ARTICLE
Received 18 Apr 2015
| Accepted 16 Feb 2016 | Published 23 Mar 2016
Water electrolysis on La
1 x
Sr
x
CoO
3 d
perovskite
electrocatalysts
J. Tyler Mefford
1
, Xi Rong
2
, Artem M. Abakumov
3,4
, William G. Hardin
5
, Sheng Dai
6
, Alexie M. Kolpak
2
,
Keith P. Johnston
5,7
& Keith J. Stevenson
1,3,5,8
Perovskite oxides are attractive candidates as catalysts for the electrolysis of water in alkaline
energy storage and conversion systems. However, the rational design of active catalysts has
been hampered by the lack of understanding of the mechanism of water electrolysis on
perovskite surfaces. Key parameters that have been overlooked include the role of oxygen
vacancies, B–O bond covalency, and redox activity of lattice oxygen species. Here we present
a series of cobaltite perovskites where the covalency of the Co–O bond and the concentration
of oxygen vacancies are controlled through Sr
2 þ
substitution into La
1 x
Sr
x
CoO
3 d
.
We attempt to rationalize the high activities of La
1 x
Sr
x
CoO
3 d
through the electronic
structure and participation of lattice oxygen in the mechanism of water electrolysis as
revealed through ab initio modelling. Using this approach, we report a material, SrCoO
2.7
,
with a high, room temperature-specific activity and mass activity towards alkaline water
electrolysis.
DOI: 10.1038/ncomms11053
OPEN
1
Department of Chemistry, The University of Texas at Austin, Austin, Texas 78712, USA.
2
Department of Mechanical Engineering, Massachusetts Institute of
Technology, Cambridge, Massachusetts 02139, USA.
3
Center for Electrochemical Energy Storage, Skolkovo Institute of Science and Technology, 143026
Moscow, Russia.
4
Electron Microscopy for Material Science, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium.
5
Texas Materials
Institute, The University of Texas at Austin, Austin, Texas 78712, USA.
6
Chemical Sciences Division, Oak Ridge National Laboratory, Oak Ridge,
Tennessee 37831, USA.
7
Department of Chemical Engineering, The University of Texas at Austin, Austin, Texas 78712 , USA.
8
Center for Nano and
Molecular Science and Technology, The University of Texas at Austin, Austin, Texas 78712, USA. Correspondence and requests for materials should be
addressed to K.J.S. (email: K.Stevenson@skoltech.ru).
NATURE COMMUNICATIONS | 7:11053 | DOI: 10.1038/ncomms11053 | www.nature.com/naturecommunications 1
T
he scarcity of fossil fuels and the increasing awareness of
the environmental and geopolitical problems associated
with their use have encouraged significant efforts towards
the development of advanced energy storage and conversion
systems using materials that are cheap, abundant and
environmentally benign. A major thrust in the field of renewable
energy has been to develop higher power and more energy-dense
storage devices, including low-temperature regenerative fuel cells
and rechargeable metal–air batteries that function through
the electrocatalysis of oxygen. Inherent to these systems are the
electrolysis of water (2H
2
O-O
2
þ4H
þ
þ4e
; oxygen
evolution reaction (OER)) and the reduction of molecular oxygen
(O
2
þ4H
þ
þ4e
-2H
2
O; oxygen reduction reaction (ORR)),
both of which require the use of an electrocatalyst due to their
slow reaction kinetics. The most active catalysts for the ORR
are Pt-alloys and other precious metals, Ir, Ru and Pd
1–3
.
However, while the Pt group metals perform well for the ORR,
the formation of an oxide surface film at high potentials,
especially in the case of Pt, decreases their ability to catalyse the
OER
4
. This problem, coupled with the Pt group metal scarcity
and restrictive cost represent major roadblocks to mass adoption
of fuel cells and metal–air batteries in renewable energy
technologies.
