Подбор оптимальных условий синтеза электродного материала Na4Nb8P4O32 для натрий-ионных аккумуляторов

All-solid-state Li batteries are regarded as a highly promising system for future electrochemical energy storage. Oxide-ceramic based all-solid-state lithium batteries (ASB) can provide high intrinsic safety, extended operational temperature range and high energy density.As the first two are intrinsic to the materials system, one prerequisite to obtain high energy densities with such ASBs is the manufacturing of thick (approx. 150 µm), dual-conducting free-standing composite cathodes, analogous to that of conventional liquid electrolyte-based lithium batteries. However, the preparation of composite cathodes using oxide-ceramic electrolytes is challenging since high temperature sintering steps are necessary during electrode and cell manufacturing to achieve proper mechanical stability, contact between the individual phases and good ionic and electronic conductivity.[1, 2] For oxide-based ceramic electrolyte materials like Li7La3Zr2O12 (LLZ) or Li1+xAlxTi2-x(PO4)3 (LATP) a sintering temperature around 1000 °C is necessary. However, cathode active materials like spinels (e. g. Li2NiMn3O8) or layered materials (e. g. Li[Ni1-x-yCoxMny]O2 (NCM)) show thermal stability only to around 700 °C and the required sintering temperature for the electrolyte exceeds their thermal stability window.To mitigate these materials interactions, advanced sintering techniques like high-pressure Field Assisted Sintering Technique/ Spark Plasma Sintering (FAST/SPS) is required, as it allows low temperature consolidation.[3]Using high-pressure FAST/SPS we have developed a method to prepare thick mixed-conducting composite cathodes and half-cells at sintering temperatures as low as 700 °C. The analysis of these composite cathodes revealed well sintered pure phases and a homogenous distribution of cubic LLZ and cathode active material. Electrochemical tests showed promising electronic and ionic conductivity (0.4 mS cm-1) values and an increased capacity (4 mA cm-2) compared to cathodes with pure active material.FAST/SPS can, therefore, open new processing windows for ceramic-based ASSLBs to alleviate the problem of interphase formation.

Simultaneous measurements of electronic and ionic conductivities have been carried out for the liquid Ag x (TlSe) 1-x system with x=0 to 0.52, applying the residual potential theory. Both the electronic and ionic conductivities increase with increasing temperature at the whole composition range investigated, except for the latter conductivity at x=0.50 (corresponding to the stoichiometric composition of AgTlSe). The system exhibits a deep minimum in the electronic conductivity and sharp maximum in the ionic conductivity at this composition. The maximum value of ionic conductivity in the system is nearly one order of magnitude smaller than the minimum one of electronic conductivity. The value of g-factor obtained by the relationship of electronic conductivity and magnetic susceptibility is roughly 0.025 at x=0.5 at 773 K, implying the occurrence of the electron localization. Further, the number density and mobility for metal ions in the system were estimated using the observed ionic conductivity under the assumption that all of Ag and Tl atoms in the system are univalently ionized and contribute to the ionic conduction. It has been concluded that the ionic mobility drastically increases, while, conduction electrons are strongly localized, at the stoichiometric composition of AgTlSe in the liquid Ag-TlSe system.

