# Competitive Oxygen Reduction Pathways to Superoxide and Peroxide during Sodium-Oxygen Battery Discharge

Zarko Jovanov, Lukas Lutz, Juan Lozano, Conrad Holc, Xiangwen Gao, Alexis Grimaud, Jean-marie Tarascon, Yuhui Chen, Lee Johnson, Peter Bruce

► To cite this version:

Zarko Jovanov, Lukas Lutz, Juan Lozano, Conrad Holc, Xiangwen Gao, et al. Competitive Oxygen Reduction Pathways to Superoxide and Peroxide during Sodium-Oxygen Battery Discharge. Batteries & Supercaps, 2022, 5 (9), pp.e202200055-1/e202200055-6. 10.1002/batt.202200055. hal-04104521

HAL Id: hal-04104521

https://hal.science/hal-04104521v1

Submitted on 16 Apr 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
<figure><img src="image_2.png" /></figure>

# Competitive Oxygen Reduction Pathways to Superoxide and Peroxide during Sodium-Oxygen Battery Discharge

Zarko P. Jovanov,[a, f] Lukas Lutz,[a, b] Juan G. Lozano,[a] Conrad Holc,[a, c] Xiangwen Gao,[a, d] Alexis Grimaud,[b] Jean-Marie Tarascon,[b] Yuhui Chen,[*e] Lee R. Johnson,[*c] and Peter G. Bruce*[a]

The sodium-air battery offers a sustainable, high-energy alternative to lithium-ion batteries. Discharge in the cell containing glyme-based electrolytes can lead to formation of large cubic NaO₂ particles via a solution-precipitation mechanism. While promising, high rates result in sudden death. The exact nature of the discharge product has been a matter of contention, and Na₂O₂ has never been directly detected in a dry glyme Na–O₂ cell. If Na₂O₂ were to form during discharge in the Na–O₂ cell it

would have a detrimental impact on cell performance. Here we show that Na₂O₂ forms during discharge at high overpotential in the glyme-based Na–O₂ batteries. Na₂O₂ formation is confirmed by spectroscopic and electrochemical analysis and electron microscopy demonstrates that it results in thin insulating films at the electrode surface. The formation of these thin films results in rapid cell death during discharge, introducing an inherent chemical limitation to the Na–O₂ battery.

## Introduction

Metal-air (O₂) batteries are promising high specific energy alternatives to lithium-ion batteries.[1–11] In the aprotic sodium-oxygen (Na–O₂) battery, sodium metal is oxidized on discharge and oxygen is reduced at the positive electrode.[12–14] In Na–O₂ batteries containing glyme-based electrolytes, NaO₂ is found to

be the major discharge product and it is commonly reported to form large cubic crystals.[12,15–17] The formation of NaO₂ crystals occurs via the dissolution of NaO₂ at the electrode surface and its subsequent precipitation away from the electrode. The solubility is controlled in part by the solvent donor number, as is the case in the Li–O₂ system,[18–20] as well as the Na⁺ activity.[21] Several publications have shown promising results with this system.[4,13,22–24] Advantages of the Na–O₂ battery compared to the Li–O₂ battery stem from the formation of a superoxide rather than a peroxide, as this simple electrochemical reaction and soluble product results in relatively low polarisation and higher rates during cycling.[2,7,25,26] However, parasitic chemical reactions can occur, impacting on the cyclability of cells. This arises from reactions of both the Na electrode and the NaO₂ discharge product, but careful selection of electrolytes and cycling parameters can mitigate these issues.[27,28]

At high rates, the capacity of typical glyme-based Na–O₂ cells is lower and sudden death of the cell is common.[2,29,30] This has been linked to the formation of NaO₂ thin films at the electrode surface.[20,29] We have shown that this is in part due to the limited solubility and rate of (de)solvation of NaO₂,[31] which must be rapid enough to sustain the dissolution-precipitation mechanism by which NaO₂ cubes grow.[20,29] However, computational studies have shown that electrochemical formation of Na₂O₂ may also be possible at higher overpotentials.[15] It is well known that in the Li–O₂ cell, O₂ is reduced to Li₂O₂, which is more stable than LiO₂.[19,32–35] Unlike the Li–O₂ system, the theoretical equilibrium potentials for the formation of NaO₂ and Na₂O₂ from oxygen are only 60 mV apart (2.27 V for NaO₂ and 2.33 V for Na₂O₂ vs. Na/Na⁺). If Na₂O₂ were to form during discharge in the Na–O₂ cell it would have a detrimental impact on cell performance and remove many of its advantages over the Li–O₂ battery. The discharge product(s) and performance of Na–O₂ cells are influenced by several factors, including electro-

