milerssliu/battery-device-data-qa huggingface dataset:
Electrochemical impedance spectroscopy (EIS) curves are presented in Fig. S8.† A larger slope, which is associated with the Warburg impedance (Ws), was observed in the low-frequency region of the curve of the Fe-intercalated ML Ti3C2Tx electrode. Such an occurrence is related to the migration of electrolyte ions within the electrode. It is inferred that the faster kinetics of EMIm+ were derived from the increased space caused by Fe pre-intercalation. Ex situ Ti 2p and Fe 2p XPS spectra are presented in Fig. S9;† no obvious shifts were observed. This finding indicates that the charge storage mechanism of the Fe pre-intercalated ML Ti3C2Tx anode is based on the intercalation/de-intercalation behavior of EMIm+ instead of a redox process. Therefore, the larger real interlayer space contributed to the enhanced electrochemical performance. The scheme in Fig. 5d shows the charge storage mechanism of this DIB. Upon charging, EMIm+ cations were inserted into Fe pre-intercalated ML Ti3C2Tx, whereas PF6− anions were intercalated into the graphite cathode. Charge–discharge profiles of the full DIB are presented in Fig. S10†. The DIB provided an energy density of 76 W h kg−1 at a power density of 360 W kg−1. The cycling stabilities and capacity retentions of the electrodes were examined via performing GCD tests at the same current density. After 50 cycles, the DIB retained 94% of its initial capacity (Fig. 5e), indicating its improved stability compared with a DIB using the ML Ti3C2Tx anode (71%). Ragone plots are presented in the inset of Fig. 5e. The full DIB displayed a good rate capacity, with energy densities of 76, 70, and 64 W h kg−1 at power densities of 360, 900, and 1800 W kg−1, respectively. However, when the power density was higher than 3600 W kg−1, the energy density was reduced. The dual-ion intercalation mechanism of this DIB provides a new route for multi-ion storage.

Question: What's the cathode?
Answer: graphite
Electrochemical impedance spectroscopy (EIS) curves are presented in Fig. S8.† A larger slope, which is associated with the Warburg impedance (Ws), was observed in the low-frequency region of the curve of the Fe-intercalated ML Ti3C2Tx electrode. Such an occurrence is related to the migration of electrolyte ions within the electrode. It is inferred that the faster kinetics of EMIm+ were derived from the increased space caused by Fe pre-intercalation. Ex situ Ti 2p and Fe 2p XPS spectra are presented in Fig. S9;† no obvious shifts were observed. This finding indicates that the charge storage mechanism of the Fe pre-intercalated ML Ti3C2Tx anode is based on the intercalation/de-intercalation behavior of EMIm+ instead of a redox process. Therefore, the larger real interlayer space contributed to the enhanced electrochemical performance. The scheme in Fig. 5d shows the charge storage mechanism of this DIB. Upon charging, EMIm+ cations were inserted into Fe pre-intercalated ML Ti3C2Tx, whereas PF6− anions were intercalated into the graphite cathode. Charge–discharge profiles of the full DIB are presented in Fig. S10†. The DIB provided an energy density of 76 W h kg−1 at a power density of 360 W kg−1. The cycling stabilities and capacity retentions of the electrodes were examined via performing GCD tests at the same current density. After 50 cycles, the DIB retained 94% of its initial capacity (Fig. 5e), indicating its improved stability compared with a DIB using the ML Ti3C2Tx anode (71%). Ragone plots are presented in the inset of Fig. 5e. The full DIB displayed a good rate capacity, with energy densities of 76, 70, and 64 W h kg−1 at power densities of 360, 900, and 1800 W kg−1, respectively. However, when the power density was higher than 3600 W kg−1, the energy density was reduced. The dual-ion intercalation mechanism of this DIB provides a new route for multi-ion storage.