For Li7La3Zr2O12 (LLZ) and Li1-3x La x TiO3, which have been investigated as electrolytes for oxide solid-state batteries[1.2], we have proposed a pyrochlore-type oxyfluoride Li2-x La(1+x)/3M2O6F (M = Nb, Ta etc.) with high ionic conductivity[3]. In this study, electrochemical properties of pyrochlore-type solid electrolyte were investigated.Li2-x La(1+x)/3Nb2O6F powder was synthesized by a solid-state reaction. The synthesized powder was identified by XRD analysis and Li2-x La(1+x)/3Nb2O6F pellets were prepared at 1000°C for 6h in order to evaluate their electrochemical properties. Impedance was measured by sputtering Ag electrodes on both surface of the pellets. The activation energy was calculated from the value of ionic conductivity at -80~120°C. Then, the Li symmetrical cell was assembled, and deposition/dissolution cycle of Li was carried out. Separator (Celgard 3501) containing liquid electrolyte (1M LiPF6 in EC/DEC) were introduced at the interface between the Li metal foil and the solid electrolyte (Li | separator | pellet | separator | Li). The test conditions were 25°C and 0.1 mA/cm2, and the cycle was carried out by repeating charging/discharging for 0.5h each. The garnet-type solid electrolyte LLZ, which is to be compared, was synthesized by a solid-state reaction in the same manner as above, and pellets were prepared at 900°C for 10h.XRD analysis of the synthesized powder confirmed the diffraction pattern from Li2-x La(1+x)/3Nb2O6F. The relative density of the pellets prepared using this powder was about 84% and LLZ pellets was about 79%. Figure 1 shows Arrhenius and Cole-Cole plots of Li2-x La(1+x)/3Nb2O6F and LLZ. The total ionic conductivities of Li2-x La(1+x)/3Nb2O6F were 3×10-5, 5×10-3, and 2×10-2 S/cm at -80°C, 25°C, and 120°C respectively, and the activation energy was 0.20 eV. Li2-x La(1+x)/3Nb2O6F showed higher ionic conductivity and lower activation energy than LLZ. Figure 2 shows the results of deposition/dissolution cycle of Li. It was observed that Li2-x La(1+x)/3Nb2O6F has a lower overpotential than LLZ. In addition, the short circuit behavior of LLZ with the increase of overpotential occurred from around 40 cycles, while Li2-x La(1+x)/3Nb2O6F was stable after 50 cycles. And it was also found that the value of the overpotential at 1cyc almost coincides with the one estimated from the cell resistance obtained by the impedance measurement. These results indicate that the mobile ion of Li2-x La(1+x)/3Nb2O6F is Li+ and the ionic conductivity is high. The cycle stability of Li2-x La(1+x)/3Nb2O6F may be due to the high ionic conductivity of Li2-x La(1+x)/3Nb2O6F, the suppression of current density heterogeneity, and the uniform progress of Li deposition/dissolution. In this presentation, the results of other physical and electrochemical properties are also reported.

Zn-MnO2 alkaline batteries are prevalent and reliable primary batteries due to their cost effectiveness, good safety characteristics, long shelf life and ability to deliver sustained power for continuous usage. This properties make alkaline batteries an ideal options to be used in remote controls, smoke detectors, wireless mice, toys and many portable electronic devices. Despite their advantages, the battery performance is constrained by the slurried Zinc anode. Because of this, researchers have prioritized the improvement of the anode, primarily aiming to enhance the capacity and stability of the active material. However, the actual behavior of the anode is limited by the ionic conductivity of Zinc slurry, which is several orders of magnitude lower than the electronic conductivity of the slurry. This causes the reacting front to begin at the anode-separator interface, leading to passivation of the anode, incomplete discharge and increased gassing.The utilization of Zinc in the slurry anode can be increased by bringing the values of ionic conductivity and electronic conductivity closer together. Unfortunately, due to the redox electrolyte mechanism that moves electrons through the anode [1], the literature does not have an established method that is able to independently measure the ionic and electronic conductivity of the slurry. Before methods can be proposed to improve ionic conductivity of slurry, a method must first be devised that is able to truly measure the ionic conductivity in isolation.In this study, a novel methodology is reported to quantify both the ionic and electronic conductivity of the slurried anode. This innovative approach aims to provide a deeper understanding of the complex interconnection between electronic and ionic conductivities of the Zinc slurry. By addressing this critical aspect, the research contributes to advancing the fundamental understanding of Zn-MnO2 alkaline batteries, potentially paving the way for significant improvements in their performance and reliability for a wide array of applications.

Study on the effect of intrinsic electrical resistivity of silicon materials on its performance for Li-ion batteries