Using alkaline electrolytes opens up the possibility to use
transition metal oxides as catalysts due to their structural stability,
resistance to electrolytic corrosion and their high activities for
both the OER and ORR
5–7
. Among the wide variety of metal
oxides available, the crystal family of perovskite oxides ABO
3±d
,
of which A is a rare-earth or alkaline earth element and B is a
transition metal, are attractive candidates due to their high
ionic and electronic conductivities, structural stability, and the
ability to substitute into the A and B sites elements of varying
valency, electronegativity or ionic size to tune the structural,
physical and electronic properties of the catalyst. Even though the
electrolysis of water to oxygen is one of the most extensively
studied reactions, predating even the fields of catalysis and
electrochemistry, the lack of a conclusive mechanism for metal
oxides in alkaline electrolyte remains a significant limitation in
the rational design of electrocatalysts for the OER
8
. Thus, much
of the research on perovskites for the OER and ORR has been
focused on identifying descriptors for the activities of perovskites
based on the electronic and structural properties of the surface or
bulk
9–11
. Since the initial discovery of La
0.8
Sr
0.2
CoO
3
as an active
ORR catalyst, many mechanistic theories have been put forward
over the past 40 years
12
. A recent review summarizes the current
understanding of mechanistic processes for the OER, specifically
highlighting correlations between bulk and surface properties of
metal oxides and their electrocatalytic activities
13
. Notably, the
idea that the e
g
filling of the transition metal in the ABO
3
perovskite controls the intermediates binding strength and thus
the electrocatalytic activity has recently gained significant
credence
14,15
. However, we have observed that among a series
of perovskites with a nominal e
g
filling of B1 (LaBO
3
, where
B ¼Mn, Co, Ni, or Ni
0.75
Fe
0.25
), there exists significant
differences in their activities for both the ORR and the OER,
indicating that the surface chemistry may not be adequately
rationalized by bulk electronic descriptions
16,17
.
A previously overlooked parameter concerns the role of oxygen
vacancy defects, which allows for crystalline oxygen to be mobile
at the surface of perovskites. It is well-known that the
stoichiometry of oxygen in the crystal structure of perovskites
often differs from the nominal value of 3 for the formula ABO
3
,
affecting both the lability of surface oxygen and reflecting the
underlying electronic structure of these materials
18–20
. The degree
of vacancy formation reflects the relative positions of the
transition metal 3d bands compared with the oxygen 2p band
in the crystal, with more covalent systems exhibiting higher
vacancy concentrations as shown in Fig. 1. In addition, it is
well-documented that the concentrations of oxygen vacancies in
perovskite electrodes can be controlled through an applied
electrical potential, with room temperature diffusion coefficients
of lattice oxygen for a number of perovskites in the range
of 10
14
to 10
11
cm
2
s
1
(refs 21–26). In a previous paper,
we demonstrated that this effect could be used as a means of
pseudocapacitive energy storage in an oxygen-deficient
LaMnO
2.91
electrode
27
. We have previously hypothesized the
role of lattice oxygen and vacancy exchange in the OER
mechanism on LaNiO
3
refs 16,17. We now revisit this idea to
investigate the role of mobile lattice oxygen in the electrolysis of
water by examining the system La
1 x
Sr
x
CoO
3 d
,0rxr1.
Through substitution of the lower valence Sr
2 þ
ion for La
3 þ
, the
amount of oxygen vacancy defects and the oxidation state of
cobalt can be tuned through the relation
28
:
LaCo
3 þ
O
3
þxSr
2 þ
xLa
3 þ
! La
1 x
Sr
x
Co
3 þ
y
Co
4 þ
1 y
O
3 d
þ
d
2
O
2
ð1Þ
where, d is the oxygen non-stoichiometry parameter, x is
the amount of Sr
2 þ
, and y is the amount of Co
3 þ
in La
1 x
Sr
x
CoO
3 d
, hereafter referred to as LSCO(1 x)x
(that is, LSCO28 for La
0.2
Sr
0.8
CoO
3 d
).
