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To confirm the improved energy density, a hybrid supercapacitor was fabricated based on a solid-state 400-KOH-Ti3C2 anode and an active catholyte containing Mn2+ in 2 M H2SO4 electrolyte. Before assembling the hybrid supercapacitor, the CF as a current collector was pretreated by electrochemical predeposition with 3 mA h cm−2 MnO2 on the surface. The electrochemical performance of the as-assembled hybrid device was systematically investigated to demonstrate its unique advantages. As shown in Fig. 4a, in situ potential detection was performed, in which the Ag/AgCl reference electrode was introduced into the two-electrode system to explore in situ the potential variation of the respective electrode. This hybrid supercapacitor achieves a wide voltage window up to 1.7 V (red line), benefiting from the potential difference between the solid 400-KOH-Ti3C2 anode (dark cyan line) and active catholyte containing Mn2+ (black line). It is worth mentioning that such a wide operating voltage window is superior to that of most recently reported hybrid supercapacitors (see Table S2†) and even comparable to that of the “water in salt” electrolyte-based hybrid capacitors.Fig. 4b displays the CV curves of the hybrid supercapacitor at various scan rates. All the CV curves exhibit an approximate rectangular shape, which suggests that the operation mechanism in this hybrid device is bonding/debonding-induced pseudocapacitance in nature. Additionally, it is found that the addition of Mn2+ does not significantly affect the reaction kinetics of the 400-KOH-Ti3C2 anode (see Fig. S16† and the corresponding discussions). On the other hand, the near linear symmetric triangular charge/discharge curves at various current densities also confirm the pseudocapacitance characteristics of the hybrid device (Fig. 4c), which agrees well with the CV results. The corresponding specific capacitance was calculated based on the charge–discharge curves (see detailed calculations in the ESI†). Impressively, the hybrid device delivers an appreciable specific capacitance of 312.8 F g−1 at 1 A g−1 and 131.2 F g−1 at 80 A g−1, showing good rate performance (Fig. S17†). The energy and power densities of the present hybrid supercapacitor were also calculated on the basis of the total mass of 400-KOH-Ti3C2 and consumed MnO2, which is given as a Ragone plot profile (Fig. 4d). As displayed in Fig. 4d, the hybrid supercapacitor exhibits a maximum energy density of 43.4 W h kg−1 at a power density of 488.7 W kg−1. Even at a high power density of 40 kW kg−1, the energy density still remains 18.2 W h kg−1. It is worth noting that the energy density achieved here is markedly superior to those of most recently reported hybrid supercapacitors, as listed in Table S3.† These facts indicate that benefiting from both the high operating voltage and specific capacity of the active catholyte containing Mn2+, the energy density of the hybrid supercapacitor is significantly enhanced. In addition, this hybrid supercapacitor exhibits prominent cycling stability with a high-capacitance retention of 75% over 20000 cycles at 4 A g−1 (Fig. 4e). The charge/discharge curves at different cycles are shown in Fig. S18† to illustrate the high coulombic efficiency of the hybrid supercapacitor.

What's the anode?

400-KOH-Ti3C2

As seen from the surface morphology images of these three electrodes (Fig. 2 and S5†), they have their own characteristics and are all totally different from the original IF (Fig. S6†). As shown in Fig. 2a, there are abundant FeCO3 cubes (8–15 μm in edge length) on the surface of the FeCO3@IF electrode. The TEM images (Fig. 2b) show that the width of the crystal lattice is about 0.278 nm, matching the (104) face of FeCO3, which confirms the formation of FeCO3 and coincides with the results in the XRD patterns and XPS spectra. According to the SEM images in Fig. 2a, S5a and b,† the FeCO3 cubes do not completely cover the surface of the IF. As shown in Fig. 2c, the surface of the area without FeCO3 cubes is rough and is clearly different from the smooth surface of IF (Fig. S6†). According to the distribution of elements (C, O and Fe) in the area without FeCO3 cubes, C, O and Fe are uniformly distributed (Fig. S7b–d†). The EDS data (Fig. S7e†) confirm that the C, O and Fe content (wt%) is 8.63%, 35.58% and 55.79%, respectively. Therefore, the atomic ratio of C, O and Fe is 1.00:3.09:1.38, which is close to the atomic ratio of FeCO3. The slightly high iron content may be due to the iron foam substrate. These results confirm that the area without FeCO3 cubes is also covered by FeCO3 compounds. In addition, the TEM results in Fig. 2 further show that aside from the FeCO3 cubes with high crystallinity according to their long-range ordered crystal structure (Fig. 2b), some compounds with short-range ordered and long-range disordered crystal structures are also observed (Fig. 2d). The widths of the crystal lattices are about 0.278 nm and 0.234 nm, which match the (104) and (110) faces of FeCO3, respectively. Therefore, the compounds on the surface of IF without FeCO3 cubes are mainly FeCO3 but with low crystallinity according to the EDS and TEM results. Unlike FeCO3@IF, Fe3O4@IF is covered with tightly packed Fe3O4 nanoparticles (Fig. 2c, S5c and d†). The morphology of Fe–O–M@IF is similar to that of FeCO3@IF, but many long and interlaced nanowires are distributed around the cubes. Moreover, the edge length of the cubes is 2–3 times that of FeCO3@IF, and the surface is uneven. Compared with the original IF with its smooth surface, the increase in roughness can improve the specific surface area (Fig. S8 and Table S2†) and promote the penetration of the electrolyte as well as mass transfer. Based on the above characterizations, three different 3D iron-based electrodes have been successfully prepared via changing the iron source.

What's the electrolyte?

Herein, we fabricated two artificial SEIs on Na metal anodes via similar chemical replacement reactions between Na metal and SnCl4 liquid or SnCl2 additive dissolved in diethylene glycol dimethyl ether (DGM). According to X-ray photoelectron spectroscopy (XPS) depth-profiling results, the SnCl4 liquid treatment leads to a uniform and high ionic diffusion artificial SEI consisting of a simultaneously generated Na–Sn alloy and NaCl, whereas the SnCl2 additive results in a disordered and heterogeneous artificial SEI (Na–Sn alloy, NaCl, Sn metal, organo-chloride, ROCO2Na and RCH2ONa) due to the heterogeneous reaction between the Na metal anode and SnCl2/DGM solvent. It is found that the cycling performance of the SnCl4–Na electrodes is highly superior to that of the SnCl2–Na electrodes, which is mainly ascribed to the uniformity and ionic diffusion distinction of the two artificial SEIs. The importance of homogeneous and high ionic diffusion properties for the artificial SEI is also demonstrated by simulation results. Overall, the artificial SEI of SnCl4–Na electrodes can afford three key advantages: (i) the simple and controllable fabrication leads to a homogeneous artificial SEI consisting of uniformly distributed Na–Sn alloy and NaCl, avoiding the non-uniform products caused by the side reactions that commonly exist in electrolyte additive methods; (ii) both Na–Sn alloy and NaCl possess much higher diffusion coefficients than organic components (ROCO2Na and RCH2ONa) and conventional SEIs, enabling fast ion diffusion and restraining the initiation of Na dendrites; (iii) the high Young's moduli of the Na–Sn alloy and NaCl can also help suppress the growth of Na dendrites. With the above excellent properties, the SnCl4–Na electrode provides an ultralong cycling stability with a stable voltage polarization (∼100 mV) for 4000 h with a cycling capacity of 3 mA h cm−2 at 2 mA cm−2. Moreover, the SnCl4–Na electrode also shows excellent cycling stability for ∼1500 h even under rather tough conditions (5 mA h cm−2, 5 mA cm−2). In addition, benefiting from the durable SEI layer, a SnCl4–Na|FeS2 full cell can deliver a stable capacity of ∼350 mA h g−1 for 380 cycles.

What's the anode?

Na metal

Herein, we fabricated two artificial SEIs on Na metal anodes via similar chemical replacement reactions between Na metal and SnCl4 liquid or SnCl2 additive dissolved in diethylene glycol dimethyl ether (DGM). According to X-ray photoelectron spectroscopy (XPS) depth-profiling results, the SnCl4 liquid treatment leads to a uniform and high ionic diffusion artificial SEI consisting of a simultaneously generated Na–Sn alloy and NaCl, whereas the SnCl2 additive results in a disordered and heterogeneous artificial SEI (Na–Sn alloy, NaCl, Sn metal, organo-chloride, ROCO2Na and RCH2ONa) due to the heterogeneous reaction between the Na metal anode and SnCl2/DGM solvent. It is found that the cycling performance of the SnCl4–Na electrodes is highly superior to that of the SnCl2–Na electrodes, which is mainly ascribed to the uniformity and ionic diffusion distinction of the two artificial SEIs. The importance of homogeneous and high ionic diffusion properties for the artificial SEI is also demonstrated by simulation results. Overall, the artificial SEI of SnCl4–Na electrodes can afford three key advantages: (i) the simple and controllable fabrication leads to a homogeneous artificial SEI consisting of uniformly distributed Na–Sn alloy and NaCl, avoiding the non-uniform products caused by the side reactions that commonly exist in electrolyte additive methods; (ii) both Na–Sn alloy and NaCl possess much higher diffusion coefficients than organic components (ROCO2Na and RCH2ONa) and conventional SEIs, enabling fast ion diffusion and restraining the initiation of Na dendrites; (iii) the high Young's moduli of the Na–Sn alloy and NaCl can also help suppress the growth of Na dendrites. With the above excellent properties, the SnCl4–Na electrode provides an ultralong cycling stability with a stable voltage polarization (∼100 mV) for 4000 h with a cycling capacity of 3 mA h cm−2 at 2 mA cm−2. Moreover, the SnCl4–Na electrode also shows excellent cycling stability for ∼1500 h even under rather tough conditions (5 mA h cm−2, 5 mA cm−2). In addition, benefiting from the durable SEI layer, a SnCl4–Na|FeS2 full cell can deliver a stable capacity of ∼350 mA h g−1 for 380 cycles.

What's the electrolyte?

Herein, we fabricated two artificial SEIs on Na metal anodes via similar chemical replacement reactions between Na metal and SnCl4 liquid or SnCl2 additive dissolved in diethylene glycol dimethyl ether (DGM). According to X-ray photoelectron spectroscopy (XPS) depth-profiling results, the SnCl4 liquid treatment leads to a uniform and high ionic diffusion artificial SEI consisting of a simultaneously generated Na–Sn alloy and NaCl, whereas the SnCl2 additive results in a disordered and heterogeneous artificial SEI (Na–Sn alloy, NaCl, Sn metal, organo-chloride, ROCO2Na and RCH2ONa) due to the heterogeneous reaction between the Na metal anode and SnCl2/DGM solvent. It is found that the cycling performance of the SnCl4–Na electrodes is highly superior to that of the SnCl2–Na electrodes, which is mainly ascribed to the uniformity and ionic diffusion distinction of the two artificial SEIs. The importance of homogeneous and high ionic diffusion properties for the artificial SEI is also demonstrated by simulation results. Overall, the artificial SEI of SnCl4–Na electrodes can afford three key advantages: (i) the simple and controllable fabrication leads to a homogeneous artificial SEI consisting of uniformly distributed Na–Sn alloy and NaCl, avoiding the non-uniform products caused by the side reactions that commonly exist in electrolyte additive methods; (ii) both Na–Sn alloy and NaCl possess much higher diffusion coefficients than organic components (ROCO2Na and RCH2ONa) and conventional SEIs, enabling fast ion diffusion and restraining the initiation of Na dendrites; (iii) the high Young's moduli of the Na–Sn alloy and NaCl can also help suppress the growth of Na dendrites. With the above excellent properties, the SnCl4–Na electrode provides an ultralong cycling stability with a stable voltage polarization (∼100 mV) for 4000 h with a cycling capacity of 3 mA h cm−2 at 2 mA cm−2. Moreover, the SnCl4–Na electrode also shows excellent cycling stability for ∼1500 h even under rather tough conditions (5 mA h cm−2, 5 mA cm−2). In addition, benefiting from the durable SEI layer, a SnCl4–Na|FeS2 full cell can deliver a stable capacity of ∼350 mA h g−1 for 380 cycles.

What's the anode?

Na metal

In this work, we carry out classical reactive molecular dynamics (MD) simulations for studying the initial formation of SEI films occurring by Li oxidation and simultaneous decomposition of electrolyte (salt and solvent) molecules in the liquid phase in contact with the Li-metal electrode. We use lithium hexafluorophosphate (LiPF6) as a salt because its reductive decomposition has been pointed as critical on the SEI formation and lithium trifluoromethanesulfonate (lithium triflate or LiTF) that has shown interesting performance in various systems. We study the behavior of the various system components (i.e. Li-metal anode slab, and salt and solvent molecules) when the electrolyte solution is put in contact with the Li metal surface. We follow the evolution of various events including the lithium metal expansion/dissolution, salt and solvent decomposition, and initial nucleation of the SEI intermediates and products as well as the electron exchange among the species. We aim to identify the effects of electrolyte composition on the Li-metal anode behavior and the SEI formation and growth at open circuit conditions. We focus only on the initial stages of SEI formation, concentrated in a specific part of the battery system, and do not examine the Li deposition events occurring when an ionic flux arrives at the anode during charge, and the effects of an applied field. To gain further understanding of the structures and mechanisms of the initial stages of SEI formation, density functional theory (DFT) calculations on initial Li–F were used to evaluate SEI fragments observed in MD simulations. Once optimized structures of such fragments are obtained, we investigated the fragment clustering processes found at initial stages of SEI nucleation. These simulations provide preliminary estimations for the energies of formation and clustering of LiF fragments.

