous RDE work by Wang et al.
68
It is proposed that the
exceptional stability and high mass and specic activities of P2-
SA-AN in PEMFCs can be attributed to the ability of surface
regions to retain Ni, which might be critical to maintain
modied electronic structures of surface Pt atoms associated
with Ni necessary to provide enhanced mass and specic
activities relative to pure Pt nanoparticle catalysts. This
hypothesis is supported by HAADF in conjunction with electron
energy loss spectroscopy on P2-SA and P2-SA-AN aer condi-
tioning, which revealed that Pt shell thicknesses (1.2 0.6 and
1.3 0.3 nm, respectively) were nearly unchanged (1.6 nm)
aer 30k cycles, as shown in Fig. S5.†Further support, from an
X-ray absorption technique that could differentiate Pt shell
structures on these samples, will be published elsewhere. This
again proves the importance of a surface shell to the activity and
stability of dealloyed Pt-alloy nano-catalysts, and indicates that
annealing could be an effective process to improve the stability
of Pt–metal core–shell catalysts.
Lastly, it is useful to note that when a PtCo
3
precursor was
prepared using an analogous synthesis method to P2 (PtNi
3
), we
were able to demonstrate an active and durable dealloyed
catalyst comparable to those described in this study.
53
Consid-
ering this, other binary and ternary systems may also be effec-
tive. This illustrates the versatility of the dealloying approach
and allows one to weigh in other selection criteria such as
supply chain, manufacturability, and the interaction with other
fuel cell components when developing a new catalyst.
Conclusions
We have demonstrated unprecedented PEMFC device perfor-
mance durability employing a family of dealloyed PtNi
3
nano-
particle catalysts, exceeding scientic and technological state-
of-art activity and stability targets. We nd that for dealloyed
catalysts larger particle size could strongly promote the forma-
tion of nanoporosity, which is associated with the rapid
reduction in Ni content. On the other hand, reducing the
particle size below a certain critical value can reduce porosity
generation and promote the formation of a Pt-rich passivation
particle shell layer. We also nd that non-porous core–shell
nanostructures can be strongly encouraged by using a less
oxidative acid leaching protocol, further improving base metal
retention and leading to exceptional activity durability. Finally,
post-acid-treatment annealing can cause the formation of
comparatively dense-Pt shells, which can provide excellent
resistance to Ni leaching, promote particle size stability, and
Fig. 5 Bright-field diffraction contrast TEM images with PSDs, Ni
atomic percentages determined from parallel-beam EDS (left) and
HAADF STEM images with point-mode EDS (right) for catalyst particles
taken from 30k-cycled MEAs. (a) and (b) P1-NA, (c) and (d) P2-NA, (e)
and (f) P2-SA, (g) and (h) P2-SA-AN. Here srepresents the standard
deviation of PSD from at least 100 measurements and the standard
deviation of EDS results among 3 different locations.
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can help to improve both specic and mass activities during
electrochemical cycling. These results highlight the critical role
of acid treatments, an aspect of catalyst design that has not
been rigorously controlled to date, in creating Pt-rich passiv-
ation layers to prevent base metal leaching and improve cata-
lytic activity.
Acknowledgements
This work was supported by the U.S. Department of Energy,
Office of Energy Efficiency and Renewable Energy under grant
DE-EE0000458. The authors thank Eric L. Thompson, Jeanette
E. Owejan, Thomas E. Moylan, Travis Downs, Peter Harvey, and
Mike Flanagan of GM for engineering support. The research
made use of the Shared Experimental Facilities supported by
the MRSEC Program of the National Science Foundation under
award number DMR 08-019762. P. S. acknowledges partial
nancial support by the German National Science Foundation
(Deutsche Forschungsgemeinscha) under grant STR 596/5-1.
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