Herein, we describe the intrinsic activities of La
1 x
Sr
x
CoO
3 d
for the OER across the full series from 0rxr1,
including the previously unreported perovskite phase
SrCoO
2.7
with the layered ordering of oxygen vacancies.
The controlled substitution of Sr
2 þ
for La
3 þ
across the full
phase space of the LSCO system while maintaining the perovskite
structure allows us to probe the effects of covalency, vacancy
defects and oxygen exchange on the electrocatalysis of the OER.
The high activities for materials with x40.4 are rationalized
through the high oxygen ion diffusivity and the covalency of the
Co 3d and O 2p bonding in these materials allowing access to a
newly hypothesized lattice oxygen-based mechanism as predicted
through DFT modelling.
Results
Crystallographic characterization. LCO, LSCO and SCO sam-
ples were synthesized using our previously developed reverse-
phase hydrolysis scheme, using a 950 °C calcination temperature
Ligand
hole
E
F
LaCo
3+
O
3
+ Sr
2+
– La
3+
– O
2
+
M 3d (*)
M 3d (*)
O 2p (*)
e
–
La
1–x
Sr
x
Co
1–x
Co
x
O
3
4+3+
La
1–x
Sr
x
Co
y
Co
1–y
O
3–
4+3+
Figure 1 | Relationship between oxygen vacancy concentration and Co–O
bond covalency. As the oxidation state of Co is increased through Sr
2 þ
substitution, the Co 3d/O 2p band overlap is increased (covalency increases)
and the Fermi level decreases into the Co 3d/O 2p p* band, creating ligand
holes. Oxygen is released from the system resulting in oxygen vacancies and
pinning the Fermi level at the top of the Co 3d/O 2p p*band
61
.
ARTICLE NATURE COMMUNICATIONS | DOI: 10.1038/ncomms11053
2 NATURE COMMUNICATIONS | 7:11053 | DOI: 10.1038/ncomms11053 | www.nature.com/naturecommunications
instead of 700 °C to ensure that the correct phase was
synthesized
16,17,27
. Figure 2a shows the powder X-ray diffraction
patterns for the system, demonstrating the successful synthesis of
the perovskite phases across the whole-composition range.
The only minor admixture found in the LCO and LSCO
samples was Co
3
O
4
. The crystal structures of all compositions
have been verified using a combination of powder X-ray
diffraction and transmission electron microscopy. The unit cell
parameters and space groups of the respective materials are given
in Supplementary Table 1. The powder X-ray diffraction and
selected area electron diffraction (SAED) patterns of the x ¼0–0.4
compositions are characteristic of the perovskite R
3c structure
with the a
a
a
tilting distortion of the octahedral framework
(Fig. 2b,e). The monoclinic distortion due to orbital ordering
reported for this compositional range was not detected
being beyond resolution of our powder X-ray diffraction
experiment
29–31
. The LSCO46 composition crystallizes in a
cubic Pm
3m perovskite structure. In the crystal structures of
LSCO28 and SCO ordering of oxygen vacancies becomes obvious
from both SAED patterns and high-angle annular dark-field
scanning transmission electron microcopy (HAADF-STEM)
images (Fig. 2c,d,f,g). Oxygen vacancies reside in the (CoO
2-d
)
anion-deficient perovskite layers alternating with the complete
(CoO
2
) layers that results in a tetragonal a
p
a
p
2a
p
(a
p
indicates the parameter of the perovskite subcell) supercell
in LSCO28. The anion-deficient layers manifest themselves as
faintly darker stripes in the HAADF-STEM images (marked with
arrowheads in Fig. 2f,g), which according to Kim et al.