Polymer nanocomposite electrolyte (PNCE) films composed of poly(ethylene oxide) (PEO), lithium salt (\( {\text{LiX}};\;{\text{X}} = ClO_4^{ - },\;BF_4^{ - },\;C{F_3}SO_3^{ - } \)) and montmorillonite (MMT) clay as nanofiller were prepared by melt-compounded hot-pressed technique at 70 °C under 3 tons of pressure. The ionic conductivity and relaxation behaviour of the films were investigated by dielectric relaxation spectroscopy in the frequency range of 20 Hz to 1 MHz at ambient temperature. The results revealed that the ionic conductivity of the PNCE films having 20:1 stoichiometric ratio of ethylene oxide monomer units to the lithium cation are governed by the size of different anions and the dissociation constant of salt, and also MMT concentration. It was found that PEO–LiBF4 film has comparative high dc ionic conductivity, whereas both the LiBF4 and LiClO4 containing PNCE films exhibit anomalous conductivity behaviour with varying MMT concentration. The PEO–LiCF3SO3 film has two orders of magnitude low value of dc ionic conductivity as compared to that of the other salts electrolyte films, but its conductivity enhances by one order of magnitude when 2 wt.% MMT is added as filler. A correlation between the values of ionic conductivity, conductivity relaxation time and the real part of permittivity at 1 MHz were found and the same was discussed in relation to the transient ion-dipolar type cross-linked structural behaviour of the polymeric nanocomposite electrolytes.

Slurry electrodes have gained traction in various electrochemical systems such as electrochemical flow capacitors, flow batteries, primary Zn-MnO2 batteries and rechargeable Zn batteries. Such electrodes are formed by suspending solid active particles in a high viscosity gel. Devices with slurry electrodes are used because they offer high volumetric capacity, high surface area and reduced complexities in manufacturing and recycling.However, charge conduction in slurry electrodes is complex, in many cases relying on both the ionic and electronic conductivity, with charge being carried through the electrolyte (as ions) and by the particles (as electrons). The dynamics of charge conduction are influenced by the nature of the suspended particles, additives, concentration of the supporting electrolyte and underlying reaction mechanisms. This is particularly evident in Zn-MnO2 alkaline batteries where the performance of the slurred zinc anode is limited by the ionic conductivity of Zinc slurry, which is several orders of magnitude lower than the electronic conductivity of the slurry. This causes the reacting front to begin at the anode-separator interface, leading to passivation of the anode, incomplete discharge and increased gassing.Therefore, it is important to increase the value of the ionic conductivity. But, due to the complex charge transport mechanism, the ionic and electronic conductivities are ultimately coupled [1]. As such, the literature does not have an established method that is able to independently determine the effective ionic and electronic conductivities when integrated into the slurry.In this study, a novel methodology is reported to quantify ionic conductivity of slurried anode electrodes, with the primary focus being the Zn anode in primary Zn-MnO2 batteries. This innovative approach aims to provide a deeper understanding of the complex interconnection between electronic and ionic conductivities of the Zinc slurry – and allows the effects of several variables to be quantified. This work can contribute not only to advancing the fundamental understanding and application of Zn-MnO2 alkaline batteries, but also might pave the way for significant improvements in the performance and reliability of a wide array of applications. References Faegh, T. Omasta, M. Hull, S. Ferrin, S. Shrestha, J. Lechman, D. Bolintineanu, M. Zuraw and W.E. Mustain, “Understanding the Dynamics of Primary Zn-MnO2 Alkaline Battery Gassing with Operando Visualization and Pressure Cells”, J. Electrochem. Soc., 165 (2018) A2528-A2535

