Trapped O2 and the origin of voltage fade in layered Li-rich cathodes

Abstract

Oxygen redox cathodes, such as Li1.2Ni0.13Co0.13Mn0.54O2, deliver higher energy densities than those based on transition metal redox alone. However, they commonly exhibit voltage fade, a gradually diminishing discharge voltage on extended cycling. Recent research has shown that, on the first charge, oxidation of O²⁻ ions forms O2 molecules trapped in nano-sized voids within the structure, which can be fully reduced to O²⁻ on the subsequent discharge. Here we show that the loss of O-redox capacity on cycling and therefore voltage fade arises from a combination of a reduction in the reversibility of the O²⁻/O2 redox process and O2 loss. The closed voids that trap O2 grow on cycling, rendering more of the trapped O2 electrochemically inactive. The size and density of voids leads to cracking of the particles and open voids at the surfaces, releasing O2. Our findings implicate the thermodynamic driving force to form O2 as the root cause of transition metal migration, void formation and consequently voltage fade in Li-rich cathodes.

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Nature Materials | Volume 23 | June 2024 | 818–825 818
nature materials
Article
https://doi.org/10.1038/s41563-024-01833-z
Trapped O2 and the origin of voltage fade in
layered Li-rich cathodes
John-Joseph Marie 1,2, Robert A. House 1,2 , Gregory J. Rees 1,2,
Alex W. Robertson 1, Max Jenkins1, Jun Chen 1, Stefano Agrestini 3,
Mirian Garcia-Fernandez 3, Ke-Jin Zhou 3 & Peter G. Bruce 1,2,4
Oxygen redox cathodes, such as Li1.2Ni0.13Co0.13Mn0.54O2, deliver higher
energy densities than those based on transition metal redox alone. However,
they commonly exhibit voltage fade, a gradually diminishing discharge
voltage on extended cycling. Recent research has shown that, on the rst
charge, oxidation of O2− ions forms O2 molecules trapped in nano-sized voids
within the structure, which can be fully reduced to O2− on the subsequent
discharge. Here we show that the loss of O-redox capacity on cycling and
therefore voltage fade arises from a combination of a reduction in the
reversibility of the O2−/O2 redox process and O2 loss. The closed voids that
trap O2 grow on cycling, rendering more of the trapped O2 electrochemically
inactive. The size and density of voids leads to cracking of the particles
and open voids at the surfaces, releasing O2. Our ndings implicate the
thermodynamic driving force to form O2 as the root cause of transition
metal migration, void formation and consequently voltage fade in Li-rich
cathodes.
Li-rich cathodes can deliver higher capacities than stoichiometric cath-
odes (up to 300 mAh g
−1
versus ~220 mAh g
−1
for LiNi
0.8
Co
0.1
Mn
0.1
O
2
),
supported by the participation of both transition metal (TM) redox
and O-redox1–5. However, the average voltage of the first discharge
(~3.6 V) gradually diminishes and the load curve develops a step-like
profile as the material is cycled. This ‘voltage fade’ phenomenon leads
to a continuous loss of energy density over cycling, a disadvantage for
commercialization of these materials.
One well-studied aspect of voltage fade is the gradual change in the
redox reactions on the TMs. Over the first cycle, Co
3+/4+
and Ni
2+/3+/4+
are
accepted to be the primary TM redox reactions in Li1.2Ni0.13Co0.13Mn0.54O2,
with Mn remaining predominantly +4 (refs. 6–9). Over cycling, several
studies have reported the increasing participation of low voltage Mn3+/4+
(refs. 10–12) and more recently Co
2+/3+
(ref. 13). Such pronounced TM
reduction is also more generally observed across a range of 3d and 4d
cathodes and has been linked with irreversible out-of-plane TM migra-
tion and voltage fade
14–17
. While it appears that the increasing contribu-
tion of lower-voltage TM redox couples accompanies voltage fade, so
far the underlying cause of voltage fade remains unclear.
A reduction in Li+ diffusivity on cycling, induced by, for example,
structure disordering to form spinel or rocksalt-like surface layers
18–24
,
could lead to a reduction in discharge voltage under the normal condi-
tions of galvanostatic cycling between fixed voltage limits. However,
such a reduction would be a kinetic overpotential. Galvanostatic inter-
mittent titration technique measurements have shown that the voltage
fade is predominantly not due to a larger overpotential, but is rather
a thermodynamic voltage loss25,26. In terms of structural changes,
previous studies have shown that the bulk structural reconfiguration
is even more severe than the surface. The formation of nanopores has
been identified in cycled materials by scanning transmission elec-
tron microscopy (STEM) and three-dimensional tomography13,27, and
He pycnometry measurements have revealed a gradual decrease in
the material density with cycling28. Chapman and co-workers also
identified the growth of nanopores within the bulk on cycling with
small-angle X-ray scattering (SAXS) measurements, but X-ray pair
distribution function measurements could not determine whether they
were filled or empty29. We showed recently, across a range of O-redox
compounds, including Li1.2Ni0.13Co0.13Mn0.54O2, that O2− oxidation results
Received: 15 December 2021
Accepted: 6 February 2024
Published online: 1 March 2024
Check for updates
1Department of Materials, University of Oxford, Oxford, UK. 2The Faraday Institution, Didcot, UK. 3Diamond Light Source, Didcot, UK.
