Material characterizations

The MnHCMn anode material is synthesized by the addition of aqueous sodium cyanide to aqueous manganese acetate. The rate of precipitation of MnHCMn from concentrated solutions is considerably slower than that of CuHCF, which results in large, sharply faceted cubic particles about 2 µm in size, as seen in the scanning electron microscope (SEM) image (Fig. 1b). X-ray diffraction (XRD) shows that the MnHCMn here has the cubic PBA structure (a 0 = 10.66 Å) (Fig. 1c). The composition of the MnHCMn was determined using inductively coupled plasma optical emission spectroscopy (ICP-OES) (ratio of Na to Mn) and Carbon, Hydrogen, and Nitrogen (ratio of C/N to Mn) analysis (Intertek) to be Na 1.24 Mn[Mn(CN) 6 ] 0.81 ·2.1H 2 O. The 13% wt. water content of MnHCMn was confirmed using thermogravimetric mass spectroscopy (TGA-MS) (Supplementary Fig. 1). We note that the balanced valence of Mn is 2+ in this formula. In particular, previously reported MnHCMn was synthesized from a sodium cyanide precursor, which leads to the monoclinic and rhombohedral structures with low vacancy15. Our materials feature a high HCMn vacancy concentration of 0.19, which contrasts the low-vacancy (0.07) monoclinic phase10. We note that the same cubic structured high-vacancy Mn(II)HCMn(III) was previously reported by Pasta et al7. Besides the differences on the morphology and the pristine material, i.e., Mn(II)HCMn(II) here, our synthesis is based on inexpensive NaCN, and the material is free from K+ that could lead to potential increase or loss of capacity. The production of the cubic MnHCMn phase is due to the limited use of cyanide during the synthesis, compared with the excessive use in previous works. Our reaction solution contains a 2.57:1 CN:Mn ratio, which is consistent with the cubic phase produced in ref. 7 (CN:Mn = 3.3) but contrasts the monoclinic phases reported in ref. 11,15 (CN:Mn > 5).

The CuHCF cathode material was synthesized by precipitation from aqueous solution at room temperature by a method similar to those previously reported6,7. Characterization of CuHCF via SEM and powder XRD show a poorly faceted, nanoparticulate (100–200 nm) morphology with the phase-pure Prussian blue structure (a 0 = 10.15 Å) (Supplementary Fig.s 2, 3). Analysis by ICP-OES and TGA-MS analysis shows a chemical formula of K 0.06 Cu[Fe(CN) 6 ] 0.67 ·3.8H 2 O (Supplementary Fig. 4).

Electrochemical property of MnHCMn anodes

We first characterize the electrochemical property (Fig. 2) and clarify the cycling mechanism (Fig. 3) of the cubic MnHCMn as a Na-ion battery anode through half-cell tests with SHE reference electrodes. The benefit of the used cosolvent electrolyte (1 M NaClO 4 , 90% acetonitrile (MeCN), 10% water) will be elaborated later in this work.

Fig. 2 Half-cell electrochemical characterization of MnHCMn vs. SHE. a The reaction potential profile of the MnHCMn electrode vs. SHE at a 1C rate. Electrolyte is 1 M NaClO 4 , 90% MeCN, 10% water. b The cycle life and coulombic efficiency of the MnHCMn electrode during 1C–1C cycling between –0.5 and –0.9 V Full size image

Fig. 3 sXAS probe of the Mn1+/Mn2+ and the strong hybridization in MnHCMn electrodes. a Mn L 3 -edge sXAS spectra collected on a series of MnHCMn electrodes at the charged (reduced), marked with “Ch”, and discharged (oxidized), marked with “D”, states of the first, second, and fortieth cycles. Calculated spectra of MnII(N), MnIII(C), MnII (C), and MnI(C) are plotted at the bottom. The sXAS features are assigned to Mn with different oxidation states at different coordination sites (dashed lines). b The N K-edge sXAS spectra collected on the same series of MnHCMn electrodes. The red and blue dashed lines indicate the chemical potential shifts between the charged and discharged electrodes. The black dashed line indicates an unoccupied low-energy electron states from the hybridization of the NC ligands and Mn-3d states at the discharged (oxidized) states. This hybridization state disappears due to the electron filling during the electrochemical charging (reducing) process Full size image

