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ISSN 2050-7488
PAPER
Kun Chang, Zhaorong Chang et al.
Bubble-template-assisted synthesis of hollow fullerene-like
MoS2 nanocages as a lithium ion battery anode material
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JPolueransael odfoMnaotetraiadljsuCshtemmairsgtriynsA
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DOI: 10.1039/C7TA09823E
Journal of Materials Chemistry A
ARTICLE
In situ diffraction studies on reversible oxygen uptake and release
in AFe2O4+ (A= Lu, Yb, Y, In).
R. S. Jayathilake, a B. D. Levitas, a and E. E. Rodrigueza*
Received 00th January 20xx,
Accepted 00th January 20xx
We have studied the reversible uptake and release of oxygen in the layered metal oxide system AB2O4 for A= Lu, Yb, Y, and
In to understand their suitability as oxygen storage materials. We first examined their structures in their pristine state of
AFe2O4 with high-resolution synchrotron X-ray diffraction. To emulate chemical looping conditions, we then monitored their
structures and reactivity under H2 and O2 atmospheres utilizing in situ synchrotron X-ray diffraction and thermogravimetric
analysis (TGA) measurements. The nature of the trivalent A cation affects the oxidation kinetics, thermal cycling stabilities,
and oxygen storage capacities, which varied between 2.20-3.13 wt.% ( 0.7-0.98 O2 mmol/g), corresponding an oxygen non-
stoichiometry () of ~0.5. These layered oxides underwent various phase transitions above 200 °C that included the creation
of a superstructure as oxygen incorporates until the maximally oxidized phase is established above ~400 °C. Bond valence
sum analysis of the Fe-O bonding across the series reveals that the more underbonded the Fe cation, the more facile the
oxygen insertion. During the cycling experiments all samples exhibited stable reversible oxygen insertion at 600 °C with the
exception of A=In, which degraded under H2. The Y analogue displayed the fastest kinetics for oxidation, which may make it
the most suitable for oxygen storage and sensing applications.
DOI: 10.1039/x0xx00000x
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redox kinetics, low operational temperatures, and high oxygen
storage capacities are all metrics for designing new OSMs. 6,8
The materials first studied as OSMs include the binary oxides
1.Introduction
In addition to their unique optical, magnetic and electronic
properties,1 oxides of transition metals in mixed valence states
also display important functions in energy storage and
conversion.2,3 In particular, their ability to reversibly uptake and
release oxygen renders them relevant for key energy
applications such as oxygen storage, energy conversion and
oxygen gas sensing.4,5 Mixed valence metal oxides can act as
oxygen storage materials (OSMs) releasing (or gaining) oxygen
under reducing (or oxidizing) environments, which is often
correlated to their oxygen non-stoichiometries. A challenge in
this field is to design materials whereby oxygen is released and
incorporated topotactically and efficiently as the redox active
metal cycles between two oxidation states.
Chemical looping combustion (CLC) and chemical looping
reforming (CLR) are two applications that utilize OSMs to
thermochemically deliver oxygen from air to fuels. These OSMs
can either completely oxidize a fuel in CLC (to generate CO2+
H2O) or partially oxidize a fuel in CLR (to generate CO+ H2).6
OSMs eliminate the need to purify and isolate oxygen by the
costly process of cryogenic distillation. Furthermore, it enables
easy product separation and prevents product contamination
by NOx gas formation.6,7 Long term stability, fast and reversible
of Mn, Fe, Co, Ni, and Cu due to their natural abundance and
variable oxidation states.9 However, their agglomeration from
sintering under the thermal work conditions have been a major
drawback. As alternatives to binaries, ternary metal oxides such
as perovskites and spinels have been researched as OSMs owing
to their high structural stability and oxygen storage capacities
(OSCs).10–13 Two types of oxygen species in these oxides take
part in the oxidation process: 1) weakly bound lattice oxygen
that completely oxidizes the fuel at lower temperatures 2)
strongly bound lattice oxygen that partially oxidizes the fuel
(chemical reforming) at higher temperatures.10,11 Hence by fine
tuning the type of ternary metal oxide and the reaction
conditions, a well-designed OSM could achieve specific product
selectivity.
The layered oxide LuFe2O4 has drawn recent attention as an
OSM and a conductometric gas sensor owing to its 1:1 mixed
valence states of Fe2+ and Fe3+.14–16 LuFe2O4 does not adopt the
typical AB2O4 spinel or ferrite aristotypes but instead a structure
consisting of bilayers where edge-sharing BO5 trigonal
bipyramids alternate with edge-sharing AO6 octahedra (Fig 1a).