4Department of Chemistry, University of Oxford, Oxford, UK. e-mail: robert.house@materials.ox.ac.uk; peter.bruce@materials.ox.ac.uk
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progression peaks (from 0.125 eV to 2.2 eV) was integrated. This feature
was chosen as the peak intensity arises solely from molecular O2, with
no contribution from the oxide ions. The analysis used to determine
the amount of O2 over cycling is described in more detail in Methods.
For this study, Li1.2Ni0.13Co0.13Mn0.54O2 was charged to specific points
along the 2nd and 100th cycles, representing quarter charge (QC), half
charge (HC), three-quarter charge (3QC), full charge (FC), quarter dis-
charge (QD), half discharge (HD), three-quarter discharge (3QD) and
full discharge (FD) (Fig. 2a,d). These points were defined on a fractional
capacity basis of the total charge/discharge capacities of the 2nd and
100th cycles accordingly. At each state of charge, multiple RIXS scans
were taken across different sample locations to minimize any effect of
sample inhomogeneity, although we note there was little difference
between the spectra (Extended Data Fig. 8). These scans were then
averaged and plotted for both the 2nd and 100th cycles (Fig. 2b,e). The
integrated signal intensity under the O
2
vibrational progression peaks
was plotted as a function of charge and discharge (Fig. 2c,f).
We observed substantial differences in oxygen activity over
cycling. Throughout the charging process on the second cycle, the
amount of O2 is seen to increase continuously over the full voltage
range, which is mirrored by the decrease in O
2
over the subsequent
discharge (Fig. 2c). In contrast, while the 100th charge appears to show
a continuous increase in O
2
, on discharge O
2
reduction to O
2−
occurs
between the FD and HD points on the load curve. The HD point is at ~3 V,
suggesting that the O
2−
/O
2
redox couple is primarily active above 3 V on
discharge, leaving the remainder of the discharge capacity to TM reduc-
tion. This accords with other studies reporting an increased contribu-
tion of lower-voltage Co
2+/3+
and Mn
3+/4+
redox couples after cycling
13,28
.
While the majority of the O2 is observed to be reduced above 3 V,
a small amount may be reduced at lower voltages.
The intensity of the oxygen signal was tracked as a function of cycle
number at FC and full discharge (Fig. 3). The data in Fig. 3 reveal an over-
all loss in the total amount of trapped O2 in the charged cathodes over
cycling, with a 44% loss from the 2nd FC to the 100th FC. Furthermore,
the amount of O
2
remaining at the end of discharge appears to increase
from the 2nd to the 100th cycle (Fig. 3). The accumulation of O2 at the
in the formation of molecular O
2
that is primarily trapped in small
voids in the bulk of these materials and can be reduced back to O2− on
discharge30–32. Following directly the fate of this O2 over cycling is
critical to understanding the growth of nanopores and ultimately to
explaining the origin of voltage fade.
In this Article, we follow directly and measure quantitatively the
trapped O
2
over cycling using high-resolution resonant inelastic X-ray
scattering (RIXS) spectroscopy. We show that the amount of trapped
O
2
formed on charge gradually diminishes over cycling and that the
trapped O
2
that is formed is increasingly difficult to reduce back to
O
2−
on discharge. As the voids containing O
2
grow,
17
O nuclear magnetic
resonance (NMR) data indicate thicker regions of insulating Li–O2− form
on the void surfaces, consistent with the O2 becoming increasingly dif-
ficult to reduce. 129Xe NMR and Brunauer–Emmett–Teller (BET) reveal
increasing amounts of open voids at or near the surfaces over cycling,
suggesting that as the voids grow large and the particle microstructure
weakens, the particles crack releasing O
2
. Together, the accumulation
of electrochemically inactive O
2
in the particles and release of O
2
from
the opening of voids near the surface leads to reduction in the O-redox
capacity. The loss of O2−/O2 redox capacity, which occurs primarily at
potentials greater than 3 V, explains the observed voltage fade.
Voltage fade characteristics in
Li1.2Ni0.13Co0.13Mn0.54O2
Samples of Li1.2Ni0.13Co0.13Mn0.54O2 were prepared by a co-precipitation
synthesis (Methods). The composition was checked by inductively cou-
pled plasma optical emission spectroscopy (Extended Data Table 1) and
energy-dispersive X-ray spectroscopy (Extended Data Figs. 1 and 2),
morphology by scanning electron microscopy (Extended Data
Figs. 1 and 2) and structure by powder X-ray diffraction (PXRD; Fig. 1b
and Extended Data Table 2). The structure of Li
1.2
Ni
0.13
Co
0.13
Mn
0.54
O
2
involves O3-type stacking of the oxide layers, with a honeycomb arrange-
ment of the TM and Li ions in the TM layer (Fig. 1a). The electrochemical
load curves from the 1st to the 100th cycle are plotted in Fig. 1c. After
the voltage plateau seen on the first charge, the load curve develops a
continuous, sloping voltage profile. We have shown that this dramatic
change is accompanied by the irreversible loss of honeycomb ordering
to form vacancy clusters driven by the formation of molecular O2 (ref. 31).