[a] Dr. Z. P. Jovanov, Dr. L. Lutz, Dr. J. G. Lozano, Dr. C. Holc, Dr. X. Gao, Prof. P. G. Bruce
Departments of Materials and Chemistry, University of Oxford
Oxford, OX2 6HT, U.K.
E-mail: peter.bruce@materials.ox.ac.uk

[b] Dr. L. Lutz, Dr. A. Grimaud, Prof. J.-M. Tarascon
Collège de France
11 place Marcelin Berthelot, 75231 Paris Cedex, France

[c] Dr. C. Holc, Dr. L. R. Johnson
Nottingham Applied Materials and Interfaces (NAMI),
Group GSK Carbon Neutral Laboratories for Sustainable Chemistry,
University of Nottingham
Nottingham, NG7 2TU, U.K.
E-mail: Lee.johnson@nottingham.ac.uk

[d] Dr. X. Gao
Materials Science and Engineering Program and Texas Materials Institute,
The University of Texas at Austin
Austin, TX 78712, U.S.A.

[e] Prof. Y. Chen
State Key Laboratory of Materials-oriented Chemical Engineering,
College of Chemical Engineering and
School of Energy Science and Engineering,
Nanjing Tech University Nanjing
Jiangsu, 211816, China
E-mail: cheny@njtech.edu.cn

[f] Dr. Z. P. Jovanov
LIQUIDLOOP GmbH
Hardenbergstr. 38, 10623 Berlin, Germany

<figure><img src="image_3.png" /></figure>

Supporting information for this article is available on the WWW under
https://doi.org/10.1002/batt.202200055

<figure><img src="image_4.png" /></figure>

© 2022 The Authors. Batteries & Supercaps published by Wiley-VCH GmbH.
This is an open access article under the terms of the Creative Commons
Attribution License, which permits use, distribution and reproduction in any
medium, provided the original work is properly cited.
<figure><img src="image_5.png" /></figure>

lyte composition and the nature of the electrode surface.[12,31,36]
Na₂O₂ has been observed in the cell with acetonitrile and NaO₂
has been shown to disproportionate to Na₂O₂ in this solvent.[37]
This is ascribed to the low donor number of acetonitrile.
Attempts have been made to form Na₂O₂ in other electrolytes,
but these require conditions which enhance the surface bind-
ing of NaO₂, such as the use of Au electrodes, in conjunction
with high overpotentials.[38] Some Pt facets and Cu nano-
particles have also been shown to favour Na₂O₂ formation on
discharge.[39,40] Na₂O₂ has been observed when discharging at
high overpotential at a TEM grid, noting that this contains both
carbon and gold.[41] Hydrated peroxide (Na₂O₂·H₂O) and NaOH
have been reported in glyme-based electrolytes due to
contamination of the system with water from sources including
the O₂ gas supply and electrolyte.[14,42] Although, trace quanti-
ties of water have also been shown to promote NaO₂ solution
growth via phase-transfer catalysis.[43] During standard cycling
conditions (dry electrolyte with a carbon electrode), it is
generally assumed that Na₂O₂ does not form during discharge
in the Na–O₂ cell containing glyme ethers.

Here we explore the impact of overpotential on the
electrochemical reduction of oxygen in the glyme-based Na–O₂
batteries and the role this plays in product selection, NaO₂
versus Na₂O₂, and cell performance. The capacity is measured

at various rates and as is the distribution, morphology, and
chemistry of the discharged product at the positive electrode.
Electrochemical analysis is used to propose discharge routes
that explain these observations. The formation of thin Na₂O₂
films is conclusively demonstrated and shown to be detrimen-
tal to cell performance.