Typically, the A cation is trivalent such as a lanthanide (Ln3+) and
B is a transition metal in a mixed valence state. LuFe2O4 is the
second member of the homologous [AO2]m[BO]n series with m=
1 and n=2,17 and its unit cell has the centrosymmetric space
group R-3m18 (Fig. 1b). However, a lowering of symmetry to P1
or C2/m has been reported due to the structural distortion
arising from charge ordering below 330K that breaks the 3-fold
a.Department of Chemistry and Biochemistry, University of Maryland, College Park,
Maryland 20742-2115, United States.
*efrain@umd.edu
Electronic Supplementary Information (ESI) available: [details of any supplementary
information available should be included here]. See DOI: 10.1039/x0xx00000x
This journal is © The Royal Society of Chemistry 20xx
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To the best of our knowledge, the In and Y analogues have never
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been studied for their reversible oDxyOgI:e1n0.10in39t/aCk7eT/Au0p9t8a2k3eE
capabilities in O2/H2 atmospheres with in situ diffraction. A
comprehensive comparison of A-site substitution effects on
oxidation kinetics and oxygen storage capability of AFe2O4
materials have never been pursued either.
2.1 Materials synthesis
All samples were prepared using solid-state reactions.
99.9% (Alfa Aesar) A2O3 (A=Yb, Lu, In, Y), 99% (Sigma Aldrich)
Fe2O3 and 99.5% (Sigma Aldrich) Fe powder were ground in an
agate mortar and pestle in 1/2: 5/6: 1/3 ratio respectively to
prepare 0.25 g of the target compound, and pressed into pellets
of 13-mm diameter. The pellets were placed in a 2 ml alumina
crucible and sealed inside evacuated 8 mm diameter quartz
ampoules. Before sealing, the ampoules were flushed with N2
gas and evacuated several times to obtain the required oxygen
partial pressure (PO2).34 These compounds are highly sensitive
to the PO2, and A2Fe3O7 often forms as an impurity if the PO2 is
not below approximately 10-7 Torr. The samples were sintered
at 1180 °C (850 °C for In) for 12 hours and then quenched to
room temperature. To make the oxidized phases, AFe2O4+ (ex
situ), the as synthesized samples were oxidized at 500 °C under
air and cooled to room temperature.
Figure 1. a) The crystal structure of AB2O4 and its unit cell in b) R-3m and c) C2/m space
groups. While R-3m is standard, C2/m is used for LuFe2O4 due to its charge order driven
monoclinic distortion below 330 K. The relation between their lattice parameters is
also shown where βM is the monoclinic angle.14
rotational symmetry.19,20 Through electron diffraction studies,
Hervieu et al. reported the monoclinic space group C2/m to
describe the charge ordered phase and its relationship to the
hexagonal unit cell (Fig. 1c).14 However, R-3m is the most
21
commonly reported space group for the isostructural YbFe2O4
YFe2O422 and InFe2O4.23
,
The AB2O4-type oxides have been heavily researched for
their physical properties such as multiferroicity, magnetism,
and linear/non-linear electronic conduction.24–30 However,
oxygen storage properties have not been explored as much in
this system. Hervieu et al. found that LuFe2O4 can be
successfully cycled between its reduced (Fe2.5+) and oxidized (Fe
3+) states at 500 °C under hydrogen and dynamic vacuum.14
Reversible oxygen intercalation was confirmed through X-ray
diffraction and thermogravimetric analysis. Despite the AFe2O4
family having the same structure type, their physical properties
vary depending on the A-site cation. For instance, LuFe2O4
exhibits charge ordering below 330 K14 and ferrimagnetism at
240 K,31 YFe2O4 displays two metal-insulator transitions at 240
K and 220 K with antiferromagnetic spin ordering,32 while
InFe2O4 exhibits both charge and magnetic orderings below
room temperature.33 This study is the first to explore the impact
of elemental substitution in the AB2O4-system on two
properties: 1) kinetics of reduction/oxidation and 2) overall
oxygen storage capacity.
We first report the effects of A-site substitution on the
structure of the four analogues, A= Lu, Yb, Y and In from room
temperature to 700 °C. Next, we relate their reactivity and
suitability as OSMs as determined by cycling experiments
carried out at 600 °C in oxidizing and reducing atmospheres
emulating both CLC and CLR conditions. The rapid data
collection of in situ synchrotron X-ray powder diffraction (sXPD)
allows to monitor the kinetics of the redox reactions by
analyzing unit cell volume and lattice parameter changes with
time. Using thermogravimetric analysis (TGA) experiments we
further investigate their oxygen storage capacities (OSCs) and
the suitability as OSMs in related oxygen-sensitive applications.