32
is related
to the structural relaxation in these planes. The anion-deficient
layers form nanoscale-twinned patterns in both the LSCO28
and SCO samples (Fig. 2f,g). In general, the crystallographic
observations on the LCO and LSCO samples are in agreement
with the La
1 x
Sr
x
CoO
3 d
phase diagram
33
. However, in contrast
to the earlier reported Sr
2
Co
2
O
5
brownmillerite or hexagonal
Sr
6
Co
5
O
15
phases
34,35
, the SCO sample demonstrates another
type of oxygen vacancy ordering. The [010]
p
SAED pattern of
SCO (Fig. 2d, top) is strongly reminiscent to that of the
Ln
1 x
Sr
x
CoO
3 d
(Ln ¼Sm-Yb, Y) perovskites with the I4/
mmm 2a
p
2a
p
4a
p
supercell
33,36,37
. A detailed deconvolution
of this SAED pattern into contributions from the twinned
domains is presented in Supplementary Fig. 1. This supercell also
allows complete indexing of the powder X-ray diffraction pattern
of SCO (Supplementary Fig. 2). The layered ordering of the
oxygen vacancies in the LSCO28 and SCO samples was directly
visualized using annular bright-field STEM (ABF-STEM) imaging
(Fig. 3a,b). In both structures the anion-complete (CoO
2
) and
anion-deficient (CoO
2 d
) layers can be clearly distinguished,
alternating along the c-axis of the tetragonal supercells. However,
3020
40 50 60 70 80
90
2θ (°)
*
*
Normalized intensity (A.U.)
x = 1
x = 0.8
x = 0.6
x = 0.4
x = 0.2
x = 0
x in La
1–x
Sr
x
CoO
3–
100
p
[010]
p
001
p
[110]
p
001
p
001
p
001
p
110
p
[010]
p
100
p
[010]
p
100
p
110
p
001
p
[110]
p
-
-
110
p
[110]
p
-
001
p
abcd
efg
Figure 2 | Structural characterization of La
1 x
Sr
x
CoO
3 d
. (a) Powder X-ray diffraction patterns for La
1 x
Sr
x
CoO
3 d
(0rxr1). The reflection from
Co
3
O
4
is marked with an asterisk. (b–d) SAED patterns of LSCO82 (b), LSCO28 (c) and SCO (d). The reflections of the basic perovskite structure are
indexed. The
110½
p
SAED pattern of LSCO82 shows weak G
p
±1/2o1114
p
-type reflections (G
p
—reciprocal lattice vector of the perovskite structure)
characteristic of the a
a
a
octahedral tilting distortion of the perovskite structure. The [010]
p
SAED pattern of LSCO28 demonstrates the
orientationally twinned G
p
±1/2o0014
p
superlattice reflections resulting in the P4/mmm a
p
a
p
2a
p
supercell. The superstructure in the [010]
p
SAED
pattern of SCO can be described with the G
p
±n/4o2014
p
(n—integer) and G
p
±1/2o1104
p
superstructure vectors corresponding to the orientationally
twinned I4/mmm 2a
p
2a
p
4a
p
supercell (see details in Supplementary Fig. 1). Note that the G
p
±1/2o1104
p
superlattice reflections are barely visible
in the
110½
p
SAED patterns of SCO, but the intensity profile (shown as insert in d) along the area marked with the white rectangle demonstrates their
presence undoubtedly. (e–g) [010]
p
HAADF-STEM images of LSCO82 (e), LSCO28 (f) and SCO (g). The image of LSCO82 shows uniform perovskite
structure, whereas the images of LSCO28 and SCO show faint darker stripes spaced by 2a
p
(marked by arrowheads) indicating nanoscale-twinned
arrangement of the alternating (CoO
2
) perovskite layers and (CoO
2 d
) anion-deficient layers. Scale bars are 5 nm.
NATURE COMMUNICATIONS | DOI: 10.1038/ncomms11053 ARTICLE
NATURE COMMUNICATIONS | 7:11053 | DOI: 10.1038/ncomms11053 | www.nature.com/naturecommunications 3
establishing the exact ordering patterns of the oxygen atoms and
vacancies in these (CoO
2 d
) layers requires more detailed
neutron powder diffraction investigation.