What's the anode?

Li-metal

In this work, we carry out classical reactive molecular dynamics (MD) simulations for studying the initial formation of SEI films occurring by Li oxidation and simultaneous decomposition of electrolyte (salt and solvent) molecules in the liquid phase in contact with the Li-metal electrode. We use lithium hexafluorophosphate (LiPF6) as a salt because its reductive decomposition has been pointed as critical on the SEI formation and lithium trifluoromethanesulfonate (lithium triflate or LiTF) that has shown interesting performance in various systems. We study the behavior of the various system components (i.e. Li-metal anode slab, and salt and solvent molecules) when the electrolyte solution is put in contact with the Li metal surface. We follow the evolution of various events including the lithium metal expansion/dissolution, salt and solvent decomposition, and initial nucleation of the SEI intermediates and products as well as the electron exchange among the species. We aim to identify the effects of electrolyte composition on the Li-metal anode behavior and the SEI formation and growth at open circuit conditions. We focus only on the initial stages of SEI formation, concentrated in a specific part of the battery system, and do not examine the Li deposition events occurring when an ionic flux arrives at the anode during charge, and the effects of an applied field. To gain further understanding of the structures and mechanisms of the initial stages of SEI formation, density functional theory (DFT) calculations on initial Li–F were used to evaluate SEI fragments observed in MD simulations. Once optimized structures of such fragments are obtained, we investigated the fragment clustering processes found at initial stages of SEI nucleation. These simulations provide preliminary estimations for the energies of formation and clustering of LiF fragments.

What's the electrolyte?

In this work, we carry out classical reactive molecular dynamics (MD) simulations for studying the initial formation of SEI films occurring by Li oxidation and simultaneous decomposition of electrolyte (salt and solvent) molecules in the liquid phase in contact with the Li-metal electrode. We use lithium hexafluorophosphate (LiPF6) as a salt because its reductive decomposition has been pointed as critical on the SEI formation and lithium trifluoromethanesulfonate (lithium triflate or LiTF) that has shown interesting performance in various systems. We study the behavior of the various system components (i.e. Li-metal anode slab, and salt and solvent molecules) when the electrolyte solution is put in contact with the Li metal surface. We follow the evolution of various events including the lithium metal expansion/dissolution, salt and solvent decomposition, and initial nucleation of the SEI intermediates and products as well as the electron exchange among the species. We aim to identify the effects of electrolyte composition on the Li-metal anode behavior and the SEI formation and growth at open circuit conditions. We focus only on the initial stages of SEI formation, concentrated in a specific part of the battery system, and do not examine the Li deposition events occurring when an ionic flux arrives at the anode during charge, and the effects of an applied field. To gain further understanding of the structures and mechanisms of the initial stages of SEI formation, density functional theory (DFT) calculations on initial Li–F were used to evaluate SEI fragments observed in MD simulations. Once optimized structures of such fragments are obtained, we investigated the fragment clustering processes found at initial stages of SEI nucleation. These simulations provide preliminary estimations for the energies of formation and clustering of LiF fragments.

What's the anode?

Li-metal

Driven by the urgent demand for electrical energy storage in electric vehicles, rechargeable devices with high energy densities and long cycle life continue to attract great interest. Lithium-ion batteries (LIBs) have attracted increasing attention over the past few decades owing to many merits, including high energy density, no memory effect, and good safety. However, making further improvements in their energy density whilst maintaining high safety remains highly challenging, inhibiting the scale-up of their applications. Developing new electrode materials is one of the most efficient ways to achieve LIBs with high energy and high safety. Silicon (Si) is considered to be a promising replacement for the graphite anodes in conventional LIBs, due to its ultrahigh specific capacity (4200 mA h g−1, ten times that of graphite), low working potential (ca. 0.4 V versus Li+/Li), natural abundance, high safety and environmentally benign nature. Unfortunately, Si has a low intrinsic electronic conductivity and this restricts charge transfer within Si anodes during the charge–discharge process. Moreover, the large volume change (>300%) of the Si anode during the lithiation/delithiation process leads to drastic particle pulverization and uncontrollable growth of a solid electrolyte interphase (SEI), leading to electrical contact loss between the active materials themselves or the active materials and current collectors. All these inherent problems lead to Si anodes having poor rate performance, low coulombic efficiency (CE) and fast capacity decay on cycling.

What's the anode?

graphite

Driven by the urgent demand for electrical energy storage in electric vehicles, rechargeable devices with high energy densities and long cycle life continue to attract great interest. Lithium-ion batteries (LIBs) have attracted increasing attention over the past few decades owing to many merits, including high energy density, no memory effect, and good safety. However, making further improvements in their energy density whilst maintaining high safety remains highly challenging, inhibiting the scale-up of their applications. Developing new electrode materials is one of the most efficient ways to achieve LIBs with high energy and high safety. Silicon (Si) is considered to be a promising replacement for the graphite anodes in conventional LIBs, due to its ultrahigh specific capacity (4200 mA h g−1, ten times that of graphite), low working potential (ca. 0.4 V versus Li+/Li), natural abundance, high safety and environmentally benign nature. Unfortunately, Si has a low intrinsic electronic conductivity and this restricts charge transfer within Si anodes during the charge–discharge process. Moreover, the large volume change (>300%) of the Si anode during the lithiation/delithiation process leads to drastic particle pulverization and uncontrollable growth of a solid electrolyte interphase (SEI), leading to electrical contact loss between the active materials themselves or the active materials and current collectors. All these inherent problems lead to Si anodes having poor rate performance, low coulombic efficiency (CE) and fast capacity decay on cycling.

What's the electrolyte?

Driven by the urgent demand for electrical energy storage in electric vehicles, rechargeable devices with high energy densities and long cycle life continue to attract great interest. Lithium-ion batteries (LIBs) have attracted increasing attention over the past few decades owing to many merits, including high energy density, no memory effect, and good safety. However, making further improvements in their energy density whilst maintaining high safety remains highly challenging, inhibiting the scale-up of their applications. Developing new electrode materials is one of the most efficient ways to achieve LIBs with high energy and high safety. Silicon (Si) is considered to be a promising replacement for the graphite anodes in conventional LIBs, due to its ultrahigh specific capacity (4200 mA h g−1, ten times that of graphite), low working potential (ca. 0.4 V versus Li+/Li), natural abundance, high safety and environmentally benign nature. Unfortunately, Si has a low intrinsic electronic conductivity and this restricts charge transfer within Si anodes during the charge–discharge process. Moreover, the large volume change (>300%) of the Si anode during the lithiation/delithiation process leads to drastic particle pulverization and uncontrollable growth of a solid electrolyte interphase (SEI), leading to electrical contact loss between the active materials themselves or the active materials and current collectors. All these inherent problems lead to Si anodes having poor rate performance, low coulombic efficiency (CE) and fast capacity decay on cycling.

What's the anode?

Si

All electrochemical measurements were carried out on a CHI 660E electrochemical workstation. The electrochemical properties of the prepared electrodes were tested using a three-electrode system in 1 M H2SO4 aqueous electrolyte. A platinum plate and a saturated calomel electrode (SCE) served as the counter electrode and the reference electrode, respectively. N-PCNFAs and PCNFAs sandwiched tightly between two pieces of 304 stainless steel meshes were directly used as the binder/additive-free working electrodes held by a clamp. To investigate the electrochemical properties, cyclic voltammetry (CV), galvanostatic charge–discharge (GCD) and electrochemical impedance spectroscopy (EIS, a frequency response analysis over the frequency range from 100 kHz to 10 mHz) were carried out. The specific capacitance of the electrode material was calculated from discharge curves according to the following equation (eqn (1)): where Cm is the specific capacitance (F g−1), I is the constant discharge current (A), Δt is the discharge time (s), m is the total mass of the active material (g) and ΔV is the set potential window (V).

What's the electrolyte?

1 M H2SO4 aqueous

Semiconducting polymers swell in typical organic solvents allowing ion penetration and high transconductance. However, for biosensing applications it is essential to operate in aqueous environments. Due to the non-polar nature of polymers, this limits their ability to swell and thus limits performance in aqueous media. A biphasic electrolyte platform developed by Duong et al. incorporates two immiscible electrolytes, an organic layer in direct contact with the polymer and aqueous layer on top. The former is selected because it allows ion penetration into P3HT so it can operate as an OECT, whilst the less dense aqueous layer interfaces with the biological environment of interest. This development opens the potential to utilise a wide range of organic semiconducting polymers that have been previously developed over the past few decades for transistor and solar cell applications.

What's the electrolyte?

The anisotropic expansion of primary grains, especially at a high state of charge (SOC), results in microcracks in NCM-based secondary particles. The morphology changes of the LLO microspheres after 200 cycles at the cut-offs of 4.8 V and 4.5 V are studied by SEM in Fig. S3.† The LLO microspheres cycled at the cut-off of 4.8 V show some obvious cracks, especially for large particles. But for the LLO particles cycled at 4.5 V, there are no obvious cracks. The structural evolutions of the LLOs at the cut-offs of 4.8 V and 4.5 V are investigated by STEM and EELS, as shown in Fig. 4 and S4–S8.† The cross-sections of both cycled LLOs are shown in Fig. S8,† in which the porous structure can still be found. The elemental maps exhibit heterogeneous distribution of TM elements in both cycled LLOs (Fig. S5 and S6†). The intergranular space is mainly filled with F and C containing compounds while the surface is mainly coated with P containing compounds for both cycled LLOs. The low magnification HAADF-STEM image (Fig. 4(a)) shows some microcracks in the cycled LLOs at the cut-off of 4.8 V. The high resolution HAADF-STEM image with the FFT and inverse FFT patterns (Fig. 4(b and c)) confirms the entire transformation from layered to rock-salt phases on the surface of LLOs at the cut-off of 4.8 V. The LLOs at the cut-off of 4.5 V undergo transitions from layered to layered/rock-salt mixed and rock-salt phases over long-term cycling (Fig. 4(d–f)). These results confirm the mitigated phase transition for LLOs over long-term cycling at the low cut-off of 4.5 V, which is probably ascribed to the reduced side reaction on the electrode/electrolyte interface.

What's the electrolyte?

Hybrid organic–inorganic crystals are unique materials that have been attracting increasing interest in the last few years due to their physical properties and potential applications in different fields, such as photovoltaics, lithium batteries, and photomechanic materials, between others. Among these applications, the development of high conductivity solid state electrolytes for fuel cells, batteries and solar cells is one of the most explored areas, in comparison with the polymers electrolytes that generally show low conductivity. In this sense, they are comparable with the classical ceramic conductors, such as doped lithium titanium phosphate, that allows fast ion transport. However, these materials present the considerable disadvantage of fragility and, in some cases, more importantly toxicity. In order to overcome this problem, a new type of materials have been proposed: the plastic hybrid materials. Although plastic crystals were described in the 60's, there is no rule to predict which cation and anion combinations will yield plastic crystalline materials and which will form salts that melt before any rotator phase is achieved. It deserves to be noted that a plastic crystal has a long-range structural order but short-range disorder, which is typically due to the occurrence of rotational motions of the constituents. This plastic phase can be seen as a mesophase between the solid and liquid phases that is often found in molecules with globular structures. This means that they are symmetrical around their center (CH4, CCl4, NH4 pentaerythritol, etc.); or they make a sphere by rotation around an axis (cyclohexane, camphor, etc.). The combination of two or more units to form hybrid materials can be desirable to enhance the multifunctionality of the system, for example doping plastic crystalline phases with a lithium imide salt for application as solid electrolytes in lithium batteries. An alternative is to combine the physical properties by tuning different constituents: ferroelectricity from polar molecules and long-range magnetic order from magnetic anions.

What's the electrolyte?