Results and Discussion

Impact of discharge rate on the capacity of the sodium-O₂
battery

To explore the impact of discharge rate, and thus overpotential,
on the discharge product and capacity of the Na–O₂ battery, a
series of Swagelok-type Na–O₂ cells were discharged at three
different current densities: 0.025, 0.25 and 0.4 mAcm⁻². The
electrolyte was 0.5 M sodium triflate (NaOTf) in diethylene
glycol dimethyl ether (DEGDME) and the electrode was carbon,
consistent with state-of-the-art cells.[³¹] Figure 1(a) shows the
discharge profiles collected from the Na–O₂ cells. At a lower
current density of 0.025 mAcm⁻², the battery achieved a
relatively large discharge capacity of over 4 mAhcm⁻², which is
typical for an Na–O₂ cell. The dominant formation of NaO₂ is

<figure><img src="image_6.png" /><figcaption>**Figure 1.** Discharge performance of Na–O₂ (Swagelok) cells: a) discharge profiles at three different current densities using O₂ saturated 0.5 M NaOTf in DEGDME; b) the corresponding PXRD patterns of the discharge positive electrodes. Three current densities are used: 0.025 mAcm⁻² (green line), 0.25 mAcm⁻² (purple line) and 0.4 mAcm⁻² (red line). Peaks marked with a black asterisk (*) correspond to the sample holder and GDL. SEM images of the positive electrode after discharge at c) 0.025 mAcm⁻²*areal* and d) 0.4 mAcm⁻²*areal*.</figcaption></figure>
<figure><img src="image_7.png" /></figure>

confirmed by powder X-ray diffraction (PXRD), Figure 1(b). Figure 1(c) shows an SEM image of the discharged positive electrode showing numerous large cube-shaped structures, consistent with the dominant formation of NaO₂ by a solution precipitation mechanism.[2,13,20] Upon increasing the current density to move to practical values, the voltage of the discharge plateaus shifted negative, discharge capacities decreased dramatically, and the cell death (rapid negative shift in voltage) rapidly occurred, Figure 1(a). This is also accompanied by an initial drop in potential, which we suggest is associated with the nucleation barrier for formation of NaO₂ on the carbon surface, Figure S1. SEM images, Figure 1(d), showed that the NaO₂ cubes became smaller and fewer in number at a higher current density. Previous studies have shown that current density has an impact on NaO₂ particle size.[20,31,44] PXRD indicates that the only product is NaO₂, Figure 1(b). Interestingly, cubic NaO₂ does not completely cover the electrode surface and large areas of apparent bare carbon remain. Passivation of the electrode surface by NaO₂ cubes at high rates cannot explain cell death.

## Soluble versus surface-grown discharge products

To investigate the electrochemistry of oxygen reduction in Na⁺-ioncontaining DEGDME in more detail, we employed a rotating ring-disk electrode (RRDE), Figure 2. Oxygen is reduced at the

rotating disk electrode ($i_{disk}$), while any soluble O₂⁻ formed is detected at a ring electrode ($i_{ring}$) by holding it at an oxidising potential (3.7 V vs. Na/Na⁺). The difference between the ring and the disk current can be assigned to insoluble products ($i_{surface}$), e.g., NaO₂(s), Na₂O₂(s), and can be used to identify the electrochemical potential at which a second reduction reaction occurs. NaO₂ does not disproportionate to form Na₂O₂ in Na⁺-ion containing DEGDME and therefore only electrochemical routes need to be considered.[37] The disk current rapidly increases between 2.5 V and 2.0 V vs. Na/Na⁺. At low overpotentials, $i_{disk}$ mirrors $i_{ring}$, indicating that the reaction is almost exclusively forming dissolved O₂⁻, rather than solid NaO₂. At higher overpotentials, $i_{ring}$ no longer fully accounts for $i_{disk}$, and this coincides with a peak in the current followed by a drop to zero current, demonstrating passivation of the electrode surface by the discharge product. The potential at which this occurs mirrors the sudden death potential of the Na-O₂ cells and indicates that it is responsible for the limited capacity seen at high overpotentials.