Introduction Lithium air batteries exhibit higher theoretical energy density than lithium ion batteries and are expected to be used in the next generation of secondary batteries. However, there are many problems, such as a decrease in discharge capacities after some cycles and a large difference between discharge and charge voltages, ΔV. A large/small ΔV means low/high round-trip (discharge-charge) energy efficiency, respectively. High-performance secondary batteries show small ΔV, i.e., high round-trip energy efficiency. Since the first report by K. M. Abraham et al. [1], various kinds of oxygen reduction/evolution catalysts[1-4] and electrolyte[1,5] materials have been intensively investigated for air batteries to improve their electrochemical properties such as the cycleability and the round-trip energy efficiency of the air batteries. The purpose of this research is to improve the round-trip energy efficiency by using highly-active catalysts for air electrodes. We are focusing on RuO2 as the catalyst and here report the performance of air batteries incorporating this oxide catalyst. Experimental Precursor powder of RuO2, Ru(OH)n, was prepared neutralizing 0.1 mol/l RuCl3 aq with 0.1 mol/l NaOH aq. RuO2 was obtained by heat-treating the hydroxide powder at 110, 200, and 500◦C. An air electrode was prepared by rolling a mixture of RuO2, KetjenBlack EC600JD (KB), and PTFE (10:54:36 in weight ratio) into a sheet about 0.5-mm thick. The lithium air battery consisted of the air electrode, a lithium metal shee,t and 1.0 mol/l lithium bis(trifluoromethanesulfonyl)imide (LiTFSI)/aprotic organic solvent as positive electrode, negative electrode, and electrolyte solution, respectively. Solvents used were propylene carbonate (PC), tetraethylene glycol dimethyl ether (TEGDME), and dimethyl sulfoxide (DMSO). The preparation of the battery is described in detail in our previous paper [4]. Electrochemical measurements were carried out under a galvanostatic condition of 0.1 mA/cm2 in a dry air atmosphere with a dew point of less than -50◦C at room temperature. The discharge/charge capacities were normalized by the weight of the air electrodes. Results and discussion Figure 1 shows XRD patterns of RuO2 heat-treated at 110, 200 and 500◦C. All the peaks correspond to the PDF data for RuO2 (#01-070-2662). The peaks become sharper as the heat-treatment temperature increases. This indicates that the particle size of RuO2 became larger. It is notable that RuO2 is crystallized even at a low temperature of 110◦C. These RuO2particles would be very fine and suitable as catalyst for the air electrode.Figure 2 shows the first discharge/charge curves of the air batteries incorporating the RuO2 catalyst heat-treated at 110, 200, and 500◦C. Compared with KB only, the air batteries with RuO2 catalyst showed larger charge capacities. The charge capacities became larger with RuO2 prepared at lower heat-treatment temperature. As for the cell voltage, the RuO2 catalyst reduced the charge overvoltage in particular, even though there were no changes in the discharge overvoltages. The charge overvoltages became smaller for RuO2 with the lower heat-treatment temperature. The tendencies observed in the capacities and voltages with the heat-treatment temperature were almost the same. These results indicate that the fine-powder RuO2catalyst prepared at the lower temperature had great effects on not the discharge but on the charge capacity/overvoltage.Figure 3 shows the first discharge/charge curves of the air batteries incorporating the RuO2catalyst in the electrolyte solution of 1.0 mol/l LiTFSI/PC, TEGDME, and DMSO. The air batteries with the TEGDME solution showed smaller discharge and charge capacities and larger discharge and charge overvoltages compared with the PC solution. On the other hand, the air batteries with the DMSO solution showed the largest discharge and charge capacities and the smallest discharge and charge overvoltages. In particular, the air batteries with the DMSO solution showed rather low average charge voltage of about 3.1 V. Such low charge voltage greatly improves the round-trip energy efficiency. This superior battery performance would be due to some properties of the DMSO solution such as its stability under the operation condition of the air battery. In conclusion, the fine RuO2powder prepared at the lower temperature reduced the charge overvoltage. Moreover, the use of the DMSO-containing solution led to the great improvement in the round-trip efficiency due to the decrease in charge overvoltage.

The development of advanced electrochemical energy storage systems is considered as one of the most important steps to drive the shift from fossil to renewable energies forward. Among the few alternatives to Li-ion batteries that would exceed significantly their specific energy, Li-oxygen and Na-oxygen batteries have attracted significant interest. In order to fully understand their electrochemical properties and to identify potential failure mechanisms, electrochemical properties and charge carrier transport of their discharge products (e.g. Na2O2) need to be known in detail. Here, we investigated the overall and electronic conductivity of nanocrystalline Na2O2 that was prepared by mechanical treatment of coarse-grained Na2O2 in a high-energy ball mill. The powder obtained is expected to be rich in defects and structural disorder. Thus, we expect that the sample investigated would be at the upper limit of reachable conductivities. Impedance measurements and DC-polarization experiments show that the overall conductivity is rather poor; it is in the order of 10−13 S/cm yielding diffusion coefficients in the order of 1.8 × 10−23 m2/s at approximately 300 K [1]. At room temperature the partial electronic conductivity ranges between 7.6 and 9.5 × 10−14 S/cm. Thus, the total conductivity in Na2O2 is largely influenced by electronic contributions rather than ionic conduction [2]. Furthermore, we also anticipate a strong coupling between ionic and electronic charge carrier transport. We measured the electronic conductivity at different temperatures to calculate the underlying activation energy. It turned out that electronic conductivity is governed by an activation energy as high as 0.95 eV while for overall charge transport values range from 0.85 and 1.03 eV depending on the samples exact properties and measurement conditions. Acknowledgement: Financial support by the Deutsche Forschungsgemeinschaft (DFG Research Unit 1277, grant no. WI3600/2-2 and 4-1 and HA6996/1-2) as well as by the Austrian Federal Ministry of Science, Research and Economy, and the Austrian National Foundation for Research, Technology and Development (CD-Laboratory of Lithium Batteries: Ageing Effects, Technology and New Materials) is greatly appreciated.