Question: What's the anode?
Answer: ML Ti3C2Tx
From the linear relationship of Ip and ν0.5 (Fig. S13†), (cathodic peak at 2.3 V), (cathodic peak at 1.9 V), and (anodic peak at 2.4 V) were obtained. The DLi+ values of the Co5.47Nx/S cathode were 219.9, 14.8, and 31.5 × 10−15 cm2 s−1 for peak A, B, and C, respectively. Importantly, the DLi+ of peak A of Co5.47Nx/S was up to four fold that of the Co5.47N/S cathode (48.6, 18.1, and 18.1 × 10−15 cm2 s−1 for peak A, B, and C, respectively), as plotted in Fig. 6d. The adsorption efficacy of cobalt nitride is attributed to its strong chemical interaction with LiPSs through both Co–S and N–Li bonds. The strong Li–N bonding in Co5.47N without nitrogen vacancies impedes direct electron transfer to the LiPSs and delays Li+ diffusion, resulting in sluggish reaction kinetics. Compared with Co5.47N, Co5.47Nx with nitrogen vacancies has a lower N content, which leads to decreased N–Li bonding and accelerated Li+ transportation. The cycling stability of Co5.47Nx and Co5.47N was evaluated at a current density of 0.5C, and the results are shown in Fig. S14.† After 200 cycles, the Co5.47Nx electrode showed a capacity retention of 85%, which was higher than the 76% for the Co5.47N electrode without nitrogen vacancies. These results suggest that Co5.47Nx facilitates Li+ transportation through its hierarchical structure and nitrogen vacancies, resulting in promising cycling and rate capacity.

Question: What's the cathode?
Answer:  Co5.47Nx/S
From the linear relationship of Ip and ν0.5 (Fig. S13†), (cathodic peak at 2.3 V), (cathodic peak at 1.9 V), and (anodic peak at 2.4 V) were obtained. The DLi+ values of the Co5.47Nx/S cathode were 219.9, 14.8, and 31.5 × 10−15 cm2 s−1 for peak A, B, and C, respectively. Importantly, the DLi+ of peak A of Co5.47Nx/S was up to four fold that of the Co5.47N/S cathode (48.6, 18.1, and 18.1 × 10−15 cm2 s−1 for peak A, B, and C, respectively), as plotted in Fig. 6d. The adsorption efficacy of cobalt nitride is attributed to its strong chemical interaction with LiPSs through both Co–S and N–Li bonds. The strong Li–N bonding in Co5.47N without nitrogen vacancies impedes direct electron transfer to the LiPSs and delays Li+ diffusion, resulting in sluggish reaction kinetics. Compared with Co5.47N, Co5.47Nx with nitrogen vacancies has a lower N content, which leads to decreased N–Li bonding and accelerated Li+ transportation. The cycling stability of Co5.47Nx and Co5.47N was evaluated at a current density of 0.5C, and the results are shown in Fig. S14.† After 200 cycles, the Co5.47Nx electrode showed a capacity retention of 85%, which was higher than the 76% for the Co5.47N electrode without nitrogen vacancies. These results suggest that Co5.47Nx facilitates Li+ transportation through its hierarchical structure and nitrogen vacancies, resulting in promising cycling and rate capacity.

Question: What's the cathode?
Answer: Co5.47N/S
In this work, a mechanism is proposed to explain the phenomenon of columnar lithium metal deposition. An electrolyte additive, such as HF, is selectively reduced at high potential vs. Li/Li+ to form uniformly distributed crystalline LiF particles with preferred crystallographic texture which are then encased in an amorphous matrix of solvent reduction products, as evidenced by systematic electrochemical analysis and in situ X-ray surface scattering. The LiF quantity and distribution are directly tunable by choosing an appropriate additive concentration and electrochemical cycling rate. Interfaces between LiF and the amorphous phase act as fast lithium-ion diffusion pathways, promoting a thin and more uniform SEI which leads to a very high lithium metal nucleation density relative to additive-free systems followed by nearly-isotropic and eventually vertical columnar growth, observed in real-time using operando small angle X-ray scattering. Understanding this process step-by-step provides new insights into the role of electrolyte additives and provides new information for the rational design of such additives. As HF is damaging to cathodes, current collectors, and other cell components, new additives which decompose at high potentials vs. lithium and direct the formation of a columnar morphology should be developed for practical use in lithium metal batteries.