## Surface confined discharge products in the sodium-O₂ battery

To identify the nature of the passivating film, Na-O₂ cells were constructed using carbon nanofiber as the electrode (prepared as described in the Supporting Information), which are suitable

<figure><img src="image_8.png" /><figcaption>**Figure 2.** RRDE voltammetry in O₂ saturated 0.5 M NaOTf in DEGDME. Measurements were recorded at Au disk-Au ring electrode at a scan rate of 20 mV s⁻¹ and a rotation speed of 2000 rpm. *i*<sub>disk</sub> (red solid line) is the overall current measured, *i*<sub>ring</sub> (red dashed line) is the current due to the formation of soluble species and *i*<sub>surface</sub> (black solid line) is the difference between the two which is assigned to insoluble species.</figcaption></figure>
for use in high resolution transmission electron microscopy (HRTEM) measurements, such that it is possible to observe any nanometer thick discharge products that form at the electrode. The cells were discharged in O₂-saturated 0.25 M NaOTf in DEGDME at a high current density (0.4 mA cm⁻²) resulting in an overpotential of approximately 500 mV and the surface was analyzed using HRTEM, Figure 3. The images unambiguously demonstrate the formation of a 20–30 nm thick film of discharge product at the electrode surface. Previously it has been proposed that surface films of NaO₂ are responsible for sudden death at high discharge rates.[29] This is a reasonable conclusion as neither PXRD nor X-ray photoelectron spectroscopy (XPS) analysis typically identifies Na₂O₂ in discharged electrodes. However, previous analyses of the Li⁺-ion containing system have shown that electrochemically formed peroxide is non-crystalline, explaining why Na₂O₂ may not be detected by PXRD.[45] In addition, examination of SEM images, Figure 1, shows that NaO₂ cubes are the major discharge product in all cases making detection of trace Na₂O₂ difficult. Here we perform XPS analysis on a positive electrode discharged at a constant potential of 1.9 V vs. Na/Na⁺, where sudden death of the cell occurs. A flooded cell containing large volumes of electrolyte was used, such that much of the NaO₂ is dissolved and the relative signal for any insoluble products (i.e., Na₂O₂) is enhanced. The oxygen 1 s region of the resulting spectrum is shown in Figure 3(c). The spectrum displayed a major peak centred at 530.9 eV, confirming the formation of peroxide during discharge at high overpotential.[7,46] No peak consistent with the formation of Na₂O (529.7 eV) was observed.[47] The peak centred at 532.9 eV is due to the formation of superoxide and appears to make up a minor contribution to the overall XPS signal.[7] However, we reiterate that in a real cell with low electrolyte volume where NaO₂ is trapped in the electrode structure, NaO₂ is the major product, Figure S2, but only small amounts of Na₂O₂ are needed to have a detrimental impact on performance.

The data demonstrates the occurrence of two reaction processes, one at lower overpotential, which results in a soluble superoxide product and the formation of the anticipated Na₂O₂ cubes, and a second at a larger overpotential that results in the formation of passivating films of product that cover the electrode surface. The results are consistent with a 1e⁻ reduction of O₂ to soluble O₂⁻ stabilised by alkali ions in the solvation shell, followed by a second reduction step to Na₂O₂, Equations (1 and 2);

$$O_2(\text{sol}) + e^- + \rightarrow O_2^- (\text{sol}) \quad (1)$$

$$O_2^{-}(\text{sol}) + e^{-} + 2 \text{Na}^{+} \rightarrow \text{Na}_2\text{O}_2 \quad (2)$$

Onset of this second reduction will inevitably result in rapid electrode passivation and cell death and explains the trend seen during discharge of Na–O₂ cells at varying rates. These results suggest that Na₂O₂ formation is unavoidable at high overpotential presenting a challenge for the development of high-rate Na–O₂ batteries. Moreover, trace amount of Na₂O₂ could accumulate during cycling and its removal will likely

require high overpotentials during charging, eliminating many of the advantages of the Na–O₂ battery over the Li–O₂ battery. We note that changing the electrolyte solution or electrode

<figure><img src="image_9.png" /><figcaption>**Figure 3.** HRTEM images and spectroscopic analysis of the thin film product formed at the positive electrode of the Na–O₂ battery: a and b) HRTEM images showing the film deposited on carbon nanofiber electrodes during discharge at high rates (0.4 mA cm⁻²). c) X-ray photoelectron spectra in the O 1 s region of Na–O₂ positive electrodes discharged at a constant potential of 1.9 V vs. Na/Na⁺. The peaks at 530.9 eV, 532.9 eV and 536 eV are due to Na₂O₂, NaO₂ and the Na auger response, respectively.</figcaption></figure>
<figure><img src="image_10.png" /></figure>

material would alter the potential at which Na₂O₂ forms and
this could be a route to avoid its detrimental impact.