Similar to graphene, carbon nanotubes (CNTs) with a unique 1D tubular structure and excellent electrical conductivity can also be introduced into electrospun CNFs in order to increase the electrochemical activity of the electrospun CNFs and influence their architecture. Growing CNTs on the surface of the electrospun CNFs could be an effective approach to enhance the capacitive performance of electrode materials in SCs. For instance, a symmetrical device based on CNF/CNT hybrid electrodes exhibited a specific capacitance of 3.35 ± 0.05 mF cm−2, which was 3.6 times higher than that of a pure CNF electrode-based device (0.91 ± 0.02 mF cm−2). A hierarchical nanostructure of CNFs/CNTs was synthesized through a CVD process by using C2H2 as the carbon precursor and Ni nanoparticles on CNFs as catalysts.Fig. 12c shows an SEM image of the CNFs/CNTs nanocomposites, demonstrating the formation of a hierarchical network structure consisting of densely grown CNTs on the surface of the CNFs. The obtained CNFs/CNT electrode showed a specific capacitance of 464.2 F g−1 at 0.5 A g−1 in 6 M KOH and excellent stability with 97% retention after 10000 cycles. The excellent performance was ascribed to the following aspects: (i) the intimate contact between the CNTs (to provide electrical conduction from the tip of each CNT to the edges and surfaces of the CNFs) and the CNFs; (ii) the interconnected and porous CNFs acting as an excellent path facilitated charge transfer and electrolyte permeation. Hierarchical CNFs/CNT hybrids were prepared via a tubular CVD growth process using Al/Fe composite catalysts. After KOH activation, the tips of CNTs in the hybrids can be well opened at 700 °C without changing the overall hierarchical structure. More interestingly, the tip-open CNFs/CNT hybrids showed a considerably improved specific capacitance, which increases to 3.3 times that of the pristine one. Moreover, sweep analysis indicated that the diffusion-type capacitance increases by 3.7 times while the Helmholtz-type capacitance increases by only 1.5 times, indicating that the tip-open CNTs contribute to ∼30% of the increase in double-layer capacitance. In another study, CNFs/CNT scaffolds were fabricated by the electrospinning technique combined with post-carbonization processes. The diameters of CNF skeletons, CNT density on CNF skeletons and length of CNTs were determined by the PAN concentration, Fe(acac)3 concentration in precursor solution and growth time of C2H2 flow, respectively. It was observed that a suitable size of CNF skeletons, CNT density and growth time of CNTs could greatly enhance the overall performance.

What's the electrolyte?

Garnet-structured oxide electrolytes (Li7La3Zr2O12, LLZO) have significant advantage of being chemically and electrochemically stable against Li metals and allow implementation in Li metal batteries. However, a short-circuit failure due to Li penetration through the LLZO electrolyte has remained a crucial issue for safety and is a major hurdle for Li-based batteries to overcome. In, we investigated a mechanism of Li dendrite formation for the crystalline Ta-doped LLZO (LLZTO) electrolyte by examining their energy band structures and defect states using reflection electron energy loss spectroscopy (REELS), scanning photoelectron microscopy (SPEM), and nanoscale charge-based deep level transient spectroscopy (Nano Q-DLTS) techniques. The experimental results revealed that the Schottky barrier height (SBH) was lowered by 0.5 eV due to defect states localized in grain boundaries and that the metallic Li propagation along the grain boundaries is caused by the SBH reduction. Based on analytical results, the laser annealing of LLZTO was performed via bandgap engineering method to suppress the Li dendrite formation by forming a mixed surface layer of amorphous LLZTO and Li2O2, which has a wide bandgap to block the electron injection into the grain boundaries. The electrochemical measurements of laser-treated LLZTO demonstrated that the stability and cycling performance were significantly improved. This study sheds light on the importance of electronic structure, in particular, the defect states to develop high-performance oxide solid electrolytes for Li metal batteries and the practicality of surface modification by laser treatment.

What's the electrolyte?

(Li7La3Zr2O12, LLZO)

Garnet-structured oxide electrolytes (Li7La3Zr2O12, LLZO) have significant advantage of being chemically and electrochemically stable against Li metals and allow implementation in Li metal batteries. However, a short-circuit failure due to Li penetration through the LLZO electrolyte has remained a crucial issue for safety and is a major hurdle for Li-based batteries to overcome. In, we investigated a mechanism of Li dendrite formation for the crystalline Ta-doped LLZO (LLZTO) electrolyte by examining their energy band structures and defect states using reflection electron energy loss spectroscopy (REELS), scanning photoelectron microscopy (SPEM), and nanoscale charge-based deep level transient spectroscopy (Nano Q-DLTS) techniques. The experimental results revealed that the Schottky barrier height (SBH) was lowered by 0.5 eV due to defect states localized in grain boundaries and that the metallic Li propagation along the grain boundaries is caused by the SBH reduction. Based on analytical results, the laser annealing of LLZTO was performed via bandgap engineering method to suppress the Li dendrite formation by forming a mixed surface layer of amorphous LLZTO and Li2O2, which has a wide bandgap to block the electron injection into the grain boundaries. The electrochemical measurements of laser-treated LLZTO demonstrated that the stability and cycling performance were significantly improved. This study sheds light on the importance of electronic structure, in particular, the defect states to develop high-performance oxide solid electrolytes for Li metal batteries and the practicality of surface modification by laser treatment.

What's the electrolyte?

LLZO (LLZTO)

Unique graphene-like metallic Co9S8 with a morphology comprised of interconnected porous nanosheets that form a 3D network has been successfully synthesized by Nazar's group through a microwave solvothermal approach based on the reaction of cobalt chloride and TAA in a mixed solvent of water and triethylenetetramine at 160 °C for 1.5 h. Microwaves utilize a solvent dipole–microwave interaction that leads to rapidly superheated regions, triggering rapid nucleation and growth of particles. These distinctive synthesis conditions result in the formation of a long-range nanosheet structure of Co9S8 with a broad pore size distribution. N2 adsorption/desorption analysis indicates a high BET surface area of 108 m2 g−1 and a very large pore volume of 1.07 cm3 g−1 with the majority of pores in the range of 1.2–10 nm and the remainder distributed over 10–80 nm, indicating a texture incorporating micro-, meso-, and macro-pores. This high surface area and pore volume along with hierarchical porosity are vital to induce high intrinsic LiPS adsorptivity and enhance electrolyte penetration across thick electrodes. Li et al. presented the preparation of a S/C–SnS2 composite by ball milling sublimed sulfur, Ketjen Black, and ultrathin SnS2 nanosheets, where the 2D ultrathin nanosheets, fabricated by a one-pot reaction of SnCl4, TAA, and acetic acid, bring about a remarkably large surface area to enhance immobilization capability for LiPSs. In addition to those conventional methods, some innovative technologies can also delicately control the surface area, porosity, and/or polarity of metal sulfides. As an example, Xiao et al. demonstrated a green water-steam-etched approach for the fabrication of H- and O-incorporated porous TiS2 (HOPT) via heat treatment of a commercial TiS2 ball-milling product. During heat treatment, a mist of droplets from ultrapure water is passed into a quartz tube containing the ball-milled TiS2 using Ar-carrier gas. Interestingly, the time and temperature of water steam etching play pivotal roles in modulating the porosity of HOPT. Increasing the temperature from 100 to 500 °C, the specific surface area and pore volume of HOPT increase from 12.7 to 47 m2 g−1 and from 0.075 to 0.25 cm3 g−1, respectively, resulting from the perforation of TiS2 caused by the introduction of hydrogen and oxygen into the S–Ti–S framework during etching. Note that the hydroxyl and thiol groups in HOPT significantly enhance the surface polarity of TiS2, which are also favorable for adsorbing the LiPSs, suppressing the shuttle effect. More laudably, this method possesses excellent generality and can be extended to other TMDs such as CoS2 and NbS2. As reported in a follow-up study, Xiao et al. presented another extremely similar multistep approach to construct sandwich-type NbS2@S@I-doped graphene, which also brings about enhanced polarity and binding affinity for layered NbS2 to ensnare LiPSs.

What's the electrolyte?

Electrochemical testing was performed in a 2032 coin-type cell using Na metal (Sigma Aldrich, USA) as the anode. Electrodes were fabricated by blending the prepared cathode powders (85 wt%), carbon black (10 wt%), and polyvinylidene fluoride (5 wt%) in N-methyl-2-pyrrolidone (Daejung Chem, Korea). The slurry was then cast on aluminum foil (Hohsen Corp., Japan) and pre-dried at 110 °C in an oven. Then, the electrode was further dried at 110 °C for 5 h in a vacuum oven, and the disks were punched out of the foil. The electrolyte solution was 0.5 M NaPF6 (Tokyo Chemical Industry, Japan) in a 1:1 volumetric mixture of ethylene carbonate (Sigma Aldrich, USA) and diethyl carbonate (Sigma Aldrich, USA) with 2 vol% fluoroethylene carbonate (Tokyo Chemical Industry, Japan). All cells were prepared in an Ar-filled glovebox (MBRAUN, Germany). The fabricated cathodes and sodium metal anodes were separated by a glass fiber (Advantec, USA) to prevent short circuiting. The loading amount of the active material for all electrodes was 3.0–4.0 mg cm−2 in the coin-type half-cell. The cells were typically tested in the constant current mode, within the voltage range of 2.0–4.3 V versus Na/Na+, where 1C = 150 mA g−1. For the full-cell test, pouch-type (3 × 5 cm) cells were fabricated and tested in the voltage range of 1.0–4.1 V at 15 mA g−1 at 25 °C. The loading amount of the active material was 9.5–10.0 mg cm−2. The anode was fabricated by blending hard carbon (provided by Aekyung Petrochemical, Korea) (80 wt%), carbon black (3 wt%), and polyvinylidene fluoride (17 wt%). The resulting slurry was covered over copper foil and dried at 110 °C for 5 h in a vacuum oven. The full cell balance was achieved by controlling the capacity ratio of anode to cathode (N/P ratio) at 1.15:1.

What's the cathode?

Electrochemical testing was performed in a 2032 coin-type cell using Na metal (Sigma Aldrich, USA) as the anode. Electrodes were fabricated by blending the prepared cathode powders (85 wt%), carbon black (10 wt%), and polyvinylidene fluoride (5 wt%) in N-methyl-2-pyrrolidone (Daejung Chem, Korea). The slurry was then cast on aluminum foil (Hohsen Corp., Japan) and pre-dried at 110 °C in an oven. Then, the electrode was further dried at 110 °C for 5 h in a vacuum oven, and the disks were punched out of the foil. The electrolyte solution was 0.5 M NaPF6 (Tokyo Chemical Industry, Japan) in a 1:1 volumetric mixture of ethylene carbonate (Sigma Aldrich, USA) and diethyl carbonate (Sigma Aldrich, USA) with 2 vol% fluoroethylene carbonate (Tokyo Chemical Industry, Japan). All cells were prepared in an Ar-filled glovebox (MBRAUN, Germany). The fabricated cathodes and sodium metal anodes were separated by a glass fiber (Advantec, USA) to prevent short circuiting. The loading amount of the active material for all electrodes was 3.0–4.0 mg cm−2 in the coin-type half-cell. The cells were typically tested in the constant current mode, within the voltage range of 2.0–4.3 V versus Na/Na+, where 1C = 150 mA g−1. For the full-cell test, pouch-type (3 × 5 cm) cells were fabricated and tested in the voltage range of 1.0–4.1 V at 15 mA g−1 at 25 °C. The loading amount of the active material was 9.5–10.0 mg cm−2. The anode was fabricated by blending hard carbon (provided by Aekyung Petrochemical, Korea) (80 wt%), carbon black (3 wt%), and polyvinylidene fluoride (17 wt%). The resulting slurry was covered over copper foil and dried at 110 °C for 5 h in a vacuum oven. The full cell balance was achieved by controlling the capacity ratio of anode to cathode (N/P ratio) at 1.15:1.

What's the anode?

Na metal

Electrochemical testing was performed in a 2032 coin-type cell using Na metal (Sigma Aldrich, USA) as the anode. Electrodes were fabricated by blending the prepared cathode powders (85 wt%), carbon black (10 wt%), and polyvinylidene fluoride (5 wt%) in N-methyl-2-pyrrolidone (Daejung Chem, Korea). The slurry was then cast on aluminum foil (Hohsen Corp., Japan) and pre-dried at 110 °C in an oven. Then, the electrode was further dried at 110 °C for 5 h in a vacuum oven, and the disks were punched out of the foil. The electrolyte solution was 0.5 M NaPF6 (Tokyo Chemical Industry, Japan) in a 1:1 volumetric mixture of ethylene carbonate (Sigma Aldrich, USA) and diethyl carbonate (Sigma Aldrich, USA) with 2 vol% fluoroethylene carbonate (Tokyo Chemical Industry, Japan). All cells were prepared in an Ar-filled glovebox (MBRAUN, Germany). The fabricated cathodes and sodium metal anodes were separated by a glass fiber (Advantec, USA) to prevent short circuiting. The loading amount of the active material for all electrodes was 3.0–4.0 mg cm−2 in the coin-type half-cell. The cells were typically tested in the constant current mode, within the voltage range of 2.0–4.3 V versus Na/Na+, where 1C = 150 mA g−1. For the full-cell test, pouch-type (3 × 5 cm) cells were fabricated and tested in the voltage range of 1.0–4.1 V at 15 mA g−1 at 25 °C. The loading amount of the active material was 9.5–10.0 mg cm−2. The anode was fabricated by blending hard carbon (provided by Aekyung Petrochemical, Korea) (80 wt%), carbon black (3 wt%), and polyvinylidene fluoride (17 wt%). The resulting slurry was covered over copper foil and dried at 110 °C for 5 h in a vacuum oven. The full cell balance was achieved by controlling the capacity ratio of anode to cathode (N/P ratio) at 1.15:1.