Figure 2a displays the cycling profile of MnHCMn vs. SHE between –0.5 and –0.9 V at 1C rate. The C-rates are defined with a reference capacity of 60 mAh g–1 (theoretical capacity 78 mAh g–1), which is typical for PBAs undergoing single electron transfer for redox reactions of the hexacyanometalate groups. The initial open circuit potential of the MnHCMn electrode in the half cell is approximately –0.1 V. The 1C cycling (CCCV charge, CV step to C/10) in the –0.5 to –0.9 V potential range yields a specific capacity of 67 mAh g–1 centered at –0.70 V, with 99.8% coulombic efficiency (Fig. 2b). This reaction corresponds to the intercalation of Na+ and the reduction of the HCMn groups:

$${\mathrm{Na}}_{{\mathrm{1}}{\mathrm{.24}}}{\mathrm{Mn}}\left[ {{\mathrm{Mn}}^{{\mathrm{II}}}\left( {{\mathrm{CN}}} \right)_{\mathrm{6}}} \right]_{{\mathrm{0}}{\mathrm{.81}}}{\mathrm{ + 0}}{\mathrm{.81\cdot }}\left( {{\mathrm{Na}}^{\mathrm{ + }}{\mathrm{ + e}}^{\mathrm{-}}} \right)\\ {\mathrm{ = Na}}_{{\mathrm{2}}{\mathrm{.05}}}{\mathrm{Mn}}\left[ {{\mathrm{Mn}}^{\mathrm{I}}\left( {{\mathrm{CN}}} \right)_{\mathrm{6}}} \right]_{{\mathrm{0}}{\mathrm{.81}}}$$ (1)

Two striking results can be seen from the electrochemical tests. First, the S-curve shape of the reaction profile indicates that the phase transformation is of a solid–solution-type reaction. Second, XRD characterization of electrodes at full charge and discharged showed that the structure remains cubic over its full state-of-charge (SOC) range with a small 1.2% decrease in lattice parameter during charging (Fig. 1c).

We note that both the solid–solution reaction and the stable structure of our cubic MnHCMn material is in contrast with the previous reports on the monoclinic MnHCMn phases, which show lower HCMn vacancy content, two-phase reactions, and a change in structure as the “A” site occupancy increases beyond 1 Na+ per interstice10,11. Retention of the cubic phase over the full SOC range of the reaction here is likely due to the high HCMn vacancy in our cubic material, which results in two Na+ per formula unit, i.e., full single occupancy of the interstitial A-sites (Fig. 1a).

Figure 2b shows the long-term 1C half-cell cycling of MnHCMn vs. SHE between –0.5 and –0.9 V. A linear loss rate of 18 ± 5 p.p.m. per cycle (N = 3, confidence interval 95%) is observed after >700 cycles and 2 months of testing. Extrapolating to 20% total capacity loss, this loss rate corresponds to a projected lifetime of 11,000 cycles and 2.5 years. In an effort to further improve the lifetime of the MnHCMn electrode, it was operated from –0.55 to –0.83 V, which is equivalent to a SOC range of 2.5–97.5%. After >900 1C cycles (test duration 3 months) in this range, a loss rate of 14 ± 1 p.p.m. per cycle (N = 4, confidence interval 95%) is observed (Supplementary Fig. 5). That improved loss rate projects to a lifetime of 14,000 cycles and 3.25 years to 80% capacity retention. We attribute the remaining capacity loss to the oxidation HCMn groups on the surface of the MnHCMn particles by a trace amount of oxygen to HCMn(III), which in turn is readily hydrolyzed16. Removal of the water from the cosolvent electrolyte would presumably eliminate this loss mechanism; however, the electrochemical characterization of MnHCMn(II/I) in a 1 M NaClO 4 , 100% MeCN electrolyte showed significantly poorer kinetics and lower capacity (Supplementary Fig. 6).