2.2 Ex situ high-resolution sXPD
Ex situ experiments with high-resolution synchrotron X-ray
powder diffraction (sXPD) were performed on the 11 BM beam
line at the Advance Photon Source (APS) at Argonne National
Laboratory. An average wavelength of 0.41458 Å was used,
covering a Q-range of approximately ~0.7 Å-1 to 10 Å-1. Powder
patterns were collected at ambient temperature and pressure
for both the pristine compounds (AFe2O4) and their oxidized
phases (AFe2O4+) prepared ex situ. The powder patterns were
analysed by Rietveld refinements for the pristine phases and Le
Bail fits for the oxidized phases using TOPAS 5.35
2.3 In situ sXPD
In situ sXPD experiments were performed in transmission
geometry on the 17 BM beam line at the APS. A 2D PerkinElmer
Si flat panel detector was used with an average wavelength of
0.45336 Å. We collected diffraction patterns every 0.33 seconds
and integrated the diffraction images with GSAS-II36 to yield
patterns with a Q-range of approximately 0.30 Å-1 to 6.6 Å-1. A
constant gas flow was maintained at a rate of 15 mL/min
through a quartz capillary sample holder which was loaded with
the sample, and a residual gas analyzer (RGA) was used to
monitor the gaseous products.
Two types of experiments were carried out for all the A-site
analogues. First, we ramped the temperature of the sample
environment up to 700 °C under 20% O2/He (air), holding at
every 100 °C for 20 minutes. We repeated this temperature
profile in pure helium to differentiate between temperature
and oxidation driven changes. Second, we performed cycling
experiments to emulate a chemical looping reactor, by
2 | J. Mater. Chem. A., 2012, 00, 1-3
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switching the atmosphere between 20 % O2/He and 3.5 % A-site analogues are isostructural with relative changes in
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H2/He. The samples were heated to 600 °C in He and held at 600 intensities, predominantly for the (00D3O),I: 1(00.01063)9/Can7TdA0(9080239E)
°C while cycling between the two atmospheres. Automated reflections, due to the difference in the A-site atomic scattering
sequential refinements37 were carried out using TOPAS 5 to factors. We observed the characteristic splitting of the (015)
study the lattice volume evolution with time.
reflection in R-3m symmetry to the (1-1-2) and (201) reflections
in C2/m symmetry for the Yb and Lu compounds but not the Y
and In analogues as shown in Fig. 2e. However, since our focus
is to compare A-site substitution effects, R-3m (No. 166) was
used as the space group for all four compounds. The lattice
parameters for Lu and Yb in the C2/m space group are included
2.4 Thermogravimetric analysis (TGA)
TGA was conducted using a SDT Q600 equipped with a TA
Discovery MKS104-S0212004 Micron Vision
2
Mass
Spectrometer. The instrument was calibrated prior to the
experiment using alumina standards for each ramp rate. Ramp
rates of 5 °C/min, 10 °C/min, and 25 °C/min were used with a
flow rate of 100 ml/min. Samples were heated to 700 °C in air,
holding at every 100 °C for 10 minutes to allow for maximum
weight gain. An experiment was also done by heating samples
up to 600 °C in air at 10 °C/min and holding the temperature for
3 hours at 600 °C to monitor the stability in the weight gained.
in the supplementary information (Table S1)
.
During the Rietveld analysis, we introduced anisotropic
atomic displacement parameters (ADPs) for the A-site, which
we found to significantly improve the fit (for A=Yb Rwp from
18.17% to 14.04 %). Table 1 shows the refinement results when
anisotropic ADPs were used with the A cation occupying the 3a
special position with 100% occupancy. However, we could
similarly improve the fit when we employed isotropic ADPs for
the A-site while it occupied the 6c site (0, 0, z) at 50% occupancy
(Table S2). The site occupancy of A=In slightly deviated from
unity, which likely occurred due to the high volatility of In. We
therefore sintered InFe2O4 at a lower temperature than the rest.
We performed Le Bail fits for the high-resolution sXPD patterns
of the maximally oxidized phase AFe2O4+ prepared ex situ (Fig.
S5-S8). P-3 was used for A = Yb, Lu and Y38, and C2/m was used
for A = In39 as the space group, and the refined lattice
parameters are shown as an inset
3. Results
3.1 Structures of AFe2O4 and AFe2O4+δ from high-resolution sXPD
We successfully fitted a structural model18 for pristine
AFe2O4 phases (before oxidation) to the high-resolution sXPD
pattern collected at room temperature (Fig. 2). A Pseudo-Voight
function was used for the peak shape profile and the intensities
were corrected for Lorentz-polarization and absorption effects.