Question: What's the cathode?
To summarize, multiple organic-based active materials showed promising characteristics in potassium-based batteries. Their capacities and potentials are typically less attractive than those of the state-of-the-art inorganic analogs, but the demonstrated rate and cycle capabilities make them highly competitive. There are still plenty of opportunities to improve the characteristics of the organic-based compounds using the truly unlimited diversity of molecular structures and tunability of their properties via functional group substitution. New advanced cathode and anode materials can be developed via rational molecular design, morphology optimization, engineering conductive fillers, and tuning the electrolyte composition.

Question: What's the cathode?
To summarize, multiple organic-based active materials showed promising characteristics in potassium-based batteries. Their capacities and potentials are typically less attractive than those of the state-of-the-art inorganic analogs, but the demonstrated rate and cycle capabilities make them highly competitive. There are still plenty of opportunities to improve the characteristics of the organic-based compounds using the truly unlimited diversity of molecular structures and tunability of their properties via functional group substitution. New advanced cathode and anode materials can be developed via rational molecular design, morphology optimization, engineering conductive fillers, and tuning the electrolyte composition.

Question: What's the anode?
To summarize, multiple organic-based active materials showed promising characteristics in potassium-based batteries. Their capacities and potentials are typically less attractive than those of the state-of-the-art inorganic analogs, but the demonstrated rate and cycle capabilities make them highly competitive. There are still plenty of opportunities to improve the characteristics of the organic-based compounds using the truly unlimited diversity of molecular structures and tunability of their properties via functional group substitution. New advanced cathode and anode materials can be developed via rational molecular design, morphology optimization, engineering conductive fillers, and tuning the electrolyte composition.

Question: What's the electrolyte?
After structural comparison, the activity of oxygen and manganese was further probed to reveal the charge compensation in electrochemistry. Operando DEMS was performed to monitor the released anionic species during the charging process. The operando DEMS data was collected with active material loading of around 8 mg (Fig. 3). In the initial charge process for typical LMO, carbon dioxide gas was evolved at the beginning, accompanied by a large amount of oxygen gas release with maximum carbon dioxide evolution. The total gas evolution was summed to be 59.3 mmol CO2 per mol of active material and 70.2 mmol O2 per mol of active material. The CO2 gas may result from the electrolyte decomposition, the reaction between the electrolyte and O2, and impurity Li2CO3 decomposition. In contrast, the T-LMO electrode delivered a minimal amount of carbon dioxide and oxygen gas release. The cumulative CO2 and O2 detected from T-LMO during the initial charge process was 10.5 mmol and 9.1 mmol per mol of active material, respectively. The reduced gas release in the treated sample indicates the improved electrochemical stability of the T-LMO material when operating at high voltage in the battery. The different outgassing behaviors of LMO and T-LMO under identical cycling conditions imply that the two materials have different oxygen activities.