Mn1+/Mn2+ redox couple in MnHCMn

The low potential cycling of our cubic MnHCMn electrode vs. SHE indicates an intriguing redox reaction in this material. Previously, the corresponding reactions at the 0 and 0.7 V potentials upon Na/Na+ reference electrode were assigned to carbon-coordinated MnIII/II(C) and nitrogen-coordinated MnIII/II(N), respectively11. The mechanism is unclear for the –0.7 V reaction that is utilized here. No data have heretofore clarified the low-potential cycling mechanism and substantiated the favorability between MnII/I(C) and MnII/I(N)10,11. The lack of understanding stems from the lack of a direct probe of the oxidation states of the Mn(C) and Mn(N) with the necessary site sensitivity.

Very recently, sXAS is demonstrated to be a site-sensitive probe of the TM redox centers in PBA materials. Such sensitivity of sXAS could clarify the N- and C-coordinated TM cation valence and spin states in PBAs17. sXAS probes the unoccupied electronic states in the vicinity of the Fermi level, which are closely related with the crystal fields, chemical bonds, formal valences, orbitals, and spin characters of the material18. sXAS is a probe with both surface and bulk sensitivity through the total electron yield (TEY) and total fluorescence yield (TFY) detection channels with about 10 and 100 nm probe depths, respectively. In particular for Mn, the TM L-edge sXAS directly probes the Mn-3d states through dipole allowed 2p-3d transitions. Such a direct probe of Mn-3d states lead to distinct spectral lineshapes of Mn with different oxidation states, which enables detailed analysis of the Mn valences19,20,21,22. Utilizing sXAS for determining Mn states in PBAs is reliable because TMs in PBAs take well-defined spin states17, which lead to distinct sXAS spectral lineshape23. In this study, we employ both sXAS and RIXS to clarify the cycling mechanism of the Mn(C) and Mn(N) groups in our cubic MnHCMn electrodes. The two sets of spectroscopic results are consistent with each other and provide the first direct experimental probe of the low-spin 3d6 state of Mn, i.e., Mn1+(C). In this section, we first discuss the sXAS results.

Figure 3a shows the comparison between the Mn L 3 -edge sXAS spectra collected on a series of cycled MnHCMn electrodes and those from multiplet calculations with site sensitivity. Atomic multiplet calculations for Mn L-edge sXAS are based on a previously developed model that incorporates both forward- and back-bonding between Mn and the octahedrally coordinated CN ligands17,24 (see methods). The calculated MnII(N) L 3 spectrum shows a characteristic three-peak structure with a main peak b. This suggests a high-spin MnII system and is consistent with other high-spin MnII systems19. On the contrary, the stronger crystal field of carbon-coordinated Mn(C) encourages low spin states, consistent with the previous experimental reports25. The lineshape of the L 3 sXAS of MnII(C) changes dramatically from that of MnII(N) and is divided into two regions (a and e in Fig. 3a) separated by >4 eV. Especially, the very low energy of the state at 638.8 eV of MnII(C) derives from the single vacant t 2g level in a low-spin 3d5 configuration. Because sXAS detects the unoccupied electron states, such a low-energy MnII(C) state is to be first filled when electrons are introduced into the system (reduction). Therefore, the C-coordinated MnII(C) is energetically favored to be reduced. The calculated reduced MnI(C) calculation is dominated by a single L 3 peak at 643.4 eV, as expected for low-spin 3d6 configurations with a fully occupied t 2g level, outside of a narrow parameter space for backhanding energetics24.

The experimental sXAS data display a reversible lineshape evolution upon sodium intercalation/extraction in accordance with the excellent cycling stability of the material. As discussed above, the absorption features b and c (Fig. 3a) originate from MnII(N) and are of high-spin character. The high energy peaks (d and e) derive from the low-spin Mn(C). Feature a contains two states from the MnII contributions of both C- and N-coordinated sites, with the states of MnII(C) sitting at a slightly lower energy. All the experimental spectra exhibit the predominant absorption peak b at 640 eV, which stems from the high-spin MnII(N), suggesting that the MnII(N) is not electrochemically active during the cycling around –0.7 V vs. SHE.