The refined structural parameters are listed in Table 1. All four
.
Figure 2. Observed, calculated, and difference curves of the AFe2O4 crystal structures at room temperature (RT) with high-resolution synchrotron X-ray powder diffraction
patterns for A = a) Lu, b) Yb, c) Y, and d) In. A wavelength of 0.41458 Å was used. e) Monoclinic distortion observed in the high resolution sXPD patterns of Lu and Yb analogues
as evidenced by the splitting of the (0 1 5) reflection of R-3m symmetry in to the (1-1-2) and (2 0 1) reflections of C2/m symmetry.
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Table 1. Structural parameters from Rietveld refinement with the high resolution (11-BM) synchrotron X-ray powder diffraction data for the AFe2O4 series at 298 K.
View Article Online
DOI: 10.1039/C7TA09823E
LuFe2O4
R-3m
a = 3.43731 (1)
Å
c = 25.2556 (1)
Å
Rwp= 13.57 %
Atom
Site
3a
x
y
z
Occ.
Beq /Uii (
Å2)
Lu
0
0
0
0
0
0
1
1
1
1
U11 = U22 = 0.00006 (9)
0.357 (11)
U33 = 0.03407 (23) U12 = 0.00003 (9)
Fe
6c
0
0.21530 (2)
0.12902 (11)
0.29206 (11)
O1
O2
6c
0
0.837 (42)
6c
0
0.837 (42)
YbFe2O4
3a
R-3m
a = 3.45879 (1)
Å
Å
Å
c = 25.1205 (1)
Å
Å
Å
Rwp= 14.39 %
Yb
Fe
0
0
0
0
0
0
1
U11 = U22 = 0.00013 (10) U33 = 0.03791 (27) U12 = 0.00007 (10)
6c
0
0.21497 (3)
0.29179 (11)
0.12837 (11)
1
1
1
0.537 (13)
1.102 (44)
1.102 (44)
Rwp = 12.03 %
O1
O2
6c
0
6c
0
YFe2O4
3a
R-3m
a = 3.51188 (1)
c =24.84257 (7)
Y
0
0
0
0
0
0
1
1
1
1
U11 = U22 = 0.00230 (8)
1.059 (7)
U33 = 0.06285 (26) U12 = 0.00115 (8)
Fe
O1
O2
6c
0
0.21454 (1)
0.29193 (6)
0.12828 (5)
6c
0
1.557 (20)
6c
0
1.557 (20)
InFe2O4
3a
R-3m
a = 3.33899 (1)
c = 26.07201 (1)
Rwp = 13.31 %
In
0
0
0
0
0
0
0
0
0.00218(6)
0.21653 (1)
0.12861 (6)
0.29290 (6)
0.970 (1)
U11 = U22 = 0.00350 (8)
0.516(9)
U33 = 0.00716 (13) U12 = 0.00175 (8)
Fe
O1
O2
6c
1
1
1
6c
0.758 (20)
6c
0.758 (20)
which we label R1, were retained. Upon further heating, extra
satellite reflections appeared around 230 °C for all compounds
except for In, which indicated the formation of the supercell
structure (R2 phase). These supercell reflections gradually
disappear around 350 °C for Y and 480 °C for Lu and Yb. In the
250 °C - 450 °C region, the diffraction peaks are not well defined,
3.2 Structural evolution with temperature: In situ sXPD
With in situ sXPD, we explored the thermal stability, optimal
operating temperatures and associated structural phase
transitions of the title compounds. Fig. 3 shows the contour
plots of the in situ sXPD patterns of AFe2O4 as they are heated
under air up to 700 °C. Below 200 °C we did not observe any
new reflections, and the room temperature phases of AFe2O4,
Figure 3. Contour plots of the sXPD data as AFe2O4 is heated in air (20% O2 in He) from 100 °C to 700 °C. a) The Lu, b) Yb, c) Y and d) In analogues. The yellow labels and blue dash
lines indicate phase co-existence. The white arrows and asterisks mark some of the reflections: (0 0 6) belongs to the R1 phase (AFe2O4) in R-3m symmetry, (0 2 4) is one of the
supercell reflections in the R2 phase that develops above 230 °C, (0 0 2) belongs to the R3 oxidized phase (AFe2O4+). The asterisks mark the satellite peaks which belong to the R2
supercell, and are observed for the Lu, Yb and Y analogues. e) The supercell satellite reflections in the isolated powder patterns of the R2 super cell structure at the end of the 20
min hold at 300 °C for Lu, Yb and Y analogues.