Introduction Lithium air secondary batteries exhibit higher theoretical energy density than lithium ion batteries and are expected to be used as the next generation of secondary batteries. However, there remain technical challenges related to their poor cycle properties. Since the first report by K. M. Abraham et al. [1], various oxygen reduction/evolution catalysts [1-4] and electrolyte [1, 5] materials have been intensively investigated for air batteries to improve their electrochemical properties, such as the their cyclability. However, there have been a few studies on support materials for air electrodes, for example, nanoporous gold [6]. We think that a decrease in the electrical contact of the air electrode during Li2O2deposition at discharge and oxygen generation at charge is one of the reasons for the poor cyclability. The purpose of our research is to improve the cyclability by using a new air electrode structure. We are focusing on a porous carbon monolith, which has a continuous interconnected network of carbon, as support material for air electrodes. We have already fabricated porous-carbon-monolith supports to replace the current carbon-powder supports. Here we report the preparation conditions for porous-carbon-monolith support materials and the performance of air batteries incorporating these supports. Experimantal A porous polyacrylonitrile (PAN) monolith was prepared as the precursor of the porous-carbon-monolith support. PAN was dispersed in dimethyl sulfoxide (DMSO). After the solution with a concentration of 94 g/l was dissolved, water was sprayed on the solution. The porous PAN monolith film was deposited from the solution due to water, which is poor solvent for PAN. The monolith was washed with methanol and dried in vacuum at room temperature. Then it was carbonized in Ar atmosphere at 1300 °C to obtain a porous carbon monolith. In the next step, for activation treatment, the porous carbon monolith was heat treated in CO2 atmosphere at 900 °C for 1 h [7]. The lithium air secondary battery consisted of the air electrode, a lithium metal sheet, and 1.0-mol/l lithium bis(trifluoromethanesulfonyl)amide (LiTFSA)/propylene carbonate (PC) as the positive electrode, negative electrode, and electrolyte solution, respectively. The battery preparation is described in detail in our previous paper [4]. Electrochemical measurements were carried out under a galvanostatic condition of 0.05 mA/cm2 in an O2atmosphere. The discharge and charge capacities were normalized by the weight of the air electrodes. Results and discussion In the SEM image of the porous carbon monolith in Fig. 1, we see the continuous interconnected network of the carbon. These particles build a 3-D disordered macroporous framework, with sizes in the range of 2-5 mm. In XRD measurements of the porous carbon monolith, all of the peaks in the XRD patterns corresponded to the ICDD data for carbon (#01-077-7164). These results indicate that our method can produce porous carbon monoliths. Figure 2 shows the pore distributions of the porous carbon monolith measured using mercury intrusion porosimetry. The porous carbon monolith has mesopores with size of ~10 nm, which correlated with first discharge capacities [8]. Figure 3 shows the first discharge/charge curves of the air batteries incorporating the as-synthesized and activation-treated porous carbon monolith. The air batteries, showed higher discharge and charge capacities, with the activation-treated porous carbon monolith. This is because activation-treated porous carbon monolith has many more active sites than the as-synthesized one. In particular, the air batteries with the activation-treated porous carbon monolith show the capacity of 12 mAh/g and the average discharge voltage of 2.6 V. This indicates that the porous carbon monolith can be used as support material for air electrodes. However, compared with Ketjen Black EC600JD (KB) powder [9], the air batteries with the porous carbon monolith show lower discharge capacities. It seems that KB-powder has many more active sites than the porous carbon monolith because its BET surface area of 1300 m2/g is larger than that of the porous carbon monolith, which is 4.3 m2/g.