By the end of the first cycle, there is little evidence of the honeycomb
superstructure peaks remaining in PXRD and negligible signal intensity
to monitor over extended cycling. For this reason, the superstructure
peaks were not included in our refinements here. From the 2nd to the
100th cycle, the load curve undergoes further changes, with a higher
proportion of capacity at lower voltage on discharge, that is, voltage
fade (Fig. 1c). This lowering of the average discharge voltage from the
2nd to the 100th cycle follows the same trend as previous reports of
voltage fade on similar compounds
16,33
. A degree of capacity fade is also
observed, similar to previously reported materials prepared in the same
way
29,34–36
. Cycling was performed at a rate of 100 mA g
−1
throughout
the study. To confirm the voltage and capacity fade observed do not
arise from kinetic limitations, cycling data were also collected at a lower
rate of 20 mA g
−1
over 100 cycles (Extended Data Fig. 2). These data
show a very similar degree of voltage and capacity fade, confirming
that these phenomena arise from bulk thermodynamic properties
rather than kinetics.
Redox changes on cycling
We showed previously, using high-resolution RIXS and
17
O NMR, that
molecular O2 is formed during O2− oxidation in Li-rich cathodes31. To
follow the changes in the amount of molecular O2 that is formed on
cycling, we employed quantitative high-resolution RIXS at the O K edge.
In the high resolution RIXS spectra, there are two main features associ-
ated with molecular O2: an energy loss feature at ~8 eV, and a series of
vibrational progression peaks propagating from the elastic RIXS peak
at 0 eV. To track the relative amount of O
2
, the area under the vibrational
a
b
a
b
c
0 50 100 150 200 250 300 350
2.0
2.5
3.0
3.5
4.0
4.5
5.0
100th
60th
20th
5th
2nd
1st
Voltage vs Li+/Li (V)
Capacity (mAh g–1)
c
b
Li/Ni
Co/Mn
Li layer
TM layer
Intensity (×105 counts)
2θ (Cu Kα1,2)
20 30 40 50 60 70 80 90
0
1
2Observed
Calculated
Background
Dierence
Bragg position
Fig. 1 | Structural characterization and electrochemical data for
Li1.2Ni0.13Co0.13Mn0.54O2. a, Li1.2Ni0.13Co0.13Mn0.54O2 with a layered
R
3m
structure,
in-plane ordering of Li/Ni and Co/Mn giving rise to the honeycomb
superstructure ordering. Li atoms are represented in blue, TM in purple and
oxygen in red. b, PXRD data and refinement to the
R
3m
crystal structure. c, Load
curves for Li1.2Ni0.13Co0.13Mn0.54O2, cycled between 2.0 V and 4.8 V at 100 mA g−1 for
100 cycles.
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end of discharge signals that the trapped O2 is becoming increasingly
electrochemically inactive. Taking the difference between the O
2
at
the end of charge and discharge reveals that the amount of O-redox
capacity diminishes from 0.48 e
−
per formula unit on the 2nd cycle
to 0.22 e− per formula unit on the 100th cycle. As a percentage of the
total charge passed O-redox diminishes from 55% on the 2nd cycle to
34% on the 100th cycle, the balance of the capacity being made up by
TM redox. The loss of O-redox capacity on cycling arises in part from
the formation of electrochemically inactive O
2
that is still trapped but
cannot be reduced on each cycle but also from the overall loss of O
2
from the particles, reflected in the 44% loss of O2 at the end of charge
after 100 cycles (Fig. 3). The remaining active O-redox capacity on the
100th cycle (0.22 e
−
per formula unit) aligns with the charge passed
between FC and HD, that is, the charge passed above 3 V. Overall the
loss of O-redox activity, through a combination of electrochemically
inactive O
2
and release of O
2
from the particles, can account for the
b
a
e
d
cf
0 50 100150200250
2
3
4
5
Voltage vs Li+/Li (V)
Capacity (mAh g–1) Capacity (mAh g–1)
QC
HC
3QC
FC
QD
HD
3QD
FD
2nd P
0 50 100150200
2
3
4
5
Voltage vs Li+/Li (V)
QC
HC
3QC
FC
QD
HD
3QD
FD
100th P
0 2.5 5.0 7.5 10.0
Energy loss (eV)
2nd P
2nd P
QC
HC
3QC
FC
QD
HD
3QD
FD
0 2.5 5.0 7.5 10.0
Energy loss (eV)
100th P
100th P
QC
HC
3QC
FC
QD
HD
3QD
FD
QC HC 3QC FC QD HD 3QD FD
0
50
100
150
200
O2 signal intensity (a.u.)
O2 signal intensity (a.u.)
QC HC 3QC FC QD HD 3QD FD
0
50
100
150
200
2nd cycle
2nd cycle
100th cycle
100th cycle
2nd cycle 100th cycle
Fig. 2 | Evolution in bulk O-redox activity over 2nd and 100th cycles. a,d, Load
curves for the 2nd (a) and 100th (d) cycles for Li1.2Mn0.54Co0.13Ni0.13O2, with the
states of charge studied. b,e, RIXS spectra at 531.5 eV collected over the 2nd
(b) and 100th (e) cycles. c,f, Variation in intensity of the molecular O2 signal in
the RIXS spectra over the 2nd (c) and 100th (f) cycles, as determined by principal
component analysis (Methods). Data are presented as mean ± standard deviation
with a sample size of 15. P, pristine.
FC
FD
O2 signal intensity (a.u.)
Cycle number
0 20 40 60 80 100
0
50
100
150
200
Fig. 3 | Evolution in amount of trapped O2 over cycling. Variation in intensity
of the molecular O2 signal from RIXS over cycling in the fully charged (FC) and
fully discharged (FD) states. The amount of O2 formed in the charged materials
decreases with cycling and there is increasing evidence of O2 that is not reduced on
discharge. Data are presented as mean ± standard deviation with a sample size of 15.