The most profound spectral evolution appears in the energy range of 642.5–646 eV, where the low-spin character of the Mn(C) is dominant. For all the charged electrodes (reduced, −0.9 V), the low-spin Mn(C) display a peak (feature d) at 643.4 eV, in contrast to the broad feature group e (642–646 eV) that evolves in discharged electrodes (oxidized, –0.5 V). The emerging peak d in the charged (reduced) electrodes aligns well with the calculated monovalent Mn1+ of the carbon-coordinated Mn(C), thus demonstrating that the redox reactions occur on the Mn(C) with a nominal oxidation state of MnII(C) when discharged and MnI(C) when charged. The overall change in the intensity of feature a is consistent with the cycling of MnII(C)/MnI(C). Because MnI(C) does not contribute to the feature at about 638.8 eV, a relatively lower intensity of feature a is observed at the charged (reduced, MnI(C)) state.

In addition to the detailed lineshape analysis of the Mn-L 3 sXAS TEY data in Fig. 3a, the spectra with both the Mn L 3 - and L 2 -edges are shown in Supplementary Fig. 7. The bulk-sensitive TFY sXAS data are shown in Supplementary Fig. 8. Overall, both the full-energy range and TFY results are consistent with the detailed Mn-L 3 analysis above (Fig. 3a). Although the TFY spectra are distorted by the so-called self-absorption effects, the enhanced TFY Mn1+(C) feature at 643.4 eV is evident in the results, suggesting that the Mn1+ is the intrinsic bulk state of the charged (reduced) samples.

Direct probe of Mn1+/Mn2+ in MnHCMn

While the combination of the sXAS experimental and theoretical results provide strong evidence of the Mn state evolution of MnHCMn during electrochemical operation, the origin of the strong sXAS peak at 643.4 eV remains unclear and its assignment to Mn1+(C) relies on theoretical calculations. In order to provide direct experimental confirmation of the Mn1+/2+ redox couple, especially the novel Mn1+ state, we have employed RIXS mapping through the newly commissioned high-efficiency iRIXS system26. Technically, RIXS is collected by changing the incident X-ray photon energy (vertical axis in Fig. 4) across the interested absorption edge (Mn-L 3 edge here) and detect the inelastically scattered photon energy (horizontal axis in Fig. 4) through a spectrometer. It has been established that RIXS features correspond to various low-energy excitations27. For 3d TM elements, RIXS signals are dominated by the strong excitations within the 3d states, called “d-d excitations” (see Chapter 8 of ref. 28 and many references therein). Because different numbers of 3d electrons (TM valence) naturally lead to different d-d excitations, RIXS is a sensitive probe of the novel Mn1+ state with further resolved emission energy that is ignored in sXAS. In a localized atomic physics model, Mn1+(C) is a low-spin 3d6 electron system. Because of the crystal field splitting of the five 3d states into t 2g and e g , Mn1+(C) features a fully occupied t 2g (6 electrons) and empty e g state in particular. As schematically displayed in Fig. 4f, such a valence electron state configuration will naturally enhance the intra-3d excitation across the large crystal field splitting between t 2g and e g , because the full t 2g in Mn1+(C) does not allow other excitations within the t 2g as in Mn2+. This leads to an enhanced feature in RIXS with a relatively large energy loss (difference between the excitation and emission energies).

Fig. 4 Direct probe of Mn1+ by soft X-ray RIXS maps. a, b Mn L 3 -edge RIXS maps on MnHCMn electrodes at the discharged (oxidized) and charged (reduced) states, respectively. c, d Corresponding RIXS calculations with the nominal Mn valence marked on the maps. The strong RIXS intensity of the charged electrode (red arrow) is reproduced only with Mn1+ at the C-coordinated site, as shown in e. Calculation results with other configurations are shown in Supplementary Fig. 15. f A schematic to explain the characteristic Mn1+ RIXS feature, which fingerprints the excitations from the fully occupied t 2g to the nominally vacant e g orbitals (d-d) and delocalized ligand band states (CT) of a low-spin 3d6 state, i.e., Mn1+(C) Full size image

As clearly shown in Fig. 4a, b, the discharged (Mn2+) sample displays several groups of the d-d excitation features (red dashed lines in Fig. 4a) that are typical for a Mn2+ system with partially occupied t 2g and e g states29. In sharp contrast, the charged sample displays a greatly enhanced RIXS feature with a relatively high energy-loss value (away from the elastic line for about 5 eV) at the excitation energy of about 643.4 eV. This excitation energy matches exactly the energy of the key Mn1+(C) sXAS feature (Fig. 3a).