Conclusion

Here we show that the product of the aprotic oxygen reduction with Na⁺ ions in glyme ethers depends on the overpotential during discharge. At low overpotential, the reaction forms the 1e⁻ reduction product NaO₂, which is soluble in glyme-based electrolyte and results in the formation of cubic NaO₂ structures in the electrode and high capacities. At high overpotential, the reaction switches to a 2e⁻ reduction, which forms thin films of passivating Na₂O₂ across the surface of the electrode. If high rates are allowed to significantly polarise the voltage of Na–O₂ cells, the formation of Na₂O₂ passivating layers result in capacity limitation and rapid cell failure. To achieve high capacity and high rates, processes that result in polarisation of the electrode potential must be minimised. Practical cells will require high rates of electrochemical reduction of oxygen, together with rapid nucleation and growth of NaO₂, and high mass-transport of oxygen and Na⁺ ions, thus avoiding voltage polarisation and detrimental Na₂O₂ formation.

Experimental Section

Materials and Methods

Diethylene glycol dimethyl ether (Diglyme or DEGDME) was distilled under argon over CaH₂. The solvent was further dried for several days over freshly activated molecular sieves (type 4 Å). The final water content was ≤ 4 ppm (determined by Karl Fischer titration). Tetrabutylammonium trifluorosulfonate/triflate (TBAOTf) was recrystallised from a mixture of dichloromethane and diethyl ether and then dried under vacuum at 80 °C for 3 days. Sodium triflate (NaOTf) was recrystallised from a mixture of dimethyl sulfoxide (DMSO) and dichloromethane and then dried under vacuum at 80 °C for 3 days. All materials were stored in an argon glove box. High purity N5 O₂ was obtained from BOC industrial gases and was used in all measurements. During measurements, O₂ was passed through a drying column consisting of freshly activated molecular sieves before entering the cell. Bubbling O₂ into the electrolyte solution for 10 mins did not change the electrolyte water content.

Electrochemical Analysis

Electrochemical measurements were performed using a VMP3 electrochemical workstation (Biologic) and a multi-necked, air-tight glass cell within a glove box. Electrochemical measurements were carried out at room temperature and IR correction was used. 3 mm diameter polycrystalline Au disks (BAS Inc.) were employed as the working electrodes. Prior to use, the working electrodes were polished in a glove box with 0.05 µm alumina slurry in ethanol and rinsed with copious amounts of ethanol followed by drying under vacuum. Platinum wire served as the counter electrode. Measurements were performed using a reference electrode based on LiFePO₄. It was first pre-oxidized (20% of total capacity) to Li₃FePO₄, which has a constant potential. This is then placed in 0.1 M LiClO₄ in DMSO separated from the bulk electrolyte with a porous frit,

thus providing a fixed potential. Potentials were corrected to Na/
Na⁺ using an internal ferrocene reference and by subtracting 3.6 V.
Electrolytes were saturated with O₂. Carbon nanofibre electrodes
were formed from a nanofiber, PTFE suspension in isopropanol
which was cast onto stainless steel mesh, dried under vacuum, and
discharged at constant current in a three-electrode cell. Rotating
ring/disk electrode, RRDE, measurements were performed using a
MSR rotator and a removable disk rotating ring/disk electrode
containing an Au disk (5 mm diameter) and a glassy carbon ring
(PINE instruments). A ring potential of 3.7 V vs. Na/Na⁺ was used.
Measurements were performed in a glove box in a round bottom
flask and all necks other than the neck containing the shaft were
sealed. The electrode was polished as described above between
measurements. The collection efficiency at the ring, N, in each
experiment was determined using an ideal redox couple.