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Article https://doi.org/10.1038/s41563-024-01833-z
reduced contribution of charge above 3 V, leading to voltage fade and
much of the capacity loss on cycling Li1.2Ni0.13Co0.13Mn0.54O2.
Void formation on cycling
To examine the origin of the diminishing extent of O
2
formation and
reduction with cycling, annular dark field (ADF)-STEM imaging was
carried out to probe changes in the particle microstructure. The images
shown in Fig. 4a–c and Extended Data Figs. 4 and 5 illustrate substantial
changes within individual grains of Li
1.2
Ni
0.13
Co
0.13
Mn
0.54
O
2
after 100
cycles, consistent with previous reports identifying void formation and
growth using SAXS, STEM and three-dimensional tomography13,27–29.
Comparing the images for the pristine and 2nd discharge with the 100th
discharge clearly shows the development of very extensive voiding
with a high density of voids, as seen by the darker areas. The voids vary
from about 4–12 nm in size and appear to be distributed throughout
the particle. Given the high density of voids it is likely that a number of
these voids are interconnected leading to pores of larger dimensions
than is apparent from STEM.
To investigate the extent to which the voids are closed or open,
129
Xe NMR was used. Xenon-129 (I = 1/2) is an inert gas with a large polar-
izable electron cloud, giving xenon a wide chemical shift range37. When
129
Xe is constrained within a void, it comes into contact with the surface
of the void causing the electron cloud to be distorted and a chemical
shift to be observed38. This shift has a well-known relationship with void
size that has been experimentally determined, described in Extended
Data Fig. 6 (ref. 39). Li
1.2
Ni
0.13
Co
0.13
Mn
0.54
O
2
was first degassed under
dynamic high vacuum for 48 h to evacuate the sample, before being
flushed with xenon gas (Fig. 4d). The xenon spectra for samples of the
pristine, 2nd and 100th discharge material in Fig. 4e all show resonances
centred at ~10 ppm, consistent with bulk xenon gas being paramagneti-
cally broadened due to interactions with the surface of the cathode
material. The signal could arise from large (>40 nm) open pores at the
surface, although no such pores are evident in the ADF-STEM. After
100 cycles, increased spectral density between 20 ppm and 50 ppm
is observed in the
129
Xe NMR spectra. This signal is not present in the
pristine or second cycled material, meaning new open voids with a
minimum size of 17 nm are forming (Fig. 4e). This new signal would
not arise if O
2
had been lost directly from the surface, only Xe atoms
in a partially confined pore open to the surface experience a chemi-
cal shift and so there must be increased surface porosity. These NMR
observations are supported by BET measurements which also show
there is an increase in the number of open pores at the surface, >20 nm
in diameter between the 2nd and 100th cycles (Extended Data Fig. 7).
The somewhat larger pore sizes from NMR and BET are consistent with
some of the pores seen in STEM being interconnected and hence larger
than the STEM images suggest.
The contents of the closed voids
In addition to the
129
Xe NMR,
6
Li and
17
O solid-state magic angle spin-
ning (MAS) NMR were used to investigate changes over cycling. The
6
Li
solid-state MAS NMR of the pristine material (Fig. 4f) shows two regions
of resonances; the resonance that arises from Li in the TM layer is at
20 nm20 nm20 nm
a b c
d fe
Pristine 2nd discharge 100th discharge
120 80 40 0 –40
δiso (129Xe (ppm))
50–20
ppm
2nd
discharge
Pristine
100th
discharge
2nd
discharge
Pristine
100th
discharge
2nd
discharge
Pristine
100th
discharge
Cycled
cathode
particle
Xe
gas
Open pore
Electrode
NMR tube
Evacuate
Xe purge
4,000 2,000 0 –2,000
δcg (6Li (ppm))
8,000 4,000 0 –4,000
δcg (17O (ppm))
g
129Xe NMR 6Li NMR 17O NMR
Fig. 4 | Formation of voids and large diamagnetic Li-rich regions over cycling.
a–c, ADF-STEM images showing single grains of the pristine (a), 2nd discharge
(b) and 100th discharge (c) material showing the formation of voids about
4–12 nm in diameter over extended cycling. d, 129Xe NMR experiments. Samples
were extracted from cells and infiltrated with Xe gas to probe the open porosity.
e, 129Xe NMR of the pristine, 2nd discharge and 100th discharge materials.
The orange region highlighted indicates the presence of open voids of 17 nm
diameter and greater after 100 cycles. δiso, isotropic chemical shift. f,g, 6Li (f) and
17O (g) NMR isolating slow and fast relaxing environments. The sharp peaks at
0 ppm in the 6Li and slow relaxing 17O NMR spectra indicate the formation of large
diamagnetic Li-rich regions on extended cycling.
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~1,500 ppm and is consistent with well-defined honeycomb ordering,
and those between 400 ppm and 900 ppm are from Li in the alkali metal
layers
31,40
. After two cycles, broadening of the signal and substantial
shift in the centre of gravity of the resonance (from ~700 to ~350 ppm) is
observed, consistent with local disordering of the cathode and cluster-
ing of Li into more diamagnetic regions with a reduced number of TM
neighbours, these changes have been discussed previously
31
. After 100
cycles, a notable sharp resonance at 0 ppm is observed that is consist-
ent with Li in an extended diamagnetic environment.