What's the electrolyte?

0.5 M NaPF6

Though oxygen anion redox processes remarkably boost the specific capacity of LLOs, they trigger some intractable issues, such as voltage/capacity fade, voltage hysteresis, low initial Coulombic efficiency, and inferior rate capability. It has been demonstrated that the voltage hysteresis is caused by the poor kinetics of anionic redox reactions and grows progressively with overoxidation of lattice oxygen anions. Continuous side reactions between overoxidized oxygen anions and organic electrolytes give rise to a thick cathode electrolyte interphase (CEI) film, which deteriorates the interfacial charge-transfer kinetics. Besides, the oxygen anion redox reactions greatly affect the stability of the cationic redox chemistry of these 3d-TM LLOs. Namely, the oxidation of oxygen anions is usually accompanied by irreversible TM migration into neighboring Li 3a vacancies, resulting in structural transformation from the layered to spinel and/or rock-salt phases, which further leads to capacity/voltage fade. Strategies for mitigation of these issues have focused on surface coating, inactive-ion doping, a composition-graded structure, and optimization of chemical composition. In particular, regulating the anionic redox behaviors by tuning the cation arrangement and oxygen stacking sequence has captured great attention recently. Although some intriguing progress has been made, it is a long way towards their real-world implementation in high energy-density LIBs.

What's the electrolyte?

For further confirmation of the above electrochemical reaction mechanism, the HRTEM images of the discharged/charged electrodes were recorded. The ex situ images (ESI Fig. S8†) clearly showed the presence of the lattice fringes related to the β-MnS phase, thereby confirming the mechanism proposed in the present study. From these results, it is evident that such reaction mechanisms demonstrate electrochemical stability for prolonged cycling of the proposed composite electrode. The early variation in the specific capacity of the present electrode during cycling (Fig. 4c) is mostly related to the gradual activation of nanostructured electrodes. However, it remains essential to identify the individual contribution of the two reactions (intercalation and conversion) in the present electrochemical regulation to explain the variation in the long-term cycling curve (Fig. 4c) of the present α-MnS@NS-C electrode. To achieve this, further studies mostly related to in situ characterizations including the galvanostatic intermittent titration technique (GITT) and potentiostatic electrochemical impedance spectroscopy (PEIS) combined with separate examinations of the voltage profiles over regular cycling spans are required. However, the high reversibility of the present intercalation-cum-conversion reaction mechanism is actively supported by the conductive carbon network formed from the co-doping of the N and S in graphitic carbon. The unique electrode morphology and carbon matrix can improve electrical connectivity. Moreover, they impart structural stability by accommodating the volume changes caused by structural transitions and prevent polysulfide dissolution into the electrolyte during consecutive lithiation/delithiation processes. This, in turn, can enhance the electrochemical performance and cyclability of the proposed composite electrode.

What's the electrolyte?

In addition to the room-temperature performance investigation, we also explored the possibility of operating the hybrid supercapacitor at low temperatures. Our previous research studies have demonstrated that rechargeable batteries using H3O+ ions as charge carriers can operate even at a low temperature of −70 °C. Therefore, it is highly possible that this hybrid supercapacitor can operate well in such cold environments. To illustrate this assumption, we evaluated its electrochemical performance at a low temperature of −70 °C. It can be observed from Fig. S19† that the acid electrolyte containing Mn2+ has frozen at −70 °C. However, the hybrid supercapacitor still displays a discharge specific capacitance of 163.6 F g−1 at 0.1 A g−1 and 62.2 F g−1 at 0.5 A g−1 (Fig. 5a). Additionally, the Ragone plot in Fig. 5b shows that the hybrid supercapacitor exhibits a maximum energy density of 27.4 W h kg−1 and power density of 605 W kg−1 at −70 °C. More surprisingly, the specific capacitance of the hybrid supercapacitor can maintain approximately 100% of the original capacitance over 1000 cycles at −70 °C (Fig. 5c). Such excellent low-temperature performance of the hybrid supercapacitor is associated with the high ionic conductivity of the electrolyte. The DSC measurements show that the freezing point of the acid electrolyte containing Mn2+ is −40.2 °C (Fig. S20†), which is lower than that of the 2 M MnSO4 solution (−11 °C, Fig. S21†). Surprisingly, it is found that this acid electrolyte containing Mn2+ can still show a decent ionic conductivity at −70 °C (2.6 mS cm−1), obtained from EIS data displayed in Fig. S22 (see the related calculations in the ESI†).

What's the electrolyte?

Herein, a Li/LiV2(PO4)3 primary battery was proposed and studied for the first time. The unique NASICON structure of monoclinic LiV2(PO4)3 facilitated Li+ diffusion, ensuring superior rate capability and low-temperature performance of the battery. However, the short shelf life of the Li/LiV2(PO4)3 primary batteries hinders their practical applications. The shelf life of the Li/LiV2(PO4)3 primary batteries was improved by optimizing the electrolyte composition. It was found that the corrosion of the Al foil triggered by the organic radical cations generated from the electrochemical oxidation of EC at high potentials leads to the self-discharge of the Li/LiV2(PO4)3 primary battery. When EC was replaced by PC, the corrosion of the Al foil was alleviated, and the shelf life of the Li/LiV2(PO4)3 primary battery was significantly improved. However, the loss of capacity still occurred after storage, which is ascribed to the oxidation decomposition of the electrolyte on the surface of the LiV2(PO4)3 cathode. A slow two-phase transition process from the LiV2(PO4)3 phase to the Li3V2(PO4)3 phase with the reduction of V4+ to low-valence V was characterized by the in situ XRD of the LiV2(PO4)3 cathode during its storage; this process was accompanied by the spontaneous insertion of Li back into LiV2(PO4)3. It was found that LiBOB as an additive effectively helps to improve the shelf life of the Li/LiV2(PO4)3 primary battery by alleviating the side reaction between LiV2(PO4)3 and electrolyte according to the XPS results, 100% of capacity could be maintained after one-month storage. Meanwhile, the Li/LiV2(PO4)3 primary battery exhibited superior rate capability and low-temperature performance: 86% energy could be maintained at 50C and 63% energy could be maintained at −40 °C and 0.1C, which was substantially better than the performance of the Li/MnO2 primary battery. Thus, our Li/LiV2(PO4)3 primary battery is an ideal candidate for applications in extreme situations.

What's the cathode?

LiV2(PO4)3

Herein, a Li/LiV2(PO4)3 primary battery was proposed and studied for the first time. The unique NASICON structure of monoclinic LiV2(PO4)3 facilitated Li+ diffusion, ensuring superior rate capability and low-temperature performance of the battery. However, the short shelf life of the Li/LiV2(PO4)3 primary batteries hinders their practical applications. The shelf life of the Li/LiV2(PO4)3 primary batteries was improved by optimizing the electrolyte composition. It was found that the corrosion of the Al foil triggered by the organic radical cations generated from the electrochemical oxidation of EC at high potentials leads to the self-discharge of the Li/LiV2(PO4)3 primary battery. When EC was replaced by PC, the corrosion of the Al foil was alleviated, and the shelf life of the Li/LiV2(PO4)3 primary battery was significantly improved. However, the loss of capacity still occurred after storage, which is ascribed to the oxidation decomposition of the electrolyte on the surface of the LiV2(PO4)3 cathode. A slow two-phase transition process from the LiV2(PO4)3 phase to the Li3V2(PO4)3 phase with the reduction of V4+ to low-valence V was characterized by the in situ XRD of the LiV2(PO4)3 cathode during its storage; this process was accompanied by the spontaneous insertion of Li back into LiV2(PO4)3. It was found that LiBOB as an additive effectively helps to improve the shelf life of the Li/LiV2(PO4)3 primary battery by alleviating the side reaction between LiV2(PO4)3 and electrolyte according to the XPS results, 100% of capacity could be maintained after one-month storage. Meanwhile, the Li/LiV2(PO4)3 primary battery exhibited superior rate capability and low-temperature performance: 86% energy could be maintained at 50C and 63% energy could be maintained at −40 °C and 0.1C, which was substantially better than the performance of the Li/MnO2 primary battery. Thus, our Li/LiV2(PO4)3 primary battery is an ideal candidate for applications in extreme situations.

What's the electrolyte?

Herein, a Li/LiV2(PO4)3 primary battery was proposed and studied for the first time. The unique NASICON structure of monoclinic LiV2(PO4)3 facilitated Li+ diffusion, ensuring superior rate capability and low-temperature performance of the battery. However, the short shelf life of the Li/LiV2(PO4)3 primary batteries hinders their practical applications. The shelf life of the Li/LiV2(PO4)3 primary batteries was improved by optimizing the electrolyte composition. It was found that the corrosion of the Al foil triggered by the organic radical cations generated from the electrochemical oxidation of EC at high potentials leads to the self-discharge of the Li/LiV2(PO4)3 primary battery. When EC was replaced by PC, the corrosion of the Al foil was alleviated, and the shelf life of the Li/LiV2(PO4)3 primary battery was significantly improved. However, the loss of capacity still occurred after storage, which is ascribed to the oxidation decomposition of the electrolyte on the surface of the LiV2(PO4)3 cathode. A slow two-phase transition process from the LiV2(PO4)3 phase to the Li3V2(PO4)3 phase with the reduction of V4+ to low-valence V was characterized by the in situ XRD of the LiV2(PO4)3 cathode during its storage; this process was accompanied by the spontaneous insertion of Li back into LiV2(PO4)3. It was found that LiBOB as an additive effectively helps to improve the shelf life of the Li/LiV2(PO4)3 primary battery by alleviating the side reaction between LiV2(PO4)3 and electrolyte according to the XPS results, 100% of capacity could be maintained after one-month storage. Meanwhile, the Li/LiV2(PO4)3 primary battery exhibited superior rate capability and low-temperature performance: 86% energy could be maintained at 50C and 63% energy could be maintained at −40 °C and 0.1C, which was substantially better than the performance of the Li/MnO2 primary battery. Thus, our Li/LiV2(PO4)3 primary battery is an ideal candidate for applications in extreme situations.

What's the cathode?

LiV2(PO4)3

The morphologies of NF, MoP/NF, NiCo-LDH/NF, and MoP@NiCo-LDH/NF-20 were also characterized by SEM (Fig. 3a, b and S2†). It can be seen that the folded lamellar structure is uniformly grown on and aggregated into a tremella shape on the smooth surface of NF, which can promote the increase in both specific surface area and active sites. In addition, the electrodeposition time was optimized to obtain MoP@NiCo-LDH/NF-x (Fig. S3†). When the electrodeposition time is 10 minutes, the MoP surface is coated with sparse NiCo-LDH nanosheets. With the increase of electrodeposition time, the MoP surface is coated with relatively dense NiCo-LDH nanosheets. However, as the electrodeposition time increases to 30 minutes, the NiCo-LDH nanosheets become more crowded. Too loose or too tight a structure is not conducive to the movement of ions in the electrolyte and charge transfer, thus affecting the performance of electrocatalysis, so MoP@NiCo-LDH/NF-20 should have relatively more active sites.

What's the electrolyte?

According to the SEM images in Fig. 4c and S16,† the basic skeleton of FeCO3@IF after OER remains almost the same compared with the electrode before OER. However, the FeCO3 cubes are covered with lots of nanosheets, which grow vertically on the cubes and are connected to each other. According to Fig. S17,† it is easy to observe the formation of a nanosheet array on the surface of the area of FeCO3@IF without FeCO3 cubes, which is the same as the surface morphology of the FeCO3 cubes after the OER process (Fig. 4c). The TEM image in Fig. 4d further confirms the formation of nanosheets. The blurry lattice fringes in the HRTEM image (the inset in Fig. 4d) suggest that the newly generated iron oxo/hydroxides have a high degree of amorphization. Therefore, the generated iron-based oxo/hydroxides are not observed in the XRD pattern (Fig. 4a). Notably, the high degree of amorphization is beneficial for the exposure of more catalytic sites and accelerating mass transfer. Furthermore, the phases with a high degree of amorphization also possess abundant unsaturated electronic configurations, which may improve the adsorption of reactants. The amorphous iron oxo/hydroxides are efficient catalytic sites for OER. More importantly, the vertical nanosheet arrays can promote the penetration of the electrolyte and bubble diffusion compared with cumulate nanoparticles and interlaced nanowires (Fig. S18†). The in situ generated interconnected nanosheets also have a strong interaction with the IF substrate, which is beneficial for maintaining good structural stability and high electron transfer efficiency. According to the SEM and XRD results, FeCO3 does not completely transform into oxo/hydroxide nanosheets, because of diffusional and electrochemical limitations. The partial surface in situ self-reconstruction causes FeCO3@IF to obtain novel and stable hierarchical structures during the OER process (Fig. 4e), with improved mechanical strength and rich diffusion pathways. Therefore, FeCO3@IF shows high OER catalytic activity and stability.