Na–O₂ cells were built from a modified Swagelok setup. The cell was vacuum-dried under elevated temperature and transferred into a glovebox (0.1 ppm H₂O, 0.1 ppm O₂). The anode was made from metallic sodium (Sigma-Aldrich), cut into a 0.5 cm² disc. 0.3 mL Dried electrolyte solution (0.5 M sodium salt in diglyme) was imbibed on two Whatman glass fibre filters (QM-A grade) (dried under vacuum at 260 °C, 24 h). A piece of GDL (1.13 cm² of surface area, 210 µm in thickness, and a weight of 10 mg) was used as cathode and held by an inox mesh current collector containing holes for gas exchange. No further insulating sleeve was used. The assembled cells were transferred from the glovebox to a filling station, and after a first evacuating step, the cells were pressurized with dry, ultrapure O₂ to 1.3 bar. To guarantee stable temperature conditions, the cells were mounted inside of a temperature-controlled incubator (25.0 ± 0.1 °C). The electrochemical measurements were performed under temperature-controlled conditions (25.0 °C) after resting for 4 h at the open-circuit voltage and using a Bio-Logic VMP3 potentiostat.

Characterization

X-ray photoelectron spectroscopy was performed on a Thermo Scientific K-Alpha X-ray photoelectron spectrometer (XPS) instrument equipped with a microfocussed monochromated Al X-ray source. The source was operated at 12 keV and a 400 micron spot size was used. The analyser operates at a constant analyser energy (CAE) 200 eV for survey scans and 50 eV for detailed scans. Charge neutralization was applied using a combined low energy/ion flood source. Powder X-ray diffraction (PXRD) measurements were performed by washing samples with dried diglyme prior to PXRD analysis. A Bruker D8 Advance diffractometer with a Cu K<sub>α</sub> radiation source (λ<sub>1</sub> = 1.5405 Å, λ<sub>2</sub> = 1.5443 Å) and a Lynxeye XE detector was used to collect the PXRD patterns. The PXRD patterns were recorded for 30 min in the 2θ range of 20°–65°. A special airtight cell with a beryllium window was used to guarantee no ambient air contamination during PXRD measurements.

Acknowledgements

P.G.B. is indebted to the EPSRC including the SUPERGEN program for financial support. L.R.J. gratefully acknowledges support from the University of Nottingham's Beacon in Propulsion Futures and the EPSRC (EP/S001611/1).
<figure><img src="image_11.png" /></figure>

Conflict of Interest

The authors declare no conflict of interest.

Data Availability Statement

The data that support the findings of this study are available
from the corresponding authors upon reasonable request.

Keywords: diglyme · oxygen reduction reaction · passivating films · sodium peroxide · sodium superoxide

[1] X. Gao, Y. Chen, L. R. Johnson, Z. P. Jovanov, P. G. Bruce, *Nat. Energy* **2017**, *2*, 17118.

[2] P. Hartmann, C. L. Bender, M. Vracar, A. K. Duerr, A. Garsuch, J. Janek, P. Adelhelm, *Nat. Mater.* **2013**, *12*, 228–232.

[3] Z. Pan, Y. Chen, F. Li, T. Zhang, C. Liu, H. Zhou, *J. Power Sources* **2014**, *251*, 466–469.

[4] Q. Sun, H. Yadegari, M. N. Banis, J. Liu, B. Xiao, X. Li, C. Langford, R. Li, X. Sun, *J. Phys. Chem. C* **2015**, *119*, 13433–13441.

[5] C. Xia, C. Y. Kwok, L. F. Nazar, *Science* **2018**, *361*, 777–781.

[6] D. Aurbach, B. D. McCloskey, L. F. Nazar, P. G. Bruce, *Nat. Energy* **2016**, *1*, 16128.

[7] J.-H. Kang, W.-J. Kwak, D. Aurbach, Y.-K. Sun, *J. Mater. Chem. A* **2017**, *5*, 20678–20686.

[8] G. Cong, W. Wang, N.-C. Lai, Z. Liang, Y.-C. Lu, *Nat. Mater.* **2019**, *18*, 390–396.

[9] J. Lu, Y. J. Lee, X. Luo, K. C. Lau, M. Asadi, H.-H. Wang, S. Brombosz, J. Wen, D. Zhai, Z. Chen, D. J. Miller, Y. S. Jeong, J.-B. Park, Z. Z. Fang, B. Kumar, A. Salehi-Khojin, Y.-K. Sun, L. A. Curtiss, K. Amine, *Nature* **2016**, *529*, 377–382.