In order to understand the effects of Sr
2 þ
substitution on
oxygen vacancy concentrations in La
1 x
Sr
x
CoO
3 d
, iodometric
titrations were performed. It should be noted that processing
conditions affect the oxygen content and oxidation state of cobalt
significantly through equation 1. The results of the iodometric
titrations are presented in Table 1. As can be seen, there is both
an increase in the bulk oxidation state of Co as well as an increase
in the concentration of oxygen vacancies as lower valence Sr
2 þ
is
substituted for La
3 þ
. The high concentration of oxygen vacancies
in SrCoO
2.7
corroborates their pronounced layered ordering.
Microstructural characterization. The overall morphology of the
LSCO series was investigated with bright-field TEM images,
presented in Supplementary Fig. 3. The samples consist of highly
agglomerated and partially sintered nanoparticles with size
ranging from 20–50 nm to few hundred nanometres. The LCO
and SCO materials demonstrate somewhat larger and more
sintered crystallites compared with those of the mixed LSCO
samples. HAADF-STEM and ABF-STEM images of the surface
structure of LCO and SCO are shown in Supplementary Fig. 4,
where the particles remain crystalline at the surface and for SCO
the anion-deficient layers, evident through the nanoscale-twinned
domain columns, extend to the surface. Brunauer–Emmett–Teller
surface areas measured through N
2
adsorption showed similar
surface areas for all samples of 3.1–4.5 m
2
g
1
(Supplementary
Table 2). This surface area is approximately half the surface area
of the materials reported in our previous studies, which results
from the higher calcination temperatures used for the LSCO
series than the previously investigated LaCoO
3
, LaNiO
3
, LaMnO
3
and LaNi
0.75
Fe
0.25
O
3
.
Electrochemical characterization. In order to better understand
the role of oxygen vacancies in La
1 x
Sr
x
CoO
3 d
during elec-
trochemical applications, the intercalation of oxygen in LSCO was
studied using cyclic voltammetry in Ar saturated 1 M KOH
solutions. The insertion and removal of oxygen ions appear as
redox peaks in Fig. 4a. It is apparent that an increase in the
oxygen vacancy concentration as Sr
2 þ
is substituted for La
3 þ
in
LSCO increases the tendency for oxygen intercalation as indicated
through the high current densities measured in the intercalation
region. In addition, it is interesting to note that the position of the
intercalation redox peaks shifts to higher potentials with
increased oxygen vacancies which can be described through the
common pseudocapacitive Nernst Equation:
E ¼ E
0
þ
RT
nF
ln
s
1 s
ð2Þ
where, E represents the measured potential for oxygen
intercalation, E
0
represents the standard potential for oxygen
intercalation, R is the universal gas constant (8.3145 J K
1
mol
1
), T is the temperature during the measurement, F is
0.0 2.2
Distance, nm
ABF intensity
Sr
Co
A
Co
O
O
AO
AO
AO AO
AO
AO
CoO
2
CoO
2–δ
CoO
2–δ
CoO
2
CoO
2
ab
C
0.77 nm
C
0.77 nm
C
C
Figure 3 | ABF-STEM imaging of oxygen vacancy ordering in La
1 x
Sr
x
CoO
3 d
(x ¼0.8, 1.0). (a) [001]
p
ABF-STEM image of LSCO28 showing the cation
and anion sublattices. The contrast is inverted in comparison with the HAADF-STEM images. The assignment of the atomic columns is shown in the
enlargement at the top right corner. Half of the perovskite (CoO
2
) layers appear brighter indicating oxygen deficiency (marked with white arrowheads). The
complete (CoO
2
) layers and anion-deficient (CoO
2 d
) layer alternate (see the ABF intensity profile below, the anion-deficient layers are marked with black
arrowheads) resulting in doubling of the perovskite lattice parameter in the direction perpendicular to the layers. (b) [001]
P
ABF-STEM image of SCO
showing layered anion-vacancy ordering. The (CoO
2 d
) layers are marked with the white arrowheads and demonstrate the contrast clearly distinct from
that of the (CoO
2
) layers. The assignment of the atomic columns is shown in the enlarged part at the bottom left.