What's the electrolyte?

The use of Mg-alloyable metallic anodes is another promising approach to enhancing the viability of MIBs. This is mainly because the Mg alloying/dealloying process occurs slightly above the Mg plating/stripping potentials, and the surface passivation problems of Mg in conventional electrolytes is expected to be alleviated to a great extent. Indeed, reversible Mg alloying/dealloying behavior on bismuth has been confirmed in conventional electrolytes by several researchers (e.g., Bi nanoparticles in Mg(N(SO2CF3)2)2/acetonitrile; Bi nanotubes in Mg(N(SO2CF3)2)2/diglyme; and Bi nanocrystals in Mg(N(SO2CF3)2)2/diglyme). Bi anodes have shown reasonable cyclability and rate performance in a nanometric form. Mg-alloyable Sn was also studied for use as MIB anodes. Sn can deliver a higher specific capacity than Bi (903 mA h g−1 for Sn + 2Mg2+ + 4e− → Mg2Sn and 385 mA h g−1 for 2Bi + 3Mg2+ + 6e− → Mg3Bi2) and can alloy/dealloy Mg at a lower potential than Bi (0.15 and 0.25 V vs. Mg/Mg2+ for Sn and Bi, respectively). Despite the theoretical prediction that nanometric Sn can act as a high-capacity MIB anode, pure Sn has shown limited electrochemical performance. For example, though pure Sn delivered a capacity of ca. 900 mA h g−1 during the 1st magnesiation (complete conversion of Sn to Mg2Sn), the capacity abruptly dropped to ca. 200–300 mA h g−1 after the 1st magnesiation, and continuously declined with repeated charge/discharge (C/D). Furthermore, pure Sn often requires the application of potentials close to, or even lower than, the thermodynamic Mg plating potential for the activation of an alloying process.

What's the anode?

Bi

The use of Mg-alloyable metallic anodes is another promising approach to enhancing the viability of MIBs. This is mainly because the Mg alloying/dealloying process occurs slightly above the Mg plating/stripping potentials, and the surface passivation problems of Mg in conventional electrolytes is expected to be alleviated to a great extent. Indeed, reversible Mg alloying/dealloying behavior on bismuth has been confirmed in conventional electrolytes by several researchers (e.g., Bi nanoparticles in Mg(N(SO2CF3)2)2/acetonitrile; Bi nanotubes in Mg(N(SO2CF3)2)2/diglyme; and Bi nanocrystals in Mg(N(SO2CF3)2)2/diglyme). Bi anodes have shown reasonable cyclability and rate performance in a nanometric form. Mg-alloyable Sn was also studied for use as MIB anodes. Sn can deliver a higher specific capacity than Bi (903 mA h g−1 for Sn + 2Mg2+ + 4e− → Mg2Sn and 385 mA h g−1 for 2Bi + 3Mg2+ + 6e− → Mg3Bi2) and can alloy/dealloy Mg at a lower potential than Bi (0.15 and 0.25 V vs. Mg/Mg2+ for Sn and Bi, respectively). Despite the theoretical prediction that nanometric Sn can act as a high-capacity MIB anode, pure Sn has shown limited electrochemical performance. For example, though pure Sn delivered a capacity of ca. 900 mA h g−1 during the 1st magnesiation (complete conversion of Sn to Mg2Sn), the capacity abruptly dropped to ca. 200–300 mA h g−1 after the 1st magnesiation, and continuously declined with repeated charge/discharge (C/D). Furthermore, pure Sn often requires the application of potentials close to, or even lower than, the thermodynamic Mg plating potential for the activation of an alloying process.

What's the electrolyte?

Mg(N(SO2CF3)2)2/acetonitrile

The use of Mg-alloyable metallic anodes is another promising approach to enhancing the viability of MIBs. This is mainly because the Mg alloying/dealloying process occurs slightly above the Mg plating/stripping potentials, and the surface passivation problems of Mg in conventional electrolytes is expected to be alleviated to a great extent. Indeed, reversible Mg alloying/dealloying behavior on bismuth has been confirmed in conventional electrolytes by several researchers (e.g., Bi nanoparticles in Mg(N(SO2CF3)2)2/acetonitrile; Bi nanotubes in Mg(N(SO2CF3)2)2/diglyme; and Bi nanocrystals in Mg(N(SO2CF3)2)2/diglyme). Bi anodes have shown reasonable cyclability and rate performance in a nanometric form. Mg-alloyable Sn was also studied for use as MIB anodes. Sn can deliver a higher specific capacity than Bi (903 mA h g−1 for Sn + 2Mg2+ + 4e− → Mg2Sn and 385 mA h g−1 for 2Bi + 3Mg2+ + 6e− → Mg3Bi2) and can alloy/dealloy Mg at a lower potential than Bi (0.15 and 0.25 V vs. Mg/Mg2+ for Sn and Bi, respectively). Despite the theoretical prediction that nanometric Sn can act as a high-capacity MIB anode, pure Sn has shown limited electrochemical performance. For example, though pure Sn delivered a capacity of ca. 900 mA h g−1 during the 1st magnesiation (complete conversion of Sn to Mg2Sn), the capacity abruptly dropped to ca. 200–300 mA h g−1 after the 1st magnesiation, and continuously declined with repeated charge/discharge (C/D). Furthermore, pure Sn often requires the application of potentials close to, or even lower than, the thermodynamic Mg plating potential for the activation of an alloying process.

What's the anode?

Mg-alloyable Sn

The use of Mg-alloyable metallic anodes is another promising approach to enhancing the viability of MIBs. This is mainly because the Mg alloying/dealloying process occurs slightly above the Mg plating/stripping potentials, and the surface passivation problems of Mg in conventional electrolytes is expected to be alleviated to a great extent. Indeed, reversible Mg alloying/dealloying behavior on bismuth has been confirmed in conventional electrolytes by several researchers (e.g., Bi nanoparticles in Mg(N(SO2CF3)2)2/acetonitrile; Bi nanotubes in Mg(N(SO2CF3)2)2/diglyme; and Bi nanocrystals in Mg(N(SO2CF3)2)2/diglyme). Bi anodes have shown reasonable cyclability and rate performance in a nanometric form. Mg-alloyable Sn was also studied for use as MIB anodes. Sn can deliver a higher specific capacity than Bi (903 mA h g−1 for Sn + 2Mg2+ + 4e− → Mg2Sn and 385 mA h g−1 for 2Bi + 3Mg2+ + 6e− → Mg3Bi2) and can alloy/dealloy Mg at a lower potential than Bi (0.15 and 0.25 V vs. Mg/Mg2+ for Sn and Bi, respectively). Despite the theoretical prediction that nanometric Sn can act as a high-capacity MIB anode, pure Sn has shown limited electrochemical performance. For example, though pure Sn delivered a capacity of ca. 900 mA h g−1 during the 1st magnesiation (complete conversion of Sn to Mg2Sn), the capacity abruptly dropped to ca. 200–300 mA h g−1 after the 1st magnesiation, and continuously declined with repeated charge/discharge (C/D). Furthermore, pure Sn often requires the application of potentials close to, or even lower than, the thermodynamic Mg plating potential for the activation of an alloying process.

What's the electrolyte?

Mg(N(SO2CF3)2)2/diglyme

The plating and stripping behaviors of the LiF@Po–Li electrodes are observed by SEM after the Li deposition of 1 mA h cm−2 at a current density of 0.5 mA cm−2 with an EC/DEC:3/7, 1.3 M LiPF6, and 5 wt% FEC electrolyte (Fig. 3). All the electrochemical experiments are performed with carbonate-based electrolytes. It has been well known that the ether-based electrolytes, especially the DME/DOL system, exhibit better electrochemical properties compared to carbonate-based electrolytes in LMBs thanks to their stable SEI layer formation and more freedom to choose additives. However, the ether-based electrolyte is not suitable for various commercial cathode material due to its instability above 4.0 V vs. Li/Li+. In this regard, carbonated-based electrolytes should be explored for their electrochemical properties in upcoming high energy density LMBs. The electrodes are washed with dimethyl carbonate (DMC) to remove residual salts before SEM observation. The Li metal plated on the bare Li electrode shows a mossy, porous and dendritic morphology with an 11.8 μm thickness and 59% porosity (Fig. 3a, b and ESI Note 1†). In contrast, the Li metal plated on the LiF@Po–Li electrode exhibits a smooth and film-like morphology with significantly reduced porosity (Fig. 3c, d and S7†). The cross-sectional view of the LiF@Po–Li electrode after plating clearly shows that a 6.8 μm-thick dense Li film is deposited on the LiF@Po–Li electrode with 28% porosity (Fig. 3d). The stripping behavior is also monitored. As shown in Fig. 3e and f, the bare Li electrode has high-density craters with sizes ranging from 10 μm to 50 μm after the Li stripping. This non-uniform surface morphology is attributed to the localized current density caused by the inhomogeneous SEI layer. In contrast, the LiF@Po–Li electrode has a smooth and conformal surface without noticeable holes over a large area after Li stripping (Fig. 3g and h). This improvement in plating and stripping behaviors of Li is attributed to the LiF@Po protective layer, which enables a uniform Li ion concentration and current density over the Li electrode by suppressing the electrolyte decomposition.

What's the cathode?

The plating and stripping behaviors of the LiF@Po–Li electrodes are observed by SEM after the Li deposition of 1 mA h cm−2 at a current density of 0.5 mA cm−2 with an EC/DEC:3/7, 1.3 M LiPF6, and 5 wt% FEC electrolyte (Fig. 3). All the electrochemical experiments are performed with carbonate-based electrolytes. It has been well known that the ether-based electrolytes, especially the DME/DOL system, exhibit better electrochemical properties compared to carbonate-based electrolytes in LMBs thanks to their stable SEI layer formation and more freedom to choose additives. However, the ether-based electrolyte is not suitable for various commercial cathode material due to its instability above 4.0 V vs. Li/Li+. In this regard, carbonated-based electrolytes should be explored for their electrochemical properties in upcoming high energy density LMBs. The electrodes are washed with dimethyl carbonate (DMC) to remove residual salts before SEM observation. The Li metal plated on the bare Li electrode shows a mossy, porous and dendritic morphology with an 11.8 μm thickness and 59% porosity (Fig. 3a, b and ESI Note 1†). In contrast, the Li metal plated on the LiF@Po–Li electrode exhibits a smooth and film-like morphology with significantly reduced porosity (Fig. 3c, d and S7†). The cross-sectional view of the LiF@Po–Li electrode after plating clearly shows that a 6.8 μm-thick dense Li film is deposited on the LiF@Po–Li electrode with 28% porosity (Fig. 3d). The stripping behavior is also monitored. As shown in Fig. 3e and f, the bare Li electrode has high-density craters with sizes ranging from 10 μm to 50 μm after the Li stripping. This non-uniform surface morphology is attributed to the localized current density caused by the inhomogeneous SEI layer. In contrast, the LiF@Po–Li electrode has a smooth and conformal surface without noticeable holes over a large area after Li stripping (Fig. 3g and h). This improvement in plating and stripping behaviors of Li is attributed to the LiF@Po protective layer, which enables a uniform Li ion concentration and current density over the Li electrode by suppressing the electrolyte decomposition.

What's the electrolyte?

5 wt% FEC

The plating and stripping behaviors of the LiF@Po–Li electrodes are observed by SEM after the Li deposition of 1 mA h cm−2 at a current density of 0.5 mA cm−2 with an EC/DEC:3/7, 1.3 M LiPF6, and 5 wt% FEC electrolyte (Fig. 3). All the electrochemical experiments are performed with carbonate-based electrolytes. It has been well known that the ether-based electrolytes, especially the DME/DOL system, exhibit better electrochemical properties compared to carbonate-based electrolytes in LMBs thanks to their stable SEI layer formation and more freedom to choose additives. However, the ether-based electrolyte is not suitable for various commercial cathode material due to its instability above 4.0 V vs. Li/Li+. In this regard, carbonated-based electrolytes should be explored for their electrochemical properties in upcoming high energy density LMBs. The electrodes are washed with dimethyl carbonate (DMC) to remove residual salts before SEM observation. The Li metal plated on the bare Li electrode shows a mossy, porous and dendritic morphology with an 11.8 μm thickness and 59% porosity (Fig. 3a, b and ESI Note 1†). In contrast, the Li metal plated on the LiF@Po–Li electrode exhibits a smooth and film-like morphology with significantly reduced porosity (Fig. 3c, d and S7†). The cross-sectional view of the LiF@Po–Li electrode after plating clearly shows that a 6.8 μm-thick dense Li film is deposited on the LiF@Po–Li electrode with 28% porosity (Fig. 3d). The stripping behavior is also monitored. As shown in Fig. 3e and f, the bare Li electrode has high-density craters with sizes ranging from 10 μm to 50 μm after the Li stripping. This non-uniform surface morphology is attributed to the localized current density caused by the inhomogeneous SEI layer. In contrast, the LiF@Po–Li electrode has a smooth and conformal surface without noticeable holes over a large area after Li stripping (Fig. 3g and h). This improvement in plating and stripping behaviors of Li is attributed to the LiF@Po protective layer, which enables a uniform Li ion concentration and current density over the Li electrode by suppressing the electrolyte decomposition.