[10] Y. Chen, Z. P. Jovanov, X. Gao, J. Liu, C. Holc, L. R. Johnson, P. G. Bruce, *J. Electroanal. Chem.* **2018**, *819*, 542–546.

[11] M. Balaish, X. Gao, P. G. Bruce, Y. Ein-Eli, *Adv. Mater.* **2019**, *4*, 1800645.

[12] C. L. Bender, D. Schroder, R. Pinedo, P. Adelhelm, J. Janek, *Angew. Chem. Int. Ed.* **2016**, *55*, 4640–4649; *Angew. Chem.* **2016**, *128*, 4716–4726.

[13] P. Hartmann, M. Heinemann, C. L. Bender, K. Graf, R.-P. Baumann, P. Adelhelm, C. Heiliger, J. Janek, *J. Phys. Chem. C* **2015**, *119*, 22778–22786.

[14] R. Pinedo, D. A. Weber, B. Bergner, D. Schröder, P. Adelhelm, J. Janek, *J. Phys. Chem. C* **2016**, *120*, 8472–8481.

[15] D. Krishnamurthy, H. A. Hansen, V. Viswanathan, *ACS Energy Lett.* **2016**, *1*, 162–168.

[16] I. I. Abate, L. E. Thompson, H.-C. Kim, N. B. Aetukuri, *J. Phys. Chem. Lett.* **2016**, *7*, 2164–2169.

[17] C. J. Allen, J. Hwang, R. Kautz, S. Mukerjee, E. J. Plichta, M. A. Hendrickson, K. M. Abraham, *J. Phys. Chem. C* **2012**, *116*, 20755–20764.

[18] I. M. Aldous, L. J. Hardwick, *Angew. Chem. Int. Ed.* **2016**, *55*, 8254–8257; *Angew. Chem.* **2016**, *128*, 8394–8397.

[19] L. Johnson, C. Li, Z. Liu, Y. Chen, S. A. Freunberger, P. C. Ashok, B. B. Praveen, K. Dholakia, J.-M. Tarascon, P. G. Bruce, *Nat. Chem.* **2014**, *6*, 1091–1099.

[20] L. Lutz, W. Yin, A. Grimaud, D. Alves Dalla Corte, M. Tang, L. Johnson, E. Azaceta, V. Sarou-Kanian, A. J. Naylor, S. Hamad, J. A. Anta, E. Salager, R.

Tena-Zaera, P. G. Bruce, J. M. Tarascon, *J. Phys. Chem. C* **2016**, *120*, 20068–20076.

[21] R. Tatara, G. M. Leverick, S. Feng, S. Wan, S. Terada, K. Dokko, M. Watanabe, Y. Shao-Horn, *J. Phys. Chem. C* **2018**, *122*, 18316–18328.

[22] S. Kang, Y. Mo, S. P. Ong, G. Ceder, *Nano Lett.* **2014**, *14*, 1016–1020.

[23] B. Lee, D.-H. Seo, H.-D. Lim, I. Park, K.-Y. Park, J. Kim, K. Kang, *Chem. Mater.* **2014**, *26*, 1048–1055.

[24] S. Yang, D. J. Siegel, *Chem. Mater.* **2015**, *27*, 3852–3860.

[25] B. D. McCloskey, J. M. Garcia, A. C. Luntz, *J. Phys. Chem. Lett.* **2014**, *5*, 1230–1235.

[26] C. L. Bender, P. Hartmann, M. Vračar, P. Adelhelm, J. Janek, *Adv. Energy Mater.* **2014**, *4*, 1301863.

[27] S. Y. Sayed, K. P. C. Yao, D. G. Kwabi, T. P. Batcho, C. V. Amanchukwu, S. Feng, C. V. Thompson, Y. Shao-Horn, *Chem. Commun.* **2016**, *52*, 9691–9694.

[28] L. Lutz, D. Alves Dalla Corte, M. Tang, E. Salager, M. Deschamps, A. Grimaud, L. Johnson, P. G. Bruce, J.-M. Tarascon, *Chem. Mater.* **2017**, *29*, 6066–6075.