The corresponding 17O MAS NMR spectra (Fig. 4g) of the pristine
material show a broad amorphous line shape with very limited resolu-
tion. This is attributed to a range of oxide-ion environments that are
broadened due to paramagnetic interactions. After 100 cycles at the
end of discharge, the spectral density of the resonance has shifted
to a slightly higher frequency, suggesting TM-rich regions are form-
ing, that is, where the oxide ions are coordinated by several TM ions
(O–TM
3
, O–TM
4
…). There is also evidence of the formation of a peak at
~0 ppm. These O atoms exhibit a similar diamagnetic chemical shift in
the 17O NMR spectrum to lithia, Li2O. As trapped O2 is reduced back to
O2−, the Li+ ions reinserted into the particles surround the O2− resulting
in the nanoscopic regions of Li
2
O (refs. 31,41). This is consistent with
the 100th cycle
6
Li spectrum and can be attributed to the formation
of Li-rich regions in the cathode (O–Li
5
and O–Li
6
). Together, the
17
O
and
6
Li NMR suggests that the materials segregates into regions of
highly paramagnetic TM-rich and diamagnetic Li-rich clusters over
extended cycling.
To probe for molecular O
2
over the 100th cycle using
17
O NMR, fast
relaxing spectra (2 ms), compared with those discussed above (100 ms),
were collected to selectively enhance the oxygens which experience
substantial paramagnetic relaxation enhancement. O2 molecules, pos-
sessing two unpaired electrons, are expected to have a much stronger
paramagnetic relaxation enhancement than oxides. The data (Fig. 5)
show a well-defined chemical environment centred at δ
cg
= 2,770 ppm,
where δ
cg
is the chemical shift centre of gravity, with a manifold of spin-
ning sidebands, consistent with previous measurements of trapped
molecular O
2
(refs. 31,42). A decrease in the intensity of the O
2
reso-
nance can be seen between the 100th charge and discharge as the O
2
is
reduced to O
2−
; however, there is evidence of some residual O
2
present
in the 100th discharged sample. These results are in accord with our
RIXS measurements (Figs. 2 and 3) and indicate that O
2
is only partially
reduced over the 100th cycle. The corresponding slow relaxing spectra
(Fig. 5b) reveal that the 17O diamagnetic peak arising from O2− sur-
rounded predominantly by Li, Li–O
2−
, is completely absent in the 100th
charge sample. This evidence supports the conclusion that the closed
voids seen in ADF-STEM are filled with O
2
and that upon discharge this
oxygen is partially reincorporated into the lattice as O
2−
in ionic Li-rich
regions (Fig. 5c), accompanied by the reinsertion of Li
+
ions into the
void space coordinated by the O2−, as we have described previously31.
O2 loss and residual trapped O2 explain
voltage fade
We showed recently, using high-resolution RIXS and
17
O NMR, that
on the first cycle of the O-redox Li1.2Ni0.13Co0.13Mn0.54O2 material, O2− is
oxidized to O
2
with the loss of honeycomb TM ordering and forma-
tion of small vacancy clusters trapping the O
2
molecules distributed
throughout the particle. The formation of trapped O
2
occurs quickly
on charging as evidenced by the lack of electron-hole states on the O
2−
sublattice. By the end of the first cycle, the trapped O
2
is completely
reduced back to O2−, accounting for the reversible O-redox capacity.
The results presented here show that, on subsequent cycling, the
O-redox mechanism is not static and continues to evolve, although
more gradually. On cycling, there is increasing accumulation of O2 at
the end of discharge (Fig. 3), indicating that not all O
2
formed on charge
is reduced on the subsequent discharge, that is, there is a decrease
Fast relaxation
(D1 = 2 ms)
Slow relaxation
(D1 = 100 ms)
ba c
Discharging
100th charge TM–O2–
TM–O2–
TM–O2–
Li-O2–
Li–O2–
O2
O2
O2
δcg (17O (ppm)) δcg (17O (ppm))
100th discharge
100th charge
100th discharge
12,000 8,000 4,000 0 –4,00012,000 8,000 4,000 0 –4,000
δcg (17O (ppm)) δcg (17O (ppm))
12,000 8,000 4,000 0 –4,00012,000 8,000 4,000 0 –4,000
Fig. 5 | Partial reduction of O2 trapped in voids to form Li-coordinated
O2− on the 100th discharge. a,b, 17O NMR spectra isolating fast (a) and slow
(b) relaxing 17O environments. The sharp peaks in a are assigned to trapped
molecular O2, which decrease in intensity on discharge. There is still evidence
of some residual molecular O2 in the discharged sample, δcg(17O2) = 2,770 ppm.
In b the slow relaxation 17O is dominated by oxide environments coordinated to
paramagnetic TM ions (TM–O2−), δcg = 2,100–2,300 ppm. After discharge, a new
17O environment is formed corresponding to oxide surrounded by Li (that is,
Li–O2−) created by the reduction of O2 in the voids and reinsertion of Li+ into the
voids coordinated by the O2−, centred at δcg = 0 ppm. D1, relaxation delay. c, Large
voids accommodating O2 are partially repopulated by Li+ on discharge. Most O2 is
reduced to O2− but some residual O2 remains.