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parameter evolution under the two atmospheresV(idewisAcruticsleseOdnlinine
DOI: 10.1039/C7TA09823E
section 4.2).
3.3 Cycling experiments with in situ sXPD
The optimal operational temperature for OSMs in chemical
looping reactions is a trade-off between fast redox kinetics of
oxygen release/uptake and prolonged structural stability. To
achieve fast reaction rates while preserving the structural
stability, we chose 600 °C as the cycling temperature for the in
situ sXPD experiment. To this end, cycling experiments were
performed under oxidizing (20% O2 /He) and reducing (3.5 %
H2/He) atmospheres. Fig. 4 shows the structural evolutions
observed in this experiment as a contour plot of some major
reflections. The cycling stability and reactivity was almost
identical for the Lu and Yb analogues (Fig. 4a, 4b). The In
analogue, however, started to degrade once it was exposed to
the H2 (Fig. 4d).
Figure 4. Contour plots of synchrotron x-ray diffraction patterns of the AFe2O4 series
during cycling between oxidizing (20% O2 in He) and reducing (3.5% H2 in He)
atmospheres at 600 °C. a) LuFe2O4 b) YbFe2O4 c) YFe2O4 and d) InFe2O4. Three
consecutive cycles were carried out for Lu and Yb and two cycles for Y and In. The Lu, Yb
and Y analogues show stable cyclability. The degradation of InFe2O4 under H2 is clearly
evidenced by the diminishing intensities of the Bragg reflections.
3.4 Thermogravimetric analysis (TGA)
To correlate oxidation driven weight changes to
crystallographic changes, we performed TGA mimicking the in
situ sXPD experiment (Fig. 3). Fig. 5 displays the TGA curves
where samples were heated up to 700 °C, equilibrating every
100 °C for 10 minutes. The R2 supercell structure that formed
above 200 °C in the in situ sXPD corresponds to a weight gain of
0.80-1.0 % (for Lu, Yb, Y). The total weight gain, once the R3
phase is formed, varies from 2.2 % to 2.6 % among the four
oxides. The derivative of the weight change is also plotted in Fig.
5 (black curve), and the highest weight gain due to oxidation
occurs between the 300 °C - 400 °C region for all samples. This
rapid weight gain, however, is most pronounced for the Y
analogue, indicating that this analogue displays the fastest
kinetics for oxide insertion. The derivative curve also
and the patterns display a large amount of phase co-existence,
most apparently for Y. Upon further heating, a major structural
phase transition occurs and the maximally oxidized phase (same
as the ex situ structure in Fig. S5-S8) is reached, which we label
the R3 phase. This transition is complete by 430 °C for Y, 500 °C
for Lu and Yb, and 600 ° C for In (Fig. 3). Evidently, the A-site
substitution has an impact on the oxidation temperatures,
controlling the window of stability for the high-temperature R3
phase. All four compounds were stable until 700 °C. The in situ
diffraction experiments performed in air were repeated in He to
isolate temperature driven changes from those driven by
oxidation. We observed different trends in the lattice
Figure 5. Thermogravimetric analysis (TGA) depicting the weight change % as a function of time and temperature for a) LuFe2O4 b) YbFe2O4 c) YFe2O4 and d) InFe2O4.
The corresponding synchrotron powder diffraction patterns are shown on the left of each TGA
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DOI: 10.1039/C7TA09823E
Figure 6. a). The linear dependence of a- and c-lattice parameters on the radii of the
A cations. The size of the spheres is to scale. b) Displacement of Fe from the center
of the trigonal bipyramid towards the AO6 octahedral layer. As a result, there are
three distinctive Fe—O bond lengths which are denoted as Q, R and S.
demonstrates that the In analogue is the most sluggish of the
series, which corroborates our in situ sXPD data. Calculations of
OSCs based on all TGA experiments are presented in section 4.4.
4. Discussion
Figure 7. Isolated in situ sXPD patterns of LuFe2O4, collected while ramping under air
from RT to 700 °C shown with the refined unit cell parameters. a) The structure of
LuFe2O4 at RT b) The super cell structure at 300 °C. The satellite reflections marked
with red arrows are due to the O2 driven modulations. c) The maximally oxidized
structure of LuFe2O4+ at 600 °C
4.1 A-site substitution effects on structure, bond valence of Fe and
kinetics of oxidation
In the AFe2O4 series, the transition metal is the redox active
centre while the A-site cation does not directly participate in the
reaction. While this is the case in many OSM’s consisting of
ternary metal oxides, the A-site cation has shown to play key
roles in determining the stability of the compound,
concentration of oxygen vacancies, and the oxidation state of
the transition metal.28 Shen et al. compared A-site substitution
effects of several AFeO3 perovskites and found that the oxygen
desorption capability varies on the A cation’s radius.40 In this
study Ba and Sr were used as the A-site cations, and the Ba
analogues showed higher oxygen desorption amounts due to
the larger lattice volumes. In a similar study of three hexagonal
AMnO3 perovskites, Abughayada et al. demonstrated that the
oxygen storage capacity is highly sensitive to the A cation’s
radius,41 where the adsorption/desorption of the oxygen
content increased with larger A cations. Hence, we anticipated
comparable results for the AFe2O4 series.