Question: What's the electrolyte?
The normalized voltage characteristics for the 1st and 2nd cycles at C/20 (standard CC protocol) for Li[Li0.2Mn0.6Ni0.1Co0.1]O2 are presented in Fig. 2. During the 1st charge, we observed a rapid increase of the potential up to 4.4 V followed by a pseudo-plateau, while during the 1st discharge, the potential declines down to 3.5 V and a subsequent long pseudo-plateau was registered. The observed substantial hysteresis between charge and discharge potentials suggests: (i) sluggish kinetics of the occurring process and (ii) structural reorganization taking place. Comparing the 1st and 2nd cycles, we noticed that the 2nd charge characteristics do not show a large pseudo-plateau between 4.4 and 4.8 V but rather a monotonic S-type shape. Interestingly the potential of the 2nd discharge curve is slightly higher than that of the 1st discharge. In turn, the lower potential for the 2nd charge characteristics (in relation to the 1st one) leads to much lower hysteresis for the 2nd cycle compared with that for the initial one. This transition from stair-case to S-type potential characteristics between the 1st and 2nd cycles is commonly observed for other Li-rich NMC compounds; however the exact mechanism is still elusive and will be investigated further in this study. Comparing the discharge capacities at various current densities (Fig. 3a), the significant discharge capacity fade from approximately 350 to 290 mA h g−1 during the first two cycles is observed. Furthermore, the specific charge only slightly responds to the increase of the current loads and each rate increase (C/20, C/10, C/5, C/2, and 1C) is followed by a minor decrease of the practical specific charge until reaching 150 mA h g−1 at a 1C rate. Going back to the C/20 rate leads to an increase of the discharge capacity slightly below 250 mA h g−1. A similar observation was made in other reports within the literature where the first discharge capacity is above 300 mA h g−1 and after a few cycles it drops to ca. 250 mA h g−1 comparable to that in our paper. We suspect that most probably the anomalously high discharge capacity is due to the nanosize of primary particles which have a very high total surface area and as such enhance the electrolyte decomposition, increasing the observed specific charge. It is also worth noticing that after rate capability tests, during the second cyclic tests at C/20 the material works with a reversible capacity of 250 mA h g−1, suggesting most probably that the observed very high initial discharge capacity is due to the parasitic reactions. Voltage profiles for the second cycles registered at various current loads are shown in Fig. 3b and exhibit almost the same S-type character, confirming that reorganization taking place between the 1st and 2nd cycles persists in the following tests.

Question: What's the electrolyte?
The electrochemical performance of LTO electrodes coated with various materials was evaluated by linear sweep voltammetry (Fig. 3). Hydrogen evolution began at ∼1.8 V vs. Li on the pristine LTO surface, which was higher than its lithiation potential (1.55 V). The HER process (rather than lithiation of LTO) dominated the cathodic reaction during the scan. The carbon coating enhanced the electronic conductivity of LTO, thus accelerating the HER (as evidenced by the higher currents and positively shifted HER potential). The TiO2 coating also positively shifted the cathodic limit due to its high catalytic activity. In contrast, the ZnO coating and Al2O3 costing negatively shifted the cathodic limit potential by 0.1 V. In addition, the HER currents on the ZnO-coated and Al2O3-coated LTO electrode were much lower compared with those on pristine LTO. The capability of HER suppression as quantified by the onset potential of the HER should increase in the order TiO2 < carbon < LTO < ZnO < Al2O3, which was consistent with the prediction in Fig. 1. The Al2O3 coating not only suppressed the HER catalytic activity but also acted as a kinetic barrier to slow down electron transfer from the electrode bulk to protons in the electrolyte. Al2O3 coating shifted the HER potential to <1.5 V, so Li+ intercalation was enabled before the HER, as evidenced by a sharp lithiation peak at 1.55 V (Fig. 4a). Using a similar approach, we also evaluated the effect of surface coating on the oxygen evolution reaction (OER). Al2O3 coating and TiO2 coating on LiNi0.5Mn1.5O4 reduced the side reactions on this high-voltage (4.8 V) cathode material only slightly (Fig. S3†).

Question: What's the cathode?
Answer: LiNi0.5Mn1.5O4
The electrochemical performance of LTO electrodes coated with various materials was evaluated by linear sweep voltammetry (Fig. 3). Hydrogen evolution began at ∼1.8 V vs. Li on the pristine LTO surface, which was higher than its lithiation potential (1.55 V). The HER process (rather than lithiation of LTO) dominated the cathodic reaction during the scan. The carbon coating enhanced the electronic conductivity of LTO, thus accelerating the HER (as evidenced by the higher currents and positively shifted HER potential). The TiO2 coating also positively shifted the cathodic limit due to its high catalytic activity. In contrast, the ZnO coating and Al2O3 costing negatively shifted the cathodic limit potential by 0.1 V. In addition, the HER currents on the ZnO-coated and Al2O3-coated LTO electrode were much lower compared with those on pristine LTO. The capability of HER suppression as quantified by the onset potential of the HER should increase in the order TiO2 < carbon < LTO < ZnO < Al2O3, which was consistent with the prediction in Fig. 1. The Al2O3 coating not only suppressed the HER catalytic activity but also acted as a kinetic barrier to slow down electron transfer from the electrode bulk to protons in the electrolyte. Al2O3 coating shifted the HER potential to <1.5 V, so Li+ intercalation was enabled before the HER, as evidenced by a sharp lithiation peak at 1.55 V (Fig. 4a). Using a similar approach, we also evaluated the effect of surface coating on the oxygen evolution reaction (OER). Al2O3 coating and TiO2 coating on LiNi0.5Mn1.5O4 reduced the side reactions on this high-voltage (4.8 V) cathode material only slightly (Fig. S3†).