What's the electrolyte?

carbonate-based

Supercapacitors are a class of energy storage devices found commonly in hybrid electric vehicles, camera components and used as backup power systems. One class of supercapacitors, called electric double layer capacitors (EDLCs), are a popular option for commercial applications. In EDLCs, electrolyte ions are adsorbed onto the electrode/electrolyte interface of high-surface area carbon and charge is stored solely via this electrostatic double layer. Due to this charge storage process, the structural integrity of the electrode is maintained resulting in high power densities and cycle lifetimes of over one million cycles. One major drawback of these devices is their low energy density, which limits their use in many applications. To improve the energy density of supercapacitors, many transition metal oxides such as V2O5, Ni(OH)2, Co3O4, RuO2, MnO2 and WO3 have been investigated. Unlike carbon-based materials, transition metal oxides undergo faradaic reactions that facilitate charge storage. Due to accessible redox states, transition metal oxides have much higher energy densities, but most suffer from low conductivity and poor cycling stability. In particular, tungsten oxide has been suggested as a potential alternative for supercapacitor electrode materials due to its redox active properties. Recent efforts have focused on improving tungsten oxide containing supercapacitors including their electrochemical performance, rate, and cycling stability in order to improve this material for energy storage applications.

What's the electrolyte?

For electrochemical activation of the CRLE two electrodes were placed in a beaker in a distance of approximately 10 cm. A reference electrode (Ag/AgCl 3 M KCl) was placed in close proximity to the electrode to be activated. The second CRLE served as inert counter electrode. In screening experiments different activation potentials and polarisation time were tested at room temperature using 0.1 M NaOH as electrolyte (the results are shown in ESI-7†). Polarisation at 1800 mV for 90 s yielded the best electrode performance at short activation time (Table 1).

What's the electrolyte?

0.1 M NaOH

The listed physical parameters are essential for designing a new generation of electronic devices with significantly improved sensitivity and time response. For instance, the high piezoelectric activity of near room temperature ferroelectrics Sn2P2S6, whose Tc ≈ 337 K, can compete with well-known BiFeO3, currently one of the top choices for switchable ferroelectric diodes. Near room temperature spontaneous polarization of Sn2P2S6 can serve as a base for non-volatile ferroelectric random-access memory (FERAM) elements, at the same time significantly miniaturizing the designed systems. The first steps of such practical exploits have been recently observed for layered room-temperature ferroelectric CuInP2S6, demonstrating that reversible polarization switchability with an on/off ratio of about ∼100 can be achieved for ferroelectric flakes with the thickness reaching an impressive limit of ∼4 nm. Due to locally controlled ion conductivity and giant negative electrostrictive coefficients, CuInP2S6 is also considered as a solid electrolyte for boosting sodium and lithium battery performance. Similarly, 2D superparaelectric Sn2P2S6 with a large dielectric constant was also proposed for the same purpose. The high photoelectric conversion efficiency of Sn2P2S6 allowed exploration of the ultrafast shift-current behaviour, comparable only to the performance of SbSI, CdSe, and CdS. These properties can be utilised for the novel design of the band edge shift current based photovoltaic devices, pushing the solar cell energy harvesting towards their theoretical limit. Additionally, Sn2P2S6 ferroelectrics have prominent photorefractive properties in the near-infrared spectral region, which makes them suitable for highly efficient acousto-optic and holographic devices.

What's the electrolyte?

CuInP2S6

7Li magic-angle spinning (MAS) ssNMR spectroscopy was performed to trace the local environment of Li. Fig. 2d and S4a† shows the 7Li projection magic-angle turning and phase-adjusted sideband separation (pjMATPASS) spectra for the pristine, two-cycled and treated LMO. The pristine LMO showed two isotropic peaks resulting from Li in the octahedral site (LiO6). The main peak with a shoulder at around 748 ppm comes from Li in the Li layer and the one at around 1500 ppm can be attributed to Li in the TM layer. For pristine LMO, the small peak at around 0 ppm was due to the residual Li2CO3. The main peaks at around 748 ppm and 1500 ppm remained in the two-cycled LMO samples (Fig. S4a†). The sharp signal with enhanced intensity at around 0 ppm was mainly due to the solid electrolyte interphase (SEI) components, such as Li2CO3 and LiF. The treated T-LMO sample exhibited that some Li-ions remained in the TM layer, as for LMO and two-cycled LMO. The remaining sidebands at around ±250 ppm in T-LMO came from a peak at around 0 ppm. Besides the similar peaks in LMO, a new peak appeared at around 529 ppm in T-LMO, which can be ascribed to Li in the tetrahedral site, like that in LiMn2O4. A broad peak at around 660 ppm corresponds well with that of the Li2Mn4O9 reference with Li in the tetrahedral site (Fig. S4b†). Although a small amount of disordered shells may exist in the T-LMO sample, their intensities are too low to be detected here, which might be overlapped by the strong NMR signal associated with Li contained in the crystalline core of the particles. The T-LMO curve was deconvoluted by the standard reference spectrum (black curve), which showed that T-LMO consists of 35.5% LMO, 16.9% LiMn2O4 and 47.6% Li2Mn4O9. The NMR spectra result provides direct evidence that new spinel phases appear in the T-LMO, which may be due to cation rearrangement.

What's the electrolyte?

For the half-cell tests, the electrochemical properties were evaluated using a CR2032 corn-type cell. The cells were prepared in an Ar-filled glove box with H2O and O2 concentrations <0.1 ppm. Si@CTSC, Si@ALGC, Si@PVDFC and Si/PVDF were used as the anodes (diameter: 14 mm), and Li metal foil was used as the counter electrodes (thickness: 300 μm and diameter: 12 mm). A Celgard 2400 polypropylene membrane was used as the separator. The electrolyte was 1 M LiPF6 in a mixture of ethylene carbonate (EC), dimethyl carbonate (DMC), diethyl carbonate (DEC) (1:1:1 by vol) with 5 wt% fluoroethylene carbonate (FEC). The electrolyte amount was 5 g A h−1. The electrochemical tests were carried out after letting the assembled cells stand for 24 hours to ensure full infiltration of the electrolyte. The GCD measurements were carried on a Land CT2001A system from 0.01 to 2.5 V. CV measurements were performed with sweep rates from 0.2 to 1.2 mV s−1 using a CHI660E work station. The frequency for the EIS measurements ranged from 100 kHz to 10 mHz. GITT tests were carried out using a constant current pulse with a duration of 10 min and a relaxation process over 10 min.

What's the anode?

Si@CTSC, Si@ALGC, Si@PVDFC and Si/PVDF

For the half-cell tests, the electrochemical properties were evaluated using a CR2032 corn-type cell. The cells were prepared in an Ar-filled glove box with H2O and O2 concentrations <0.1 ppm. Si@CTSC, Si@ALGC, Si@PVDFC and Si/PVDF were used as the anodes (diameter: 14 mm), and Li metal foil was used as the counter electrodes (thickness: 300 μm and diameter: 12 mm). A Celgard 2400 polypropylene membrane was used as the separator. The electrolyte was 1 M LiPF6 in a mixture of ethylene carbonate (EC), dimethyl carbonate (DMC), diethyl carbonate (DEC) (1:1:1 by vol) with 5 wt% fluoroethylene carbonate (FEC). The electrolyte amount was 5 g A h−1. The electrochemical tests were carried out after letting the assembled cells stand for 24 hours to ensure full infiltration of the electrolyte. The GCD measurements were carried on a Land CT2001A system from 0.01 to 2.5 V. CV measurements were performed with sweep rates from 0.2 to 1.2 mV s−1 using a CHI660E work station. The frequency for the EIS measurements ranged from 100 kHz to 10 mHz. GITT tests were carried out using a constant current pulse with a duration of 10 min and a relaxation process over 10 min.

What's the electrolyte?

1 M LiPF6

Ti3C2Tx-based materials have been widely reported as anode materials for supercapacitors and hybrid supercapacitors with the acid electrolyte. Here, the modified Ti3C2Tx MXene (400-KOH-Ti3C2) samples were synthesized based on a previous report. The XRD pattern (Fig. 1a and S1†) reveals that the diffraction peaks of 400-KOH-Ti3C2 shift to a lower 2θ angle compared with those of the original Ti3C2Tx, indicating the increase of interlayer distances after KOH treatment and calcination. In addition, the intensity of partial diffraction peaks for 400-KOH-Ti3C2 is slightly weaker than that of original Ti3C2Tx, suggesting that some feature of the Ti3C2Tx diffraction peaks is shielded by the K+ intercalation into the Ti3C2 MXene layer. Raman spectroscopy (Fig. 1b) and XPS tests (Fig. S2†) were carried out to further manifest the purity and chemical composition of the as-obtained 400-KOH-Ti3C2 samples. To clearly reflect the interlayer spacing for 400-KOH-Ti3C2 samples, the SEM and TEM images are given in Fig. S3† and 1c, respectively. It can be found that 400-KOH-Ti3C2 consisted of stacked MXene nanosheets. The corresponding HRTEM image shows that the interlayer distance of 400-KOH-Ti3C2 is 1.17 nm (Fig. 1d), much larger than that of the original Ti3C2Tx (0.98 nm, Fig. S4†), further confirming the expansion of the interlayer distance after KOH treatment and calcination. In addition, the scanning electron microscopy-energy dispersive X-ray spectroscopy (SEM-EDX) mapping images of 400-KOH-Ti3C2 clearly indicate that Ti, C, O, F and K elements are homogeneously distributed in all MXene sheets (Fig. S5†).

What's the anode?

Ti3C2Tx-based materials

The M2+–Fe LDHs (M = Ni, Zn) nanosheets were prepared using a facile electrosynthesis method. Typically, nickel foam (NF, 10 × 30 × 0.05 mm3) was sonicated in 2 M HCl solution for 15 min and subsequently rinsed with water and ethanol to ensure a clean surface. The electrodeposition was carried out in a standard three-electrode system cell containing NF as the working electrode, Pt plate as the counter electrode and Ag/AgCl (3 M KCl) as the reference electrode. The electrolyte for the electrosynthesis of NiFe LDH contained Ni(NO3)2·6H2O (0.03 M) and Fe(NO3)3·9H2O (0.01 M). The total cation concentration in the electrolyte was maintained at 0.04 M. The potentiostatic deposition was then carried out at −1.0 V vs. Ag/AgCl for 300 s at room temperature. The obtained NiFe LDH nanosheets were carefully withdrawn and rinsed thoroughly with water and ethanol and left to dry at 45 °C. The NiZnFe LDHs nanosheets were prepared via a similar method, by replacing Ni(NO3)2·6H2O (0.03 M) with a mixture of Ni(NO3)2·6H2O (0.027 M) and Zn(NO3)2·6H2O (0.003 M).

What's the electrolyte?

Ni(NO3)2·6H2O (0.03 M) and Fe(NO3)3·9H2O (0.01 M)

EELS spectra are obtained from the thin edge of cycled LLOs as labelled by the orange lines in Fig. 4(a and d) to study oxygen vacancies and TM valence states. Detailed comparisons of O-K, Ni-L, Co-L and Mn-L spectra are presented in Fig. 4(g–j) and S7.† For O-K edges, the intensity of the prepeak (∼530 eV) is sensitive to the valence states of TM ions. The intensity of this prepeak decreases drastically from the interior of the particle to the surface and almost disappears at the surface of the cycled LLOs at the cut-off of 4.8 V (Fig. S7†), indicating a significant decrease of TM valence states. This matches well with the fine changes of Mn-L edges in Fig. 4(h), which show a chemical shift (∼2.5 eV) toward the lower binding energy on the surface (0–10 nm depth), as compared with the cycled bulk (≥30 nm depth). These signatures demonstrate that the majority of Mn ions start to shift to a lower valence at a depth of 20 nm beneath the surface of cycled LLOs at the cut-off of 4.8 V, while for the LLOs cycled at 4.5 V, the reduced valence of dominant Mn ions occurs in the top 10 nm layer beneath the surface (Fig. 4(j)). The valence decrease of Mn ions in LLOs is usually accompanied by Mn dissolution into the electrolyte during cycling, which is measured by ICP-OES as shown in Fig. S8.† The Mn concentration dissolved in the electrolyte at the cut-off of 4.8 V after 100 and 200 cycles reaches 139 and 267 ppm, respectively, which are much higher than the 56 and 97 ppm at 4.5 V. These results confirm the reduced Mn dissolution into the electrolyte for the LLOs cycled at the low cut-off of 4.5 V.