Composite electrodes containing active materials, carbon and binder are widely used in Li-ion batteries. It is well known that their morphology influences the electrochemical performance of batteries. For the improvement of performance in practically used Li-ion batteries, it is necessary to tune the parameters of composite electrodes such as porosity, compounding ratio and thickness. The morphology of composite electrodes varies the effective electronic and ionic conductivity in composite electrodes. However the tuning of composite electrodes has been mainly based on the intuition and experiences, since it is difficult to distinguish the electronic conductivity and the ionic conductivity in composite electrodes. For composite electrodes, the traditional 4-probe method cannot be applied because charge-discharge currents are flew during applying voltage. Therefore, a measurement method for distinction between electronic and ionic conductivities in composite electrodes is required. In this study, we have developed a method for simultaneous measurement of electronic and ionic conductivities in composite electrodes. This method is applied to the various porosities composite electrodes. Fig. 1 shows the measurement setup and the arrangement of the electrodes and samples. Two Al foils were bonded with polypropylene. 75 wt% carbon-coated LiFePO4 powder, 10 wt% acetylene black and 15 wt% PVDF were mixed in 1-methyl-2-pyrrolidinone anhydrous (NMP, Sigma-Aldrich) solvent. The slurries were coated onto the aluminum foils. Drying was done at 70 °C to remove solvent and additional drying was performed at 80 °C in a vacuum oven to vaporize the residual solvent. These LiFePO4-based composite electrodes were pressed at 0 kgf, 300 kgf, 600 kgf, 900 kgf, 1200 kgf pressures to control their porosity. As shown in Fig.1, the electrochemical cell contains 6 probes connecting with the composite electrode. For the electronic and the ionic conduction electrodes, Al foils and Li foils were used, respectively. Two potentiostats and bias voltage were connected to the cell. After open circuit voltage is measured, the two potentiostats were operated with this voltage as the set point. And then, a bias voltage was applied between the two working electrodes. The ionic current was measured at A1 or A3 current meter, and the electronic current was measured at A2 or A4 current meter. The measurement principle is explained in the literature [1]. Fig. 2 shows the effective electronic and ionic conductivities for various porosity composite electrodes. The electronic conductivity is at least 10 times higher than the ionic conductivity. This indicates the dominant contribution to the electrochemical performance is the ionic conductivity. For the high porosity electrodes, the effective ionic conductivity is almost constant. On the other hand, when the porosities of the electrodes are less than 47%, the ionic conductivity decreased. This indicates that the compressed composite electrode structure narrowed the ionic conduction path, resulting the decrease of the ionic conductivity. From the result of the rate capability test in the LiFePO4 composite electrodes, the decrease of porosity causes the decrease of discharge capacity at 10 C rate. Therefore, the low ionic conductivity from low porosity lowers the rate capability of LiFePO4composite electrodes.

Comparison of phase composition, morphology and electrochemical property for Li3−xNaxV2(PO4)3 (x=0.5, 1.5 and 2.0) as lithium storage cathode materials

All-solid-state batteries have attracted attention because of their high energy density, safety, and long cycle life. Sulfide active materials exhibit high capacities and enable an enhanced energy density in all-solid-state batteries. In this study, we synthesized electrode-electrolyte bifunctional materials in the system Li2S-V2S3-LiX (X = F, Cl, Br, or I) through a mechanochemical process. In addition, the effects of the addition of lithium halides on the electrochemical properties were investigated. All-solid-state batteries with the Li2S-V2S3-LiI electrode showed the highest capacity of 400 mAh g-1 among all the cells, even though their electronic and ionic conductivities were the same. From the point of view of the ionic conductivity and structure of the electrodes during cycling, it was clarified that a high reversible capacity was achieved not only by high ionic and electronic conductivities before cycling but also by maintaining the ionic conductivity even at the deep state of charge. Furthermore, high-loading all-solid-state cells were fabricated using the Li2S-V2S3-LiI materials with a mass loading of 37.3 mg cm-2, exhibiting a high areal capacity of approximately 11.5 mAh cm-2 at 60 °C and good cycle performance.