Question: What's the electrolyte?
As ideal self-template precursor, metal–organic frameworks (MOFs) with regular compositions can be converted into porosity-tunable materials with a uniform distribution of metal-based components and carbon. The nano-sized metal-based components sustaining the framework of MOFs are able to curtail the distance of ion diffusion and create more electrolyte-accessible sites. Porous carbon, derived from decomposition of the organic linkers, can efficiently alleviate the volume variation of the metal-based components during the ion insertion/extraction process, and it can also reinforce electron transmission. In addition, the incorporation of heteroatoms can be easily realized by the modification of organic linkers in MOFs. Recently, some studies have reported that MOF-derived nanocomposites are promising anode materials for LIHCs and SIHCs, and they can facilitate boosting the pseudocapacitive-dominant charge storage as well as maintain the structure stability.

Question: What's the anode?
Answer: MOF-derived nanocomposites
As ideal self-template precursor, metal–organic frameworks (MOFs) with regular compositions can be converted into porosity-tunable materials with a uniform distribution of metal-based components and carbon. The nano-sized metal-based components sustaining the framework of MOFs are able to curtail the distance of ion diffusion and create more electrolyte-accessible sites. Porous carbon, derived from decomposition of the organic linkers, can efficiently alleviate the volume variation of the metal-based components during the ion insertion/extraction process, and it can also reinforce electron transmission. In addition, the incorporation of heteroatoms can be easily realized by the modification of organic linkers in MOFs. Recently, some studies have reported that MOF-derived nanocomposites are promising anode materials for LIHCs and SIHCs, and they can facilitate boosting the pseudocapacitive-dominant charge storage as well as maintain the structure stability.

Question: What's the electrolyte?
Based on the above discussion, a driving mechanism of the EC procedures in the shoulder-by-shoulder structure was proposed (Fig. 2d) which is similar to that of a planar-arrangement supercapacitor. When both WO3 films were transparent during the first process, in essence, the conductive ITO layers were likely acting as two split capacitance plates. Once a working voltage was applied, Li ions were driven by the edge effect of the electric field of the capacitance plates and intercalated into the WO3 film at the negative potential, namely PI. Because the edge effect of the electric field was intense in the middle of the two EC films, more Li ions were injected into this area than into other areas, showing an uneven color distribution. These concentration differences resulted in the self-diffusion of Li ions in the colored film, namely PIII. However, when one EC unit was colored in the subsequent processes, it was likely acting as a charged capacitor. Once a reverse driving voltage was employed, the charge transfer procedure (namely PII) preferentially occurred. Electrons were transferred through the outer loop, while Li ions were transferred through the electrolyte, and they simultaneously intercalated into another EC film that was at a negative potential and colored it. The ions (electrons and Li ions) in the area of the colored film that was next to the middle of the two EC units were first transported into its counterpart and placed in the area that was also next to the seam because the electric field in the seam area was very intense. It is reasonable that PII is much quicker than PI because discharging a capacitor is a rapid process. When PII finished, the EC film was partly colored. Next, PI and PIII both occurred, and the partly colored film became completely colored. When the working voltage was turned off, the colored film became evenly colored, driven by PIII.

Question: What's the electrolyte?