What's the electrolyte?

For the first cycle in LP30 a more apparent “pitting” peak is observed that occurs at an earlier time compared to LP30 + FEC (occurring at ∼78% and ∼92% capacity, respectively for 0.5 mA cm−2). Other studies have suggested that this is due to inhomogeneous dissolution of the lithium whiskers that result in dead Li formation and early peaking behaviour. However, the lower plating efficiency quantified with in situ NMR can also lead to the early peaking behaviour observed when lower amounts of microstructures are present. With 1 and 2 mA cm−2, the peaking in the first cycle (where the stripping current is kept at 1 mA cm−2) occurs at 85% and 89% capacity respectively. This correlates well with the in situ NMR, which indicated higher plating efficiencies for the higher current densities in LP30. The voltage traces are flatter for the LP30 + FEC electrolyte, as compared to those for LP30, consistent with both the higher plating efficiency seen in the in situ NMR and of studies showing minuscule dead lithium formation in LP30 + FEC. The lower overpotential observed for LP30 + FEC is somewhat consistent between cells (Fig. S2–S4†), but the overpotential is affected both by the resistances in the cell (in particular of the SEI) and the surface area (which increases during electrodeposition) accounting for variations between cells.

What's the electrolyte?

LP30 + FEC

(2) A Ti metal precursor is anodized in an alkali electrolyte (e.g. NaOH and Ba(OH)2 solutions). The alkaline anodization (Fig. 6B) could be considered an electro-assisted hydrothermal-like process with 2 stages where TiO2·2H2O anodized from the titanium anode reacts with the hot alkaline surroundings generated by a continuous electric field. In the first stage, a fast anodic reaction on the titanium surface creates a thin TiO2·2H2O nanosheet layer (eqn (1) and (2)) which would react with alkali to form titanate nanowire bundles on the titanium anode (eqn (3)). In particular, it should be mentioned that the OH− concentration should be above 4 M to ensure the formation of titanate rather than titania. In the second stage, the anodized TiO2·2H2O would also break down and disperse into the electrolyte to form a 3D hierarchical porous structure with a large specific surface area.

What's the anode?

titanium

(2) A Ti metal precursor is anodized in an alkali electrolyte (e.g. NaOH and Ba(OH)2 solutions). The alkaline anodization (Fig. 6B) could be considered an electro-assisted hydrothermal-like process with 2 stages where TiO2·2H2O anodized from the titanium anode reacts with the hot alkaline surroundings generated by a continuous electric field. In the first stage, a fast anodic reaction on the titanium surface creates a thin TiO2·2H2O nanosheet layer (eqn (1) and (2)) which would react with alkali to form titanate nanowire bundles on the titanium anode (eqn (3)). In particular, it should be mentioned that the OH− concentration should be above 4 M to ensure the formation of titanate rather than titania. In the second stage, the anodized TiO2·2H2O would also break down and disperse into the electrolyte to form a 3D hierarchical porous structure with a large specific surface area.

What's the electrolyte?

alkali

(2) A Ti metal precursor is anodized in an alkali electrolyte (e.g. NaOH and Ba(OH)2 solutions). The alkaline anodization (Fig. 6B) could be considered an electro-assisted hydrothermal-like process with 2 stages where TiO2·2H2O anodized from the titanium anode reacts with the hot alkaline surroundings generated by a continuous electric field. In the first stage, a fast anodic reaction on the titanium surface creates a thin TiO2·2H2O nanosheet layer (eqn (1) and (2)) which would react with alkali to form titanate nanowire bundles on the titanium anode (eqn (3)). In particular, it should be mentioned that the OH− concentration should be above 4 M to ensure the formation of titanate rather than titania. In the second stage, the anodized TiO2·2H2O would also break down and disperse into the electrolyte to form a 3D hierarchical porous structure with a large specific surface area.

What's the anode?

titanium

(2) A Ti metal precursor is anodized in an alkali electrolyte (e.g. NaOH and Ba(OH)2 solutions). The alkaline anodization (Fig. 6B) could be considered an electro-assisted hydrothermal-like process with 2 stages where TiO2·2H2O anodized from the titanium anode reacts with the hot alkaline surroundings generated by a continuous electric field. In the first stage, a fast anodic reaction on the titanium surface creates a thin TiO2·2H2O nanosheet layer (eqn (1) and (2)) which would react with alkali to form titanate nanowire bundles on the titanium anode (eqn (3)). In particular, it should be mentioned that the OH− concentration should be above 4 M to ensure the formation of titanate rather than titania. In the second stage, the anodized TiO2·2H2O would also break down and disperse into the electrolyte to form a 3D hierarchical porous structure with a large specific surface area.

What's the electrolyte?

(e.g. NaOH and Ba(OH)2 solutions)

According to the simulation, the number of moles per surface area formed in the two electrolytes, NSEI (Fig. S17a†) is also greater for the LP30 + FEC electrolyte, indicating a thicker SEI is being formed. We have estimated the thickness of the SEI by using eqn (13) and assuming it to be pure Li2CO3 so as to provide a qualitative understanding of the extent of SEI formation. Averaging over the whole experiment (74 hours) the SEI formation rate is 6.1 nm h−1 in LP30 and 12 nm h−1 for LP30 + FEC. The values are relatively large and seem to overestimate the thickness of the SEI but are on a similar scale to what was estimated in an earlier isotope exchange study (14 nm h−1 in LP30). One possible reason for this overestimation of thickness is the assumption that all of the reduced electrolyte species are deposited to form the SEI. However, it has been shown experimentally that a wide range of the reduced electrolyte species are soluble and go into the electrolyte. We also note that the comparison of SEI thicknesses needs to be interpreted with caution as the chemical composition and density is expected to differ between the two electrolytes.

What's the electrolyte?

LP30 + FEC

Next, the annealed WO3−x film was assembled into an electrochromic device to test its functionality. The device structure is shown in Fig. 3a, in which the FTO glass is used as the transparent electrode. The electrolyte with 1.25 M LiClO4 in a mixed solvent was used in the electrochromic device to achieve the insertion and extraction of Li+ ions in the WO3−x film. The device in the bleached state shows a high transmittance in both the visible and infrared wavelength range (Fig. 3b), while in the colored state, the transmittance of the electrochromic device decreases significantly. At the wavelength of 680 nm, the transmittance difference between the two curves is around 70%. In the wavelength range of 700–1000 nm, the transmittance of the colored state is even less than 3%, indicating the excellent near infrared light shielding ability. Hence, in the scorching summer, smart windows made of this electrochromic WO3−x film can keep the room cool by blocking heat radiation. Fig. 3c shows the photographs of the packaged electrochromic device in the coloring and bleaching process. A series of different color states from light to deep blue can be easily obtained by adjusting the applied bias from −3.5 V to 3.5 V, which may have potential applications in the field of electrochromic displays.

What's the electrolyte?

1.25 M LiClO4

Intermetallic Mg2Sn alloyed with extra Mg was presented as a new high-performance anode for MIBs. The 3Mg/Mg2Sn was composed of c-Mg, a-Mg, and Mg2Sn, and showed excellent electrochemical performance when cycled in Mg(HMDS)2:MgCl2/THF. During the 1st de-magnesiation, 3Mg/Mg2Sn first dissolved Mg2+ from c-Mg, which was then followed by the complete conversion of Mg2Sn to Sn with a partial release of a-Mg. The irreversible dissolution of c-Mg facilitated the reversible de-magnesiation/magnesiation in Mg2Sn and a-Mg, and allowed 3Mg/Mg2Sn to demonstrate unprecedented electrochemical properties. Thanks to the high surface area and pore volume ascribed to the irreversible dissolution of c-Mg, the 3Mg/Mg2Sn anode delivered reversible capacities of 805 and 430 mA h g−1 at 100 and 1500 mA g−1, respectively, with reasonable cycling stability. De-magnesiated 3Mg/Mg2Sn also showed similar electrochemical performance in other conventional electrolyte systems (Mg(TFSI)2:MgCl2/diglyme and Mg(TFSI)2/acetonitrile). In addition to the beneficial role of c-Mg with respect to the electrochemical properties of 3Mg/Mg2Sn in a half-cell configuration, irreversibly dissolved Mg2+ ions could ideally balance the Mg2+ ions trapped in a Mo6S8 cathode in a full-cell.

What's the cathode?

Mo6S8

Intermetallic Mg2Sn alloyed with extra Mg was presented as a new high-performance anode for MIBs. The 3Mg/Mg2Sn was composed of c-Mg, a-Mg, and Mg2Sn, and showed excellent electrochemical performance when cycled in Mg(HMDS)2:MgCl2/THF. During the 1st de-magnesiation, 3Mg/Mg2Sn first dissolved Mg2+ from c-Mg, which was then followed by the complete conversion of Mg2Sn to Sn with a partial release of a-Mg. The irreversible dissolution of c-Mg facilitated the reversible de-magnesiation/magnesiation in Mg2Sn and a-Mg, and allowed 3Mg/Mg2Sn to demonstrate unprecedented electrochemical properties. Thanks to the high surface area and pore volume ascribed to the irreversible dissolution of c-Mg, the 3Mg/Mg2Sn anode delivered reversible capacities of 805 and 430 mA h g−1 at 100 and 1500 mA g−1, respectively, with reasonable cycling stability. De-magnesiated 3Mg/Mg2Sn also showed similar electrochemical performance in other conventional electrolyte systems (Mg(TFSI)2:MgCl2/diglyme and Mg(TFSI)2/acetonitrile). In addition to the beneficial role of c-Mg with respect to the electrochemical properties of 3Mg/Mg2Sn in a half-cell configuration, irreversibly dissolved Mg2+ ions could ideally balance the Mg2+ ions trapped in a Mo6S8 cathode in a full-cell.

What's the anode?

3Mg/Mg2Sn

Intermetallic Mg2Sn alloyed with extra Mg was presented as a new high-performance anode for MIBs. The 3Mg/Mg2Sn was composed of c-Mg, a-Mg, and Mg2Sn, and showed excellent electrochemical performance when cycled in Mg(HMDS)2:MgCl2/THF. During the 1st de-magnesiation, 3Mg/Mg2Sn first dissolved Mg2+ from c-Mg, which was then followed by the complete conversion of Mg2Sn to Sn with a partial release of a-Mg. The irreversible dissolution of c-Mg facilitated the reversible de-magnesiation/magnesiation in Mg2Sn and a-Mg, and allowed 3Mg/Mg2Sn to demonstrate unprecedented electrochemical properties. Thanks to the high surface area and pore volume ascribed to the irreversible dissolution of c-Mg, the 3Mg/Mg2Sn anode delivered reversible capacities of 805 and 430 mA h g−1 at 100 and 1500 mA g−1, respectively, with reasonable cycling stability. De-magnesiated 3Mg/Mg2Sn also showed similar electrochemical performance in other conventional electrolyte systems (Mg(TFSI)2:MgCl2/diglyme and Mg(TFSI)2/acetonitrile). In addition to the beneficial role of c-Mg with respect to the electrochemical properties of 3Mg/Mg2Sn in a half-cell configuration, irreversibly dissolved Mg2+ ions could ideally balance the Mg2+ ions trapped in a Mo6S8 cathode in a full-cell.

What's the electrolyte?