Table 1 | Oxygen vacancy concentration, d, and cobalt
oxidation state, y.
xinLa
1 x
Sr
x
CoO
3 d
d y
0 0.01±0.01 3.01±0.01
0.2 0.01±0.01 3.18±0.02
0.4 0.05±0.04 3.30±0.08
0.6 0.09±0.01 3.43±0.01
0.8 0.16±0.01 3.48±0.02
1.0 0.30±0.03 3.40±0.06
Error is based on the s.d. of triplicate measurements.
ARTICLE NATURE COMMUNICATIONS | DOI: 10.1038/ncomms11053
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Faraday’s constant (96,485 C mol
1
), and s is the occupancy
fraction of accessible lattice vacancy sites
38
for the reaction:
La
1 x
Sr
x
CoO
3 d
þ2sOH
ÐLa
1 x
Sr
x
CoO
3 d þs
þsH
2
O þ2se
ð3Þ
This type of Nernst Equation is commonly associated with
pseudocapacitive-type intercalation mechanisms, indicative of
facile oxygen ion diffusion.
The diffusion rates of oxygen ions in LSCO were measured
chronoamperometrically based on a bounded 3D solid-state
diffusion model with a rotating disk electrode (RRDE)
rotating at 1,600 r.p.m. in Ar saturated 1 M KOH
39–41
.
These results are presented in Fig. 4b, and a more detailed
description of the theory behind the model is included as
Supplementary Fig. 5. It was found that SCO, with a vacancy
concentration of d ¼0.30±0.03, had a diffusion rate of
D ¼1.2±0.1 10
12
cm
2
s
1
at room temperature, which is
B40 faster than for LCO, with a complete oxygen sublattice
and a diffusion rate of D ¼3±1 10
14
cm
2
s
1
. As a general
comment, diffusion coefficients in the range of 10
9
to
10
14
cm
2
s
1
have been found as usual values for the short
circuit diffusion of oxygen along high-diffusivity pathways,
including grain boundaries
24
. Although it is unclear whether
the measured diffusion rates are from bulk diffusion or along
grain boundaries, isotope tracer studies have shown that diffusion
rates trend in the order of surface oxygen4oxygen at grain
boundaries4bulk oxygen in perovskite systems, and thus the fast
diffusion rates found in this study represent the lower boundary
on the mobility of oxygen at the surface
42
. Further, the crystallite
size and density of grain boundaries is relatively consistent across
the LSCO series due to the similar synthetic conditions, indicating
that the diffusion rates can at least be compared against each
other. The results indicate that the diffusion rates scale with Sr
concentration because of the correlation with vacancies and Sr
content. The results highlight the benefit of substitution of a
lower valence ion into the A-site as an effective means of
increasing the mobility of oxygen in perovskite oxide electrodes.
The electrolysis of water. The OER activities for LSCO and for a
commercial IrO
2
sample were quantified through cyclic
voltammetry in O
2
saturated 0.1 M KOH at 1,600 r.p.m., as
shown in Fig. 5a. Each material was mixed at a mass loading of
30 wt% perovskite on a mesoporous nitrogen-doped carbon (NC)
or onto Vulcan Carbon XC-72 (VC) for stability measurements.