This is the first time low-spin 3d6, i.e., Mn1+, is directly fingerprinted through such strongly enhanced RIXS excitation feature, which is expected from a localized atomic state model with fully occupied t 2g states (Fig. 4f). Furthermore, our theoretical calculation confirms such assignment by reproducing the same contrast between the Mn2+(C)Mn2+(N) and Mn1+(C)Mn2+(N) system, as shown in Fig. 4c, d, with an enhanced 5 eV energy-loss feature from the Mn1+(C) configuration (Fig. 4e).

However, we note that RIXS is a two-step process, and inaccuracies in the simulations of the core-hole resonance and the final low energy excitations are compounded, therefore, it is hard to get as good an agreement with experiments as for sXAS. Although our theoretical calculations include the charge-transfer (CT) channels in the form of t 2g back-bonding17,24, features from delocalized states are largely underestimated and the atomic picture described above considers only the localized 3d states. The detailed analysis of the RIXS results is not a topic of this work, but as shown in Supplementary Fig. 16, the specific Mn1+ feature at 643.4 eV excitation energy displays an emission energy aligned with a long broad vertical feature around 638.3 eV emission energy in the full Mn-L RIXS map, which is typically called fluorescence in the RIXS community and represents the coupling of local multiplets to (ligand) band-like itinerant states. This indicates that this strong Mn1+ feature also contains significant contributions from CT between Mn 3d and the ligand band states. Additionally, a cross-over from d-d type to non-resonant fluorescence signals is seen around 641.7 eV excitation energy, which could be due to the quantum interference effects when the excitations involve two unoccupied states close in energy, e.g., Mn from different sites. However, the intensity drop of the quasi-elastic feature around 643.4 eV excitation energy (Supplementary Fig. 16) can be interpreted by the multiplet model and is due to the absence of a low-energy spin degree of freedom in the filled t 2g shell of Mn1+(C). Compared with the almost purely localized d-d excitations in the RIXS of the discharged (Mn2+) system (Fig. 4a), the Mn1+ feature (red arrow in Fig. 4b) aligns with the fluorescence feature and displays unusually strong intensity. This suggests a strong overlap of the ligand and Mn-3d states, i.e., enhanced hybridization, as also elaborated in the next section.

Mn1+ in such coordinated compounds has been proposed since 192830,31, and it was also speculated in other type (monoclinic/rhombohedral) of MnHCMn electrodes10,11; our consistent results of sXAS and RIXS here finally provide the long-awaiting experimental verification.

Strong hybridization and covalency in MnHCMn

A quantitative simulation of the sXAS of the fully charged and discharged MnHCMn samples is provided in Supplementary Fig. 9 through a linear combination of the calculated MnII(N), MnII(C), and MnI(C) spectra. The lineshape evolution upon the SOC can be seen from the simulated spectra that are in good agreement with the experimental results, which testifies the validity of our theoretical calculations. Interestingly, the calculated d electron occupancies, as listed in Supplementary Table 1, for MnII(C) and MnI(C) ions are 4.658 and 5.530, respectively. These values are much lower than the nominal values of 5 and 6 for Mn2+ and Mn1+. Such divergences of the electron occupation numbers indicates a strong covalency in the system, in which the hybridization between the Mn 3d and the C-N molecular orbitals results in the delocalization of the Mn-3d valence electrons32.

Because the electron delocalization is a critical parameter to determine the electronic conductivity, which is directly related with the rate performance of a battery material, we further testify the scenario of the strong hybridization and delocalization in the MnHCMn system by measuring the N-2p electron states through the N-K sXAS. Figure 3b displays the bulk-sensitive TFY signal of N-K sXAS of the same batch of cycled electrodes. The predominant absorption feature is located around 400 eV and can be assigned to transitions from N 1s core electrons to the unoccupied CN π* orbitals32,33. This feature is shifted slightly to higher energy for the discharged (oxidized) electrodes as compared with those of the charged (reduced) ones, consistent with the overall chemical potential change in the system. The most profound spectral evolution is found at around 397.5 eV, where the discharged (oxidized) electrodes exhibit an absorption peak that is attributed to the hybridized state of CN π* with the Mn 3d orbitals32. The variation of this low-energy N-K feature upon SOC is also shown in the surface-sensitive TEY spectra (Supplementary Fig. 10), but the stronger effect in the TFY data (Fig. 3b) indicates that this is a bulk property of the discharged (oxidized) MnHCMn.