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Article https://doi.org/10.1038/s41563-024-01833-z
in the reversibility of the O
2−
/O
2
transformation. The decrease in the
amount of O
2
trapped at the top of charge also signals a loss of O
2
from
the particles. Together, the loss of O2 and the decreasing reversibility
of the O
2−
/O
2
transformation lead to a loss of 0.26 e
−
per formula unit in
O-redox capacity, corresponding to a reduction in the percentage of the
capacity due to O-redox from 55% on the 2nd cycle to 34% on the 100th
cycle. As the O
2−
/O
2
couple occurs above 3 V, the loss of O-redox capac-
ity at the expense of a higher proportion of TM capacity leads to the
overall voltage loss on cycling. The decrease in the reversible O-redox
capacity is also commensurate with the capacity fading on cycling,
that is, the capacity loss on cycling is associated with reduction in the
O-redox activity at the higher voltages and a lower average voltage.
Accompanying the loss of O-redox capacity, Li-rich materials
exhibit pronounced changes in the cathode particle microstruc-
ture. There have already been a number of reports of voids forming
on cycling in Li1.2Ni0.13Co0.13Mn0.54O2 using STEM, ptychography and
small-angle scattering
27–29
. Voids also manifest as a reduction in average
particle density, which has been recently observed
28
. Our ADF-STEM
images (Fig. 4a–c and Extended Data Figs. 4 and 5) provide additional
evidence for this, showing voids develop that are about 4–12 nm in
diameter within individual particles after 100 cycles. Previous stud-
ies proposed void formation on the first charge corresponding to a
few vacant cation sites and therefore approximately 1 nm in diameter,
implying the voids grow on cycling. Reduction of O2 trapped in these
larger voids is expected to be more difficult than in the much smaller
voids present on the first and second cycles. This is in accord with the
RIXS observations at the end of discharge showing increasing amounts
of unreduced O
2
over cycling and by
17
O NMR, which also shows evi-
dence of residual O2 at the end of the 100th discharge. Furthermore,
our 17O NMR study of the 100th cycle discharge process reveals that, as
the trapped O2 is reduced to O2−, diamagnetic 17O environments form
that were not present in the charged sample, indicative of O2− in ionic,
2nd cycle
2nd discharge
• Formation of small
diamagnetic Li-rich regions
• Dense particle
structure remains
2nd charge
2nd discharge
2nd charge
• O2 formation in
small vacancy clusters
100th charge
• O2 formation in
larger vacancy clusters
• Particle fracture
• O2 loss from open voids
• Formation of larger
diamagnetic Li-rich regions
•Incomplete reduction of
trapped O2 in larger closed
voids
100th charge
100th discharge
100th discharge
100th cycle
O2 loss from void
coarsening and
opening
a b
c d
O2
reduced
to O2–
O2
trapped
in bulk
Densified
rocksalt/spinel
surface
Vacancy
clusters
Li-rich
region
O2
O2
Closed
void
Open
void
Residual
trapped O2
Some O2
reduced
to O2–
Fig. 6 | Voltage fade mechanism. a,b, Second cycle: reversible O-redox involves
the formation of molecular O2 trapped in small vacancy clusters throughout
the particle. O2 molecules are fully reduced to O2− on discharge forming small
diamagnetic Li-rich regions. c,d, One-hundreth cycle: further TM migration
leads to agglomeration and coarsening of clusters into larger voids driven by the
formation of more O2. The large voids and their high density in the particles lead
to a weakening of the latter, cracking and O2 release. It is also more difficult to
reduce O2 in larger voids.
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Article https://doi.org/10.1038/s41563-024-01833-z
Li-rich environments. This evidence further supports the conclusion
that the closed voids formed on cycling are filled with O2 and that on
discharge some of the O
2
is reduced to O
2−
, which is reincorporated
into the lattice along with the charge compensating reinserted Li+
(Fig. 6). The insulating nature of these Li–O2− regions that will form
on the walls of the void where the electrons to reduce O2 on discharge
are supplied offers an explanation for why it is increasingly difficult to
reduce O2 in larger voids.
The 129Xe NMR and BET data show that the cathode particles
develop increased density of open voids >17 nm in diameter at the
particle surfaces on cycling. There is also a wider body of evidence
showing that Li-rich NMC suffers from particle cracking, which had
commonly been associated with increased lattice strain
43
. Together
this suggests that the increased density of relatively large voids filled
with O
2
upon cycling may result in weakening and hence fracturing
of the particles, releasing O
2
from open or partially open voids at or
near the surfaces of the particles and explaining its loss on extended
cycling release (Fig. 6).
Implications
Suppressing the release of O
2
from particles by protecting the surface
with coatings is known to be an effective strategy to prevent capacity
fade in Li-rich cathodes and it can also suppress voltage fade to an
extent. However, a key implication of our study is that surface coatings
cannot eliminate voltage fade. Efforts must be directed towards bulk
mitigation strategies such as avoiding O
2
formation and the appearance
of voids in favour of stabilised hole states on O (ref. 32).
The oxygen redox process, which proceeds by the formation
and reduction of trapped O
2
molecules, becomes less prevalent on
cycling Li1.2Ni0.13Co0.13Mn0.54O2. On charging, the O2 formed is trapped in
closed voids within the particles. The trapped O
2
becomes increasingly
electrochemically inactive because the growth in size of these closed
voids makes electron tunnelling between the O
2
and the void edges
more difficult. The voids at or near the surface, including any new
fracture surfaces due to particle cracking, are open and can vent O2.