An evaluation of the carbon loading and total mass loading is
presented in Supplementary Fig. 6, Supplementary Table 5 and
the Supplementary Discussion. There is a shift towards more
active Tafel slopes with increasing Sr content, with LCO
and IrO
2
having similar Tafel slopes of @ V/@ ln i ¼58 mV dec
1
(E2RT/F) which decreases towards SCO with a Tafel slope of
@ V/@ ln i ¼31 mV dec
1
(ERT/F). This shift of Tafel slope for
the OER may be indicative of the facile surface kinetics for oxygen
exchange with increasing vacancy content, whereby OER kinetics
that are limited by high-coverage Langmuir like behaviour where
surface oxygen is not exchanged rapidly (y-1) show Tafel slopes
of 2RT/F. In contrast, those materials showing more rapid surface
oxygen exchange in the intermediate coverage Temkin condition
range (0.2oyo0.8) have slopes of RT/F
9
. The specific activities at
an overpotential of 400 mV, based on perovskite surface area
from BET, are presented in Fig. 5b. It is clear that substitution of
Sr
2 þ
for La
3 þ
in LSCO, and thereby the creation of oxygen
vacancies, is beneficial to the OER, with the fully substituted
SrCoO
2.7
at 28.4 mA cm
2
ox
which is B6 more active than
LaCoO
3.005
(4.3 mA cm
2
ox
), B23 more active than the
commercial IrO
2
sample (1.2 mA cm
2
ox
), and B1.5 more
active than previously reported high-vacancy concentration
cobaltite perovskites (Ba
0.5
Sr
0.5
Co
0.8
Fe
0.2
O
2.6
: B20 mA cm
2
ox
;
Pr
0.5
Ba
0.5
CoO
2.85
: B20 mA cm
2
ox
) (refs 14,43). In addition, due
to the small particle size from the reverse-phase hydrolysis
synthesis, SrCoO
2.7
(3.6 m
2
g
1
) had a mass activity of
1,020±20 mA mg
1
ox
at þ1.63 V versus the reversible
hydrogen electrode (RHE), which is B2 more active than
BSCF with a similar surface area (B500 mA mg
1
ox
, 3.9 m
2
g
1
)
(ref. 14). To verify that the measured current was due only to the
OER, and not to side-reactions or corrosion of the electrode
material, rotating-ring-disk (RRDE) cyclic voltammetry was
performed with a Pt ring poised at þ0.4 V versus RHE,
whereby O
2
generated at the disk from the OER is collected
and reduced at the ring. The results for SrCoO
2.7
/NC and IrO
2
/
NC are shown in Fig. 5c. The collection efficiency for both
SrCoO
2.7
/NC and IrO
2
/NC was 37%, which was equal to the
collection efficiency measured during calibration of the RRDE for
the oxidation of 0.3 mM ferrocene-methanol in 0.1 M KCl.
Therefore, we can confirm that the current is exclusively due to
the generation of oxygen on the SCO or the IrO
2
surface within
the precision of the RRDE measurements.
The stability of SrCoO
2.7
and of the carbon supports under
OER conditions were tested galvanostatically at 10 A g
1
ox
and
1,600 r.p.m., shown in Fig. 5d. As is readily apparent, both the
a
b
1.5
1.0
0.5
0.0
–0.5
–1.0
1.6
1.2
0.8
0.4
0.0
0.0 0.2
0.4
0.6
0.8
1.0
0.4 0.6
0.8 1.0 1.2
1.4
E (V vs RHE)
x in La
1–x
Sr
x
CoO
3–δ
D
o
2– (×10
–12
cm
2
s
–1
) Current density (mA cm
–2
geom.
)
SCO
LSCO28
LSCO46
LCO
LSCO82
LSCO64
Figure 4 | Electrochemical oxygen intercalation into La
1 x
Sr
x
CoO
3 d
.
(a) Cyclic voltammetry at 20 mVs
1
for each member of LSCO in Ar
saturated 1 M KOH. The redox peaks, indicative of the insertion and removal
of oxygen from the crystal, shift to higher potentials with increasing Sr
2 þ
and oxygen vacancy concentrations. (b) Oxygen diffusion rates measured
at 25 °C chronoamperometrically. The diffusion rate increases with Sr
2 þ
and oxygen vacancy concentrations as well. Error bars represent the
standard deviation of triplicate measurements.
NATURE COMMUNICATIONS | DOI: 10.1038/ncomms11053 ARTICLE
NATURE COMMUNICATIONS | 7:11053 | DOI: 10.1038/ncomms11053 | www.nature.com/naturecommunications 5