[29] J. E. Nichols, B. D. McCloskey, *J. Phys. Chem. C* **2017**, *121*, 85–96.

[30] Y. Li, H. Yadegari, X. Li, M. N. Banis, R. Li, X. Sun, *Chem. Commun.* **2013**, *49*, 11731–11733.

[31] L. Lutz, D. A. D. Corte, Y. Chen, D. Batuk, L. R. Johnson, A. Abakumov, L. Yate, E. Azaceta, P. G. Bruce, J.-M. Tarascon, A. Grimaud, *Adv. Energy Mater.* **2018**, *8*, 1701581.

[32] M. J. Trahan, I. Gunasekara, S. Mukerjee, E. J. Plichta, M. A. Hendrickson, K. M. Abraham, *J. Electrochem. Soc.* **2014**, *161*, A1706–A1715.

[33] C. M. Burke, V. Pande, A. Khetan, V. Viswanathan, B. D. McCloskey, *Proc. Natl. Acad. Sci. USA* **2015**, *112*, 9293–9298.

[34] B. D. Adams, C. Radtke, R. Black, M. L. Trudeau, K. Zaghib, L. F. Nazar, *Energy Environ. Sci.* **2013**, *6*, 1772–1778.

[35] B. Horstmann, B. Gallant, R. Mitchell, W. G. Bessler, Y. Shao-Horn, M. Z. Bazant, *J. Phys. Chem. Lett.* **2013**, *4*, 4217–4222.

[36] Z. Zheng, C. Wu, Q. Gu, K. Konstantinov, J. Wang, *Energy Environ. Mater.* **2021**, *4*, 158–177.

[37] C. Sheng, F. Yu, Y. Wu, Z. Peng, Y. Chen, *Angew. Chem. Int. Ed.* **2018**, *57*, 9906–9910; *Angew. Chem.* **2018**, *130*, 10054–10058.

[38] S. Ma, W. C. McKee, J. Wang, L. Guo, M. Jansen, Y. Xu, Z. Peng, *Phys. Chem. Chem. Phys.* **2017**, *19*, 12375–12383.

[39] T. A. Galloway, J.-C. Dong, J.-F. Li, G. Attard, L. J. Hardwick, *Chem. Sci.* **2019**, *10*, 2956–2964.

[40] S. Han, C. Cai, F. Yang, Y. Zhu, Q. Sun, Y. G. Zhu, H. Li, H. Wang, Y. Shao-Horn, X. Sun, M. Gu, *ACS Nano* **2020**, *14*, 3669–3677.

[41] I. Landa-Medrano, A. Sorrentino, L. Stievano, I. Ruiz de Larramendi, E. Pereiro, L. Lezama, T. Rojo, D. Tonti, *Nano Energy* **2017**, *37*, 224–231.

[42] C. Liu, M. Carboni, W. R. Brant, R. Pan, J. Hedman, J. Zhu, T. Gustafsson, R. Younesi, *ACS Appl. Mater. Interfaces* **2018**, *10*, 13534–13541.

[43] C. Xia, R. Black, R. Fernandes, B. Adams, L. F. Nazar, *Nat. Chem.* **2015**, *7*, 496–501.

[44] N. Ortiz-Vitoriano, T. P. Batcho, D. G. Kwabi, B. Han, N. Pour, K. P. C. Yao, C. V. Thompson, Y. Shao-Horn, *J. Phys. Chem. Lett.* **2015**, *6*, 2636–2643.

[45] K. U. Schwenke, M. Metzger, T. Restle, M. Piana, H. A. Gasteiger, *J. Electrochem. Soc.* **2015**, *162*, A573–A584.

[46] H. Yadegari, Y. Li, M. N. Banis, X. Li, B. Wang, Q. Sun, R. Li, T.-K. Sham, X. Cui, X. Sun, *Energy Environ. Sci.* **2014**, *7*, 3747–3757.

[47] A. Barrie, F. J. Street, *J. Electron Spectrosc. Relat. Phenom.* **1975**, *7*, 1–31.

Manuscript received: January 31, 2022

Revised manuscript received: April 14, 2022

Accepted manuscript online: May 10, 2022

Version of record online: June 15, 2022