Based on our Rietveld refinements, both structural and cell
parameters are affected by the nature of the A-site in the title
compounds. The a-lattice parameters follow the order Y3+>
Yb3+> Lu3+> In3+, while the c-lattice parameters behave
conversely. The ionic radii of the A-site cation clearly determine
these trends (Fig. 6a). Rietveld refinements of the R1 phase also
revealedsignificant A-site disorder from the special position 3a.
This characteristic disorder has been reported previously for
LuFe2O4.18,13 Infrared spectroscopy has also revealed a splitting
of the Lu—O vibrational modes,42 which arises from the
interactions between the polar FeO5 bilayers with the LuO6
octahedral layers. The polarity in FeO5 bilayers is attributed to
charge ordering of the Fe3+ (electron-deficient) and Fe2+
(electron-rich) cations. In our study, we observed this
anisotropic disorder along the z axis for the A cations in all four
compounds.
same supercell reflections by X-ray and electron diffraction,
indexed them in the rhombohedral R-3m space group,14 and
attributed their appearance to oxygen insertion in a manner
commensurate to the metal sublattice. The refined powder
pattern of the R2 phase for LuFe2O4, isolated at 300 °C, is shown
in Fig. 7b along with the refined unit cell parameters.
Interestingly, we observed the formation of the R2 phase for the
Lu, Yb and Y analogues but not for the case of A = In.
To extend our understanding of the A-site substitution
effects on the structure of the title compounds, we carried out
bond valence (BV) calculations using bond lengths from the X-
ray diffraction results. BVs were calculated using the formula of
Brown and Altermatt,43 and the empirical values R0 and B were
adopted from Gagné et al. (Used R0 and B values are listed in
Table S3).44 Bond lengths and BVs of the Fe—O bonds are
tabulated alongside the A-site ionic radii in Table 2. In the FeO5
trigonal bipyramidal polyhedron, the Fe atoms are off-centre
and shifted towards the AO6 octahedral layers to create two
unequal Fe—O axial bond lengths (Fig. 6b). While the axial Fe—
O bonds showed no significant trend with the choice of A-site
cation, the equatorial Fe—O bond lengths showed a direct
correlation to the A cationic radii. This is also reflected in the
overall bond valence sum (BVS) of Fe—O bonds. The BVS of Fe-
O in the Y and In analogues depict the highest and lowest
deviation from the average valence state of 2.5 in Fe. Hence, in
the Y analogues, iron is the most under-bonded, and in the case
of In, it is the most over-bonded before the oxidation step.
The mixed valence state (+2.5) of the R1 phase has been
observed by Fe-2p, O-1s X-ray absorption spectroscopy (XAS)
and Mössbauer spectroscopy,31,45 and other studies have
confirmed that when forming the R3 phase, the mixed valence
state goes fully to Fe3+.46,38 One noticeable difference across our
series is the temperature at which oxidation to the R3 phase is
Upon heating in air, we observed the supercell structure of
the R2 phase around 250 °C (Fig.3a-c). Hervieu et al. found the
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Figure 9. Comparison of unit cell volume change of the Lu and Yb analogues for 3
consecutive cycles. Refined unit cell volumes from sequential refinements are plotted
with the error bar. Both compounds exhibit
a similar stability and reversibility
throughout the three cycles. Error bars are smaller than the markers.
Figure 8. The evolution of lattice parameters as a function of temperature in He and Air
obtained from sequential refinements for Lu and Yb. While thermal expansion is
observed under He, an anisotropic lattice parameter change is observed under air due
to oxidation driven effects. The error bars are shown and are smaller than the marker
attributed to the decreasing ionic radii of iron upon oxidation.
Approximately 35-44% of Fe-sites are already oxidized at 300 °C
(in air) in the Lu and Yb compounds based on the ionic radii
change (~0.1 Å) and c-lattice parameter change (0.035-0.044 Å).
completed. As mentioned in section 3.2 (Fig. 3), Y is the first to
fully transform to the R3 phase (~ 430 °C) while In is the last
(~600 °C). This difference in reactivity towards oxidation clearly
reflects the difference in the deviation in the BVS of Fe from the
mixed valence (+2.5) state. In the Y analogue, Fe is found in the
most under-bonded environment, and hence oxidizes the
fastest. The reverse is found for the In analogue, where the
sluggish kinetics of oxygen insertion is likely due to Fe being
over-bonded.