Mg(TFSI)2:MgCl2/diglyme and Mg(TFSI)2/acetonitrile

A button-sized hybrid electrolyte cell (Fig. S4†), in which an ionic liquid-infiltrated NCM with a LLZTO electrolyte and a 20 μm–thick Li metal were respectively used as a cathode and a Li metal anode, was fabricated to check out the improvement of the electrochemical performance after laser-annealing treatment. For comparison, two kinds of LLZTO pellets were used: a polished pellet without laser treatment and a laser-treated pellet. In Fig. 6a and b, the Nyquist plots of the AC-impedance spectra from the samples obtained at an open circuit from the cells were compared. As shown in Fig. 6a and b, the ohmic losses calculated from the high-frequency intercept with the real axis were almost the same in the range of 3.2–4.0 Ω cm2 for both the cells. Based on the ohmic resistance, the ionic conductivities were expected to be in the range of 7.5–9.4 10−3 S cm−1 at 60 °C for the LLZTO pellets with and without laser treatment. These conductivity values are consistent with the reported values of Ta-doped LLZO and also with that value (∼8.6 × 10−3 S cm−1 at 60 °C) experimentally obtained from the Au-sputtered symmetric cells (Au/LLZTO/Au), as shown in Fig. S5a.† Thus, it can be inferred that the cathode and ionic liquid catholyte are not contributing significantly to the ohmic losses in these cells. Furthermore, the amorphous layer formed by laser irradiation had no significant effect on the ohmic resistance due to its low thickness of about 400 nm. The AC-impedance spectra of both cells consist of a depressed arc in the high-frequency range, a narrow line inclined at a constant angle to the real axis (Warburg impedance), and a capacitive line (blocking region) in the low-frequency range, which are typical features of a NCM-Li cell. The high-frequency arc could be associated with the charge-transfer reactions at the cathode and anode interface. Because the relaxation times for the cathodic and anodic reactions are generally overlapped over the high-frequency ranges for a NCM-Li full cell, it is difficult to separate the charge-transfer resistance at the Li-LLZTO interface from the resistance at the NCM-ionic liquid. The AC-impedance spectrum for Li/LLZTO/Li cell is also shown in Fig. S5b.† Considering that the electrochemical kinetics at the cathode would be almost the same for both cells, the variation in the high-frequency arc could be attributed to the charge-transfer resistance at the Li metal and LLZTO interface. The values of the overall interfacial resistance, Rct, were quantitatively determined from complex non-linear least-square (CNLS) fitting of the measured impedance spectra based on a typical equivalent circuit of the NCM-Li full cell. The Rct (Rc + Ra) value was reduced from ∼80 Ω cm2 to ∼26 Ω cm2 when the LLZTO surface was treated by a laser beam. In addition, as shown in Fig. S5c,† the electronic conductivity of LLZTO was significantly reduced by the laser treatment, which can originate from the formation of a wide band gap surface layer.

What's the cathode?

NCM

A button-sized hybrid electrolyte cell (Fig. S4†), in which an ionic liquid-infiltrated NCM with a LLZTO electrolyte and a 20 μm–thick Li metal were respectively used as a cathode and a Li metal anode, was fabricated to check out the improvement of the electrochemical performance after laser-annealing treatment. For comparison, two kinds of LLZTO pellets were used: a polished pellet without laser treatment and a laser-treated pellet. In Fig. 6a and b, the Nyquist plots of the AC-impedance spectra from the samples obtained at an open circuit from the cells were compared. As shown in Fig. 6a and b, the ohmic losses calculated from the high-frequency intercept with the real axis were almost the same in the range of 3.2–4.0 Ω cm2 for both the cells. Based on the ohmic resistance, the ionic conductivities were expected to be in the range of 7.5–9.4 10−3 S cm−1 at 60 °C for the LLZTO pellets with and without laser treatment. These conductivity values are consistent with the reported values of Ta-doped LLZO and also with that value (∼8.6 × 10−3 S cm−1 at 60 °C) experimentally obtained from the Au-sputtered symmetric cells (Au/LLZTO/Au), as shown in Fig. S5a.† Thus, it can be inferred that the cathode and ionic liquid catholyte are not contributing significantly to the ohmic losses in these cells. Furthermore, the amorphous layer formed by laser irradiation had no significant effect on the ohmic resistance due to its low thickness of about 400 nm. The AC-impedance spectra of both cells consist of a depressed arc in the high-frequency range, a narrow line inclined at a constant angle to the real axis (Warburg impedance), and a capacitive line (blocking region) in the low-frequency range, which are typical features of a NCM-Li cell. The high-frequency arc could be associated with the charge-transfer reactions at the cathode and anode interface. Because the relaxation times for the cathodic and anodic reactions are generally overlapped over the high-frequency ranges for a NCM-Li full cell, it is difficult to separate the charge-transfer resistance at the Li-LLZTO interface from the resistance at the NCM-ionic liquid. The AC-impedance spectrum for Li/LLZTO/Li cell is also shown in Fig. S5b.† Considering that the electrochemical kinetics at the cathode would be almost the same for both cells, the variation in the high-frequency arc could be attributed to the charge-transfer resistance at the Li metal and LLZTO interface. The values of the overall interfacial resistance, Rct, were quantitatively determined from complex non-linear least-square (CNLS) fitting of the measured impedance spectra based on a typical equivalent circuit of the NCM-Li full cell. The Rct (Rc + Ra) value was reduced from ∼80 Ω cm2 to ∼26 Ω cm2 when the LLZTO surface was treated by a laser beam. In addition, as shown in Fig. S5c,† the electronic conductivity of LLZTO was significantly reduced by the laser treatment, which can originate from the formation of a wide band gap surface layer.

What's the anode?

Li metal

A button-sized hybrid electrolyte cell (Fig. S4†), in which an ionic liquid-infiltrated NCM with a LLZTO electrolyte and a 20 μm–thick Li metal were respectively used as a cathode and a Li metal anode, was fabricated to check out the improvement of the electrochemical performance after laser-annealing treatment. For comparison, two kinds of LLZTO pellets were used: a polished pellet without laser treatment and a laser-treated pellet. In Fig. 6a and b, the Nyquist plots of the AC-impedance spectra from the samples obtained at an open circuit from the cells were compared. As shown in Fig. 6a and b, the ohmic losses calculated from the high-frequency intercept with the real axis were almost the same in the range of 3.2–4.0 Ω cm2 for both the cells. Based on the ohmic resistance, the ionic conductivities were expected to be in the range of 7.5–9.4 10−3 S cm−1 at 60 °C for the LLZTO pellets with and without laser treatment. These conductivity values are consistent with the reported values of Ta-doped LLZO and also with that value (∼8.6 × 10−3 S cm−1 at 60 °C) experimentally obtained from the Au-sputtered symmetric cells (Au/LLZTO/Au), as shown in Fig. S5a.† Thus, it can be inferred that the cathode and ionic liquid catholyte are not contributing significantly to the ohmic losses in these cells. Furthermore, the amorphous layer formed by laser irradiation had no significant effect on the ohmic resistance due to its low thickness of about 400 nm. The AC-impedance spectra of both cells consist of a depressed arc in the high-frequency range, a narrow line inclined at a constant angle to the real axis (Warburg impedance), and a capacitive line (blocking region) in the low-frequency range, which are typical features of a NCM-Li cell. The high-frequency arc could be associated with the charge-transfer reactions at the cathode and anode interface. Because the relaxation times for the cathodic and anodic reactions are generally overlapped over the high-frequency ranges for a NCM-Li full cell, it is difficult to separate the charge-transfer resistance at the Li-LLZTO interface from the resistance at the NCM-ionic liquid. The AC-impedance spectrum for Li/LLZTO/Li cell is also shown in Fig. S5b.† Considering that the electrochemical kinetics at the cathode would be almost the same for both cells, the variation in the high-frequency arc could be attributed to the charge-transfer resistance at the Li metal and LLZTO interface. The values of the overall interfacial resistance, Rct, were quantitatively determined from complex non-linear least-square (CNLS) fitting of the measured impedance spectra based on a typical equivalent circuit of the NCM-Li full cell. The Rct (Rc + Ra) value was reduced from ∼80 Ω cm2 to ∼26 Ω cm2 when the LLZTO surface was treated by a laser beam. In addition, as shown in Fig. S5c,† the electronic conductivity of LLZTO was significantly reduced by the laser treatment, which can originate from the formation of a wide band gap surface layer.

What's the electrolyte?

LLZTO

A button-sized hybrid electrolyte cell (Fig. S4†), in which an ionic liquid-infiltrated NCM with a LLZTO electrolyte and a 20 μm–thick Li metal were respectively used as a cathode and a Li metal anode, was fabricated to check out the improvement of the electrochemical performance after laser-annealing treatment. For comparison, two kinds of LLZTO pellets were used: a polished pellet without laser treatment and a laser-treated pellet. In Fig. 6a and b, the Nyquist plots of the AC-impedance spectra from the samples obtained at an open circuit from the cells were compared. As shown in Fig. 6a and b, the ohmic losses calculated from the high-frequency intercept with the real axis were almost the same in the range of 3.2–4.0 Ω cm2 for both the cells. Based on the ohmic resistance, the ionic conductivities were expected to be in the range of 7.5–9.4 10−3 S cm−1 at 60 °C for the LLZTO pellets with and without laser treatment. These conductivity values are consistent with the reported values of Ta-doped LLZO and also with that value (∼8.6 × 10−3 S cm−1 at 60 °C) experimentally obtained from the Au-sputtered symmetric cells (Au/LLZTO/Au), as shown in Fig. S5a.† Thus, it can be inferred that the cathode and ionic liquid catholyte are not contributing significantly to the ohmic losses in these cells. Furthermore, the amorphous layer formed by laser irradiation had no significant effect on the ohmic resistance due to its low thickness of about 400 nm. The AC-impedance spectra of both cells consist of a depressed arc in the high-frequency range, a narrow line inclined at a constant angle to the real axis (Warburg impedance), and a capacitive line (blocking region) in the low-frequency range, which are typical features of a NCM-Li cell. The high-frequency arc could be associated with the charge-transfer reactions at the cathode and anode interface. Because the relaxation times for the cathodic and anodic reactions are generally overlapped over the high-frequency ranges for a NCM-Li full cell, it is difficult to separate the charge-transfer resistance at the Li-LLZTO interface from the resistance at the NCM-ionic liquid. The AC-impedance spectrum for Li/LLZTO/Li cell is also shown in Fig. S5b.† Considering that the electrochemical kinetics at the cathode would be almost the same for both cells, the variation in the high-frequency arc could be attributed to the charge-transfer resistance at the Li metal and LLZTO interface. The values of the overall interfacial resistance, Rct, were quantitatively determined from complex non-linear least-square (CNLS) fitting of the measured impedance spectra based on a typical equivalent circuit of the NCM-Li full cell. The Rct (Rc + Ra) value was reduced from ∼80 Ω cm2 to ∼26 Ω cm2 when the LLZTO surface was treated by a laser beam. In addition, as shown in Fig. S5c,† the electronic conductivity of LLZTO was significantly reduced by the laser treatment, which can originate from the formation of a wide band gap surface layer.

What's the cathode?

Prior to its integration with the electrodes, various properties of the LLZO SE such as the microstructure, crystalline phase and ionic conductivity were characterized. To improve the Li ion conductivity in the parent LLZO structures, Nb5+ was doped at the Zr4+ sites to create Li+ vacancies, and Ba2+ was doped at the La3+ sites to increase the Li+ concentration.Fig. 1a shows the cross-sectional SEM image of the LLZO SE pellet. A dense microstructure was formed by the uniaxial pressing and the subsequent sintering steps. The relative density of the LLZO pellet was ∼95% which was the ratio of the theoretical density calculated using the lattice parameter derived from XRD and the real density measured using a He gas pycnometer. XRD analysis of the LLZO SE (Fig. 1b) reveals the cubic phase of LLZO, where most of the diffraction peaks relate to the parent Li–garnet phase “Li5La3Nb2O12” with the space group Iad. A small peak at ∼2θ = 30 degree corresponds to the presence of BaZrO3 impurities. LLZO exists in cubic and tetragonal polymorphs. The former is typically obtained at higher sintering temperatures (>1000 °C), and exhibits about two orders of magnitude higher Li-ion conduction.Fig. 1c shows the Nyquist plots of a Au|LLZO|Au quasi-blocking cell obtained at room temperature. Au metal was sputtered on either side of the LLZO pellet to obtain uniform contacts. The plots were fitted with a resistor–capacitor (constant phase element) circuit and a series capacitor. By fitting the high frequency (5–0.5 MHz) semi-circle, we obtain a capacitance of ∼2.7 × 10−10 F which could be attributed to the bulk and grain-boundary contributions to Li+ transport in the SE. In the low frequency range (<100 Hz), we observe a linear increase which reflects the quasi-blocking effect of the Au electrodes. From the low-frequency intercept of the semi-circle with the x-axis we determine a bulk resistance (RB) of ∼640 ± 1.4 Ω cm2. When using l as the thickness of the SE pellet and a as the electrode surface area, the ionic conductivity σ was calculated using the formula: σ = l/RBa, which was found to be ∼1.5 × 10−4 S cm−1. Further, σ is plotted as a function of increasing temperature T in Fig. 1d and the relationship follows an Arrhenius behavior as expressed by the following equation where σ is the conductivity of the electrolyte (S cm−1), A is the pre-exponential factor, T is the temperature (Kelvin), Ea is the activation energy expressed (eV), and kb is Boltzmann's constant. The activation energy Ea calculated from the Arrhenius plot is 0.35 eV, which is consistent with values reported in the literature. The EIS plots used to measure Ea are given in Fig. S1.†

What's the electrolyte?

Bulk electrolysis was conducted to test the performance of CRLEs under conditions of electrochemical turnover. A series of charge/discharge experiments was performed for the negative and the positive electrolyte separately. The coulombic efficiency (CE) was calculated from photometric and coulometric data according to eqn (5), where npt are the number of moles obtained with photometry (c is the concentration, V is the volume, and 0.8 is the SoC difference), and nec the number of moles calculated from Faraday's law (I is the current, t is the time and F is the Faraday's constant).

What's the electrolyte?