Together these two mechanisms result in the loss of O-redox capacity
on cycling. The gradual loss of O2 participating in the charge compen-
sation reaction over extended cycling offers an explanation for the
voltage fade phenomenon which draws together the observations
of structural reorganisation, void formation, void opening and TM
reduction into a single mechanism. The implication is that voltage
fade mitigation strategies should focus on the bulk and suppressing
the formation of O2.
Online content
Any methods, additional references, Nature Portfolio reporting sum-
maries, source data, extended data, supplementary information,
acknowledgements, peer review information; details of author contri-
butions and competing interests; and statements of data and code avail-
ability are available at https://doi.org/10.1038/s41563-024-01833-z.
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Methods
Co-precipitation synthesis
Ni0.13Mn0.54Co0.13 CO3 precursors were prepared by a coprecipita-
tion route. NiNO
3
·6H
2
O (≥98%, Sigma-Aldrich), MnNO
3
·4H
2
O (≥99%,
Sigma-Aldrich) and CoNO
3
·6H
2
O (≥99%, Sigma-Aldrich) were dissolved
in de-ionized water with a molar ratio of 0.13:0.54:0.13 to prepare a 1.5 M
solution. In addition, a 1.5 M solution of Na2CO3 (≥99.5%, ACS reagent,
Sigma-Aldrich) was prepared. The solutions were added dropwise into
a beaker under continuous stirring, at a constant temperature of 40 °C
and pH 7.6. After full addition of the TM solution, the beaker was left
covered overnight under stirring. The resulting carbonate mix was then
washed with de-ionized water, filtered and dried at 120 °C overnight.
The dried mixed metal carbonate precursor was then mixed with Li
2
CO
3
(≥99%, ACS reagent, Sigma-Aldrich) using a mole ratio of (Li:TM of 1:1.5)
and calcined at 900 °C for 15 h under continuous O2 flow to obtain the
desired compound. A heating and cooling rate of 5 °C min−1 was used
during the synthesis. 17O-labelled samples were prepared in the same
way except the final calcination step was performed under a sealed
atmosphere of O2 gas (CortecNet >70 atom% 17O).
Electrochemical characterization
The electrodes were prepared by combining the active material
(80 wt%), Super P carbon (10 wt%) and polytetrafluoroethylene
binder (10 wt%) using a mortar and pestle. The mixture was then
rolled to a thickness of about 100 μm to form self-supporting films.
Electrodes were assembled into coin cells using Whatman glass fibre
separators and 1 M LiPF6 in ethylene carbonate:dimethyl carbonate
50:50 (battery grade, Sigma-Aldrich) electrolyte, with a Li metal
counter electrode. A typical coin cell has an active mass loading of
~10 mg. Galvanostatic cycle testing was carried out using Maccor
Series 4000. The cells and electrodes were prepared and assembled/
disassembled in the glove box under an inert atmosphere and all
cycling for the characterization studies was performed using the
same conditions. Cells were cycled between 2.0 V and 4.8 V versus
Li+/Li at a rate of 100 mA g−1 without voltage holds, rests or forma
-
tion cycling.
Inductively coupled plasma optical emission spectroscopy
The pristine cathode material was dissolved in aqua regia
(HCl:HNO3/25:75), before diluting the solution for measurement. A
calibration curve was created using standard solutions. Elemental
analysis was carried out by ion-coupled plasma optical emission spec-
troscopy using a PerkinElmer Optima 7300DV ion-coupled plasma
optical emission spectroscope.
PXRD
Diffraction data were collected on a Rigaku 9 kW SmartLab Cu-source
diffractometer equipped with a Hypix 2D detector.
ADF-STEM
ADF-STEM micrographs were measured using an aberration-corrected
JEOL ARM 200F microscope operated at 200 kV. A convergence
semi-angle of 22 mrad was used, with a collection semi-angle of 69.6–
164.8 mrad (ADF). Sets of fast-acquisition multiframe images were
taken and corrected for drift and scan distortions using SmartAlign43.
To avoid the exposure to air, sample transfer to the STEM microscope
was carried out with a vacuum transfer suitcase.
RIXS
High-resolution RIXS data were collected using the I21 beamline at
Diamond Light Source
44
. To produce the data sets for the quantitative
analysis, scans at 531.5 eV were recorded at 15 different sample loca-
tions and averaged together, with little inhomogeneity in the signal
observed (Extended Data Fig. 8). The line scan data from 0.1 eV up to
13.0 eV, excluding the signal from the elastic peak, were z-scored by
dividing each scan by its standard deviation. Then, the area under the
vibrational peak progression (from 0.13 eV to 2.2 eV) was integrated to
measure the relative amount of O2. The areas for each scan were aver-
aged to create a measure of oxygen intensity, with errors coming from
the standard deviation of the mean for each data set.
Solid-state 17O and 6Li MAS NMR spectroscopy
All 6Li and 17O MAS (υR = 37037 Hz) solid-state NMR were completed at
9.45 T (υ06Li = 58.92 MHz, υ017O = 54.25 MHz) using a Bruker Avance III
HD spectrometer and a 1.9 mm double air bearing MAS probe, where
υR is the MAS frequency and υ0 is the Larmor frequency. All 6Li and 17O
spectra are referenced to 1 M
6
LiCl
(aq)
and H
217
O, respectively, at 0 ppm.