4.3 Thermochemical cycling capability and stability
The only study14 to test the oxygen storage capacity and cycling
stability of LuFe2O4 utilized 500 °C as the cycling temperature
where the reaction kinetics were slow. Hence, we chose 600 °C
as the cycling temperature, which did in fact increase the
reaction rate, noticeably for the reduction step. Fig. 9 depicts
the refined cell volume change with time for the Lu and Yb
analogues under air and H2. The rate of cell volume change in
air is relatively faster than that in H2 illustrating the faster
reaction kinetics for oxidation. This is a clear indication of the
difference in the thermodynamic stability of the two phases.
We also found the rate of cell volume change for both
compounds to be nearly the same. This is promising since Yb is
less expensive than Lu, and the Yb analogue could therefore be
proposed as a cost-effective alternative to Lu. While YFe2O4
appeared stable towards cycling (Fig. 4c), it was rather difficult
to calculate the accurate cell volume change due to highly
disordered oxidized phases.
Under air, we observed that these layered oxides were
quite stable but started to decompose above 700 °C. This
observation agrees with the report by Hervieu et al.14 However,
the stability under H2 at 600 °C was not the same for all the
samples. The In analogue started to degrade once it was
exposed to the reducing atmosphere. Although the A-site cation
typically does not take part in the redox reaction in many OSMs,
In3+ reacts with H2 to destabilize the structure and therefore
leads to phase degradation (Fig. S11).
4.2 Lattice expansion from oxygen insertion vs thermal expansion
Upon heating the samples in He, we did not observe any of
the above discussed phase transitions, confirming that they are
completely driven by oxygen insertion. To further isolate the
thermal expansion effects from oxidation ones, the evolution of
refined unit cell parameters with time was analyzed in air versus
He. Fig. 8 depicts the evolution of cell parameters for Lu and Yb
in both atmospheres. In air, anisotropic cell volume changes
were observed as opposed to the isotropic thermal expansion
under He. The a-lattice parameter gradually increased as a
function of temperature while the c-lattice parameter
decreased. A similar temperature dependence for lattice
parameters was reported previously for the Yb and Tm
analogues in the temperature range of 50-400 K.47 In our
samples, the overall cell volume change was almost equal in
both environments as the a-parameter expansion rate in air was
higher than that in He, compensating for the decrease in the c-
parameter in air (Fig. S9). The increase in the a-parameter can
be attributed to the bilayer shift mechanism in the ab-plane due
¯
to O2 insertion, which has been evidenced from electron
diffraction experiments.14 The decrease in c-parameter can be
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Table 2. Comparison of bond lengths, unit cell volume and A3+ cationic radii of the A-site analogues at RT. Q, R and S are labels for the Fe-O bond lengths as shown in Figure 6b.
A-site cation
In3+
Fe-O
R
Bond valence
sum Fe
Shannon radii48 of A
site cation (Å)
Q
S
BL (Å)
1.9521 (2)
0.5515 (3)
2.0079 (5)
0.4803 (6)
2.0164 (5)
0.4703 (6)
2.0392 (2)
0.4445 (2)
2.294 (2)
1.991 (2)
0.501 (2)
1.936 (4)
0.574 (6)
1.919 (4)
0.599 (6)
1.985 (2)
0.508 (3)
1.06
1.117
1.125
1.159
BV (v.u)
BL (Å)
0.237 (1)
2.155 (4)
0.334 (3)
2.155 (4)
0.333 (3)
2.153 (2)
0.335 (2)
2.392 (3)
2.349 (7)
2.343 (7)
2.177 (3)
Lu3+
BV (v.u)
BL (Å)
Yb3+
BV (v.u)
BL (Å)
Y3+
BV (v.u)
BV = bond valence, BL = bond length, v.u = valence unit
ability to understand how the non-redox active metal center
can affect important properties such as oxygen insertion
kinetics and cycling stability. The A-site cationic radii
demonstrate a clear correlation to the lattice parameter change
of the pristine structures and to the iron valence state as
evidenced by the BVS analysis. The deviation of the BVS of Fe
from the ideal 2.5 state is directly proportional to the A-cationic
radii, while the lattice parameters show an anisotropic
correlation. When heated under air all four analogues undergo
a phase transition above ~400 °C from R-3m (R1 phase) to P-3
or C2/m symmetry (R3-phase). For Lu, Yb and Y, the formation
of a supercell structure is observed around 250 °C as evidenced
by satellite reflections in sXPD data. Sequential refinements of
in situ diffraction data indicate an anisotropic cell volume
change in air contrary to the isotropic thermal expansion under
He. As the deviation of the BVS of Fe increases, the temperature
4.4 Comparison of oxygen storage capacity
The maximum oxygen non-stoichiometry (
reported in this series is ~0.5 for LuFe2O4+

) that has been
14
.