Although a similar evolution in the C K-edge sXAS spectra are expected, the C K-edge data (Supplementary Fig. 11 and 12) are dominated by the signals from carbon black and binder on the surface of the composite electrodes (Supplementary Fig. 13). Nonetheless, the bulk-sensitive TFY data (Supplementary Fig. 11) show a weak but clear low-energy state only in the discharged (oxidized) samples, resembling that of the N-K lineshape evolution. During the Na+-intercalation charging (reducing) process, these low-energy hybridized state is favored in energy and gets filled in conjunction with the Mn 3d orbital, as indicated by its disappearance in the N-K and C-K sXAS data of the charged electrodes (Fig. 3b, S11).

The evolving low-energy features in the N-K and C-K sXAS data reveal the strong hybridization and thus the delocalization of the Mn-3d electrons. Additionally, the Mn1+ RIXS features also indicate strong overlap of the ligand and Mn-3d states, as briefly described above and will be detailed elsewhere. These evidences are consistent with the much lower Mn-3d orbital occupancy compared with the nominal values (Supplementary Table 1). The strong hybridization of the CN molecular and Mn 3d orbitals implies that the electron states in the MnHCMn material are intrinsically delocalized, which will fundamentally enhance the electron transport and improve the performance of the electrode at high cycling rates.

Co-solvent electrolyte and PBA solubility

While PBA-based electrodes are stable and of low cost, the material dissolution into aqueous electrolyte severely limits the lifetime of PBA electrodes, as reported previously for CuHCF6,13. Previously, electrolyte salt concentration was tuned to mitigate the solubility issue of CuHCF7. Here an aqueous–organic cosolvent electrolyte was developed to solve this problem. We employ ultraviolet-visual spectroscopy (UV-vis) to characterize the dissolved ferricyanide in aqueous and organic–aqueous electrolytes. Figure 5a shows the clear contrast on the solubility of CuHCF between the water and the 90% MeCN:10% water electrolyte. After the CuHCF electrode is soaked in the electrolyte for >48 h, the charged CuHCF has a solubility of >10 p.p.m. in water; however, in the organic–aqueous co-solvent solution, zero dissolution is detected. Electrochemical characterization of CuHCF cathode with our cosolvent electrolyte is performed in half-cells containing a 1 M NaClO 4 cosolvent electrolyte (90% MeCN, 10% water). Figure 5b displays a solid–solution reaction centered at 0.85 V vs. SHE and providing 62 mAh g−1 was observed, which is consistent with the oxidation and reduction of the hexacyanoferrate groups as previously observed in aqueous electrolytes6. The coulombic efficiency of the cycling is 99.9%. During cycling at a symmetric 1C rate (CCCV charge, CV step to C/10) between 0.6 and 1.15 V, which corresponds to a SOC range of approximately 0–95%, a specific capacity of 58 mAh g–1 was observed. Extended half-cell cycling under these conditions resulted in >2300 cycles with zero capacity loss during over 6 months of testing (Fig. 5c). This is in sharp contrast with the CuHCF electrode cycled in aqueous electrolyte, which typically shows clear capacity decrease within several days (Supplementary Fig. 17). This superior long cycle- and calendar-life was achieved because of the insolubility of CuHCF in our co-solvent electrolyte, which stands in contrast to the relatively short calendar lives of PBA cathodes in aqueous electrolytes6,13.