All spectra were recorded using a Hahnecho (π/2–τ–π–τ) sequence,
where τ is 1/υR and π/2 is 250 kHz; the resultant free induction decay is
processed as a half echo. The 6Li spectra were achieved with a recycle
delay of 300 ms. These spectra were completed with relaxation times
of 2 ms (fast relaxation) and 100 ms (slow relaxation).
129Xe static NMR
A J-Young NMR tube containing the Li-rich NMC cathode, was degassed
under dynamic high vacuum using a turbo pump for 48 h and then
infilled with natural abundance xenon gas (BOC) at 1 atm of pressure
for 48 h. The
129
Xe NMR (298.1 K, 1 atm) spectrum were completed at
9.45 T (υ0 = 110.69 MHz) using a 5 mm solution-state NMR probe at a
controlled temperature of 298.1 K. A 25 kHz pulse was utilized for all
experiments with a recycle delay of 0.5 s. All shifts are referenced to
natural abundance Xe (gas, 1 atm and 298.1 K) at 0 ppm.
BET
Nitrogen adsorption/desorption analysis was carried using a Micromer-
itics 3Flex Adsorption Analyser. Samples were dried via in situ degas-
sing at 70 °C for 5 h before measurement.
Data availability
All the data generated or analysed during this study are included within
the paper and its Extended Data figures and tables. Source data are
available from the corresponding authors upon reasonable request.
References
44. Zhou, K-J. et al. I21: an advanced high-resolution resonant
inelastic X-ray scattering beamline at Diamond Light Source.
J. Synchrotron Radiat. 29, 563–580 (2022).
45. Terskikh, V. V., Mudrakovskii, I. L. & Mastikhin, V. M. 129Xe nuclear
magnetic resonance studies of the porous structure of silica gels.
J. Chem. Soc. Faraday Trans. 89, 4239–4243 (1993).
Acknowledgements
P.G.B. is indebted to the EPSRC, the Henry Royce Institute for
Advanced Materials (EP/R00661X/1, EP/S019367/1, EP/R010145/1
and EP/L019469/1) and the Faraday Institution (FIRG016) for inancial
support. R.A.H. acknowledges funding from the Royal Academy of
Engineering under the Research Fellowship scheme. We acknowledge
Diamond Light Source for time on I21 under proposal MM25785.
Author contributions
J.-J.M. conducted the synthesis and characterization work. J.-J.M.
prepared the 129Xe-iniltrated samples, R.A.H. prepared the 17O-labelled
samples and G.J.R. performed and itted the MAS NMR. R.A.H. and
J.-J.M. in close collaboration with S.A., M.G.-F. and K.-J.Z. conducted
the RIXS measurements. A.W.R. and J.C. conducted the ADF-STEM
measurements. M.J. performed the BET measurements. J.-J.M., R.A.H.
and P.G.B. wrote the paper with contributions from all authors.
Competing interests
The authors declare no competing interests.
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Additional information
Extended data is available for this paper at
https://doi.org/10.1038/s41563-024-01833-z.
Correspondence and requests for materials should be addressed to
Robert A. House or Peter G. Bruce.
Peer review information Nature Materials thanks William Chueh,
Naoaki Yabuuchi and the other, anonymous, reviewer(s) for their
contribution to the peer review of this work.
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Extended Data Fig. 1 | SEM and EDX images for Li1.2Ni0.13Co0.13Mn0.54O2showing elemental distribution. Spherical particles of 3–4 μm in diameter can be seen,
while the EDX elemental analysis confirms a homogenous distribution of metals within the grains.
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Extended Data Fig. 2 | Electrochemical cycling data at different rates. (a) Cycling data collected at different current rates of 100 and 20 mA/g (C/3 and C/15).
A similar degree of voltage (b) and capacity (c) fade is observed.
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Extended Data Fig. 3 | SEM and EDX for the mixed metal carbonate precursor. Spherical particles of 3–4 μm in diameter can be seen in the SEM and EDX shows a
homogenous distribution of metals within grains and between particles.
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Extended Data Fig. 4 | Additional ADF-STEM images. (a) Pristine, (b) 2nd cycle and (c) 100th cycle. Again, well defined atomic layers can be seen within the grains in the
pristine and 2nd cycles, while the images from the 100th cycle highlight the presence of voids.
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Extended Data Fig. 5 | ADF-STEM images for samples collected after the 100th cycle. Voids are outlined in orange and range in dimensions from 4 to 12 nm. The likely
interconnection of several of these voids is apparent.
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Extended Data Fig. 6 | 129Xe NMR data. 129Xe NMR data at different stages of cycling alongside a plot of void size vs chemical shift45. The minimum void size in the
samples is around ~17 nm, as taken from a maximum chemical shift of 50 ppm.
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Extended Data Fig. 7 | BET data. (a) BET data showing the pore size distributions for the pristine, 2nd cycle and 100th cycle samples. To remove the contribution from
carbon and binder which were present in all samples, the pristine data were subtracted from the 2nd and 100th cycle data to better observe the changes in cathode
particle porosity, (b).
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Extended Data Fig. 8 | RIXS line scans. RIXS line scans collected at 531.5 eV at different sample locations for the charged samples of (a) 2nd and (b) 100th cycles.
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Extended Data Table 1 | ICP-OES data for Li1.2Ni0.13Co0.1 3Mn0.54O2
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Extended Data Table 2 | Rietveld Reinement parameters of powder X-ray diffraction data for Li1.2Ni0.13Co0.13 Mn0.54O2
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1.
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