Fig. S12 a-c shows
the change of with temperature, calculated with TGA
measurements performed with three different ramp rates. The
maximum obtained is close to ~0.5, which is reached by 400
°C for Y and by 600 °C for the rest of the oxides. Furthermore,
across all ramp rates, the maximum OSCs vary between 2.20-
3.13 wt.% (0.7-0.98 O2 mmol/g) at 600 °C, which are listed in
Table 3. Taylor et al. has reported the OSC of the La1-xSrxFeO3
perovskite system together with OSCs of other metal oxides
from literature.49 In these oxide systems the OSCs range from
1.7-3.7 wt. %. Therefore, AFe2O4 compounds in this study have
comparable OSCs to these perovskites, and structural
modifications could be explored to further increase their OSCs.
Figure S12d shows that once the maximum oxygen gain is
reached at 600 °C, the OSCs remain stable over a time span of 3
hours. While oxidation kinetics do correlate with A-site ionic
radii, the final OSCs do not. However, variations between OSCs
could be complicated by other parameters such as particle size
and surface area. While beyond the scope of this study, these
parameters should be analyzed to better understand how to
improve OSCs in the AFe2O4 system.
at which the phase transition from R1R3 completes in air,
decreases due to the more under-bonded Fe.
These layered oxides exhibit reversible oxygen insertion at
600 °C in alternating 20% O2/He and 3.5% H2/He atmospheres
transitioning between the two phases (R1 and R3). Increasing
cycling temperature from 500 °C to 600 °C increased the rate of
reaction, and the rate of oxidation was faster than that of
reduction. While Lu, Yb and Y showed stable cycling, the In
analogue showed signs of degradation under H2. OSCs
calculated based on thermogravimetric analysis with different
ramp rates vary in the range of 2.20-3.13 wt.% at 600 °C, but no
clear trend with A-cation size could be established for this
metric. However, the oxidation kinetics showed a direct
correlation to the ionic radii across all ramp rates. Based on
these gathered metrics, the Yb and Lu analogues have
comparable reaction rates, cycling stability and OSCs, making Yb
a cost-effective alternative for Lu. Considering the fast oxidation
kinetics of the Y analogue at low temperatures along with its
ability to cycle, it could be a suitable candidate for applications
such as gas sensing in addition to chemical looping.
Table 3. OSCs at 600 °C calculated from TGA experiments shown in Fig. S12.
Ramp rate
OSC
5 °C/min
10 °C/min
25 °C/min
wt.%
O2
wt.%
O2
wt.%
O2
mmol/g
0.70
mmol/g
0.72
mmol/g
0.69
LuFe2O4
YbFe2O4
YFe2O4
2.29
2.41
2.99
2.64
2.20
2.36
3.07
2.60
2.25
2.21
3.13
2.38
0.75
0.93
0.82
0.73
0.96
0.81
0.69
0.98
0.74
InFe2O4
Conclusions
Acknowledgements
We have studied four A-site analogues of AFe2O4 (A=Lu, Yb,
Y, In) as OSMs. The Y and In compounds have never been
explored as OSMs before, and we have found that Y and In
express the fastest and the slowest kinetics, respectively, for
oxygen uptake across all ramp rates. This study affords us the
We acknowledge the Department of Commerce/NIST award
70NANB12H238 for support. Use of the Advanced Photon
Source at Argonne National Laboratory was supported by the
U.S. Department of Energy, Office of Science, Office of Basic
Energy Sciences, under Contract No. DE-AC02-06CH11357. We
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thank W. Xu and A. Yakovenko at 17 BM, ANL, D. Taylor and X.
Wang at university of Maryland for their help with sXPD data
collection, and P. Zavalij at university of Maryland for his help
on X-ray data analysis.
Izumi and H. Asano, Phys. B Condens. Matter,V1ie9w9A2rt,ic1le8O0n–line
181, 603–605.
DOI: 10.1039/C7TA09823E
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Kakurai, Y. Murakami, K. Yoshii, S. Mori, Y. Horibe and H.
Kitô, Nature, 2005, 436, 1136–1138.
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