Fig. 5 Dissolution and half-cell tests of CuHCF in co-solvent electrolyte. a The UV-vis spectra of aliquots of water and the 90% MeCN + 10% water cosolvent electrolyte, in which a CuHCF electrode was stored at room temperature for at least 48 h. A 10 p.p.m. K 3 Fe(CN) 6 standard sample is measured for comparison. The peak at 420 nm corresponds to the dissolved Fe(CN) 6 3–. No dissolution is observed for the co-solvent electrolyte. b The reaction potential profile of the CuHCF cathode at a 1C rate with the co-solvent electrolyte. c The cycle life and coulombic efficiency of the CuHCF electrode during 1C–1C cycling between 0.6 and 1.15 V over 6 months. The noise is due to the precipitation of salt in the Ag/AgCl reference electrode frit, resulting in an inconsistent junction potential between the reference electrode-filling solution and the electrolyte Full size image

An all-PBA low-cost Na-ion battery full cell. Full pouch cells containing the MnHCMn anode, the CuHCF cathode (1.5× capacity), and the cosolvent electrolyte were cycled between 1.74 and 1.23 V (Fig. 6a). This full-cell voltage range resulted in 95% anode capacity utilization, 99.8% coulombic efficiency, and 97.7% round-trip energy efficiency during symmetric 1C cycling. During a 12C discharge from full charge, the cell delivers 74% of its total energy >1.23 V. Even with the symmetric 12C cycling, the cell still provides 55% of its total energy with 88% round-trip energy efficiency. The lost energy is due to the high internal resistance of the prototype pouch cell, which contained a mesh reference electrode between the active electrodes. We expect that the optimization of cell design for low impedance will further improve the energy retention during rapid cycling. At the 1C rate, by assuming 100% capacity utilization, 2/3 active cell mass, and 2 g mL–1 cell density, the full cell delivers a theoretical specific energy and energy density of 33 Wh kg–1 and 66 Wh L–1, respectively. We note that engineering optimization of the battery device is necessary to achieve this goal. Although this energy density is lower than that of Li-ion cells, it is comparable to that of the lead acid cells presently used in stationary storage systems.

Fig. 6 Na-ion battery full-cell tests based on CuHCF cathode, MnHCMn anode, and co-solvent electrolyte. a The voltage profiles of the full CuHCF/MnHCMn cell during 1C–1C cycling, 12C discharge from 1.74 V, and 12C-12C cycling. The full-cell-specific energy and energy density are 33 Wh kg–1 and 66 Wh L–1, respectively. b The cycle life and coulombic efficiency of the full cell during 1C–1C cycling between 1.74 and 1.23 V. For the first 400 cycles, galvanostatic charging with no potential hold (CV step) was used. After cycle 400, a CV step to a C/10 current, typically of 1–2 min in duration, was performed. The oscillations in capacity and coulombic efficiency are diurnal and correspond to temperature fluctuations from 23 to 25 °C in the cell testing station Full size image

After 1000 cycles at 1C (3.5-month test duration) utilizing 95% of the total anode capacity, the full cell retains 95% of its initial discharge capacity, which projects to 4000 cycles to 80% capacity retention (Fig. 6b). The full-cell capacity loss rate of 50 p.p.m. per cycle is approximately three times the capacity observed during the half-cell anode test. The higher loss rate observed in the full cell is attributed to at least two sources. First, oxygen may slowly diffuse into even a well-sealed pouch cell due to the permeability of the polymer laminate to oxygen. Additionally, it is possible that trace amount of oxygen is generated at the cathode at high potentials and migrates to the anode. Both effects could result in the oxidation of the active material surface to HCMn(III), which may react with water to form manganese oxide phases16. Second, any imbalance in the coulombic efficiencies of the two electrodes results in a “slip” in their relative SOC ranges, which results in a decrease in the accessible full-cell capacity within the fixed upper and lower voltage cutoffs. Nonetheless, the full cell data shown here represent a full Na-ion battery cell that achieves high-rate and long-life electrochemical cycling and with space for further optimizations in design and balancing.

Furthermore, a recent analysis of Li-ion cells showed that the active materials comprised about 50% of the total manufactured cell cost34. For the all-PBA cells here, the material cost is significantly lower. The other manufacturing costs will be further lowered because the electrolyte contains water, so the costly dry room is not required (see P.A. Nelson et al.35, Modeling the Performance and Cost of Lithium-Ion Batteries for Electric-Drive Vehicles, 2011). Therefore, the developed all-PBA Na-ion battery system provides a unique opportunity for dramatically decreasing the cost for grid-scale energy storage (see Supplementary Notes 1 and 2 for detailed cost estimation).