탄화수소가스를 고온<TEX>$(1000^{\circ}C)$</TEX>에서 열분해 하여 고상화하는 기상 열분해법을 사용하여 저결정질 탄소재를 제조하고 같은 방법으로 붕소를 첨가한 저결정질 탄소재<TEX>$C_{l-x}B_x(x=0.05,\;0.10,\;0.20)$</TEX>를 제조하여, 리튬 이온 이차전지의 부극으로서의 전기화학적 특성을 조사하였다. 시료 대 PVDF를 95:5의 무게비로 첨가한 경우, 붕소를 첨가하지 않은 저 결정질 탄소재(x=0.00)는 초기 방전용량 374mAh/g을 나타내었으며, 제 2싸이클부터는 싸이클 성능이 비교적 우수하여 제 10싸이클에서 258mAh/g의 방전용량을 나타내었다. 시료 대 PVDF를 95:5의 무게비로 첨가한 경우, <TEX>$C_{1-x}B_x(x=0.00,\;0.05,\;0.10\;0.20)$</TEX> 시료들 중에서 x=0.05 조성의 시료는 가장 큰 초기 방전용량 860mAh/g을 나타내었으며, 10번째 싸이클에서 181mAh/g의 방전용량을 나타내었다. 제 2싸이클부터 싸이클 성능은 모두가 비슷하게 나타났다 초기방전 용량(PVDF <TEX>$10wt.\%$</TEX> 사용시, 853mAh/g), 싸이클 성능, 방전용량(PVDF <TEX>$10wt.\%$</TEX>사용시 10번째 싸이클에서 400mAh/g)면에서 <TEX>$C_{0.90}B_{0.10}$</TEX> 시료가 리튬이온 이차전지의 부극으로서의 가장 우수한 전기화학적 특성을 나타내었다. 합성한 탄소에 NMP를 용매로 한 액상 혼합 바인더(PVDF)를 90:10의 무게비로 첨가한 경우가 95:5의 무게비로 첨가한 경우보다 대체로 모든 조성에서 충<TEX>$\cdot$</TEX>방전용량이 크게 나타났다. 붕소가 첨가되어 덜 disordered된 구조가 됨으로써 1.25V보다 낮은 전압 부분에서 평탄구역이 증가하는 것으로 판단된다. 붕소가 첨가된 경우 충<TEX>$\cdot$</TEX>방전용량이 제 2싸이클에서부터 급격히 감소하였는데, 이는 첨가된 붕소가 제 1싸이클에서 삽입되는 Li과 일부는 강하게 결합하여 추출이 안되고 일부만이 다시 가역적으로 추출<TEX>$\cdot$</TEX>삽입되기 때문으로 생각된다. 붕소 첨가에 의한 충<TEX>$\cdot$</TEX>방전용량의 증가는, 붕소가 electron acceptor로 작용하여 삽입된 Li와 붕소-탄소 host 사이의 결합 강도를 증가시킴으로써 붕소치환 된 탄소에서 Li의 전위를 상승시키기 때문에 일어난다고 사려된다. Disordered carbon and boron-substituted disordered carbons <TEX>$C_{l-x}B_x(x=0.05,\;0.10,\;0.20)$</TEX> were synthesized by Pyrolysis of LPG(liquid Propane gas)and <TEX>$BCl_3$</TEX>. Their electrochemical properties as anode materials for Li-ion secondary batteries were then investigated. When PVDF is added to the sample in a weight ratio 5 : 95, the disordered carbon with x=0.00 had the first discharge capacity 374 mAh/g. Its cycling performance was relatively good from the second cycle and it had the discharge capacity 258 mAh/g at the 10th cycle. When PVDF is added to the sample in a weight ratio 5 : 95, the sample with x=0.05 among the samples <TEX>$C_{l-x}B_x(x=0.05,\;0.10,\;0.20)$</TEX> exhibited the largest first discharge capacity 860 mAh/g and discharge capacity 181 mAh/g at the 10th cycle. All the samples had similar cycling performances from the second cycle. The sample <TEX>$C_{0.90}B_{0.10}$</TEX> showed the best electrochemical properties as a anode materials fur Li-ion secondary battery from the view points of the first discharge capacity(853 mAh/g when <TEX>$10w1.\%$</TEX> PVDF is used), cycling performance, discharge capacity(400mAh/g at the 10th cycle when <TEX>$10wt.\%$</TEX> PVDF is used). All the samples showed generally larger charge and discharge capacities when <TEX>$10wt.\%$</TEX> PVDF ratter than <TEX>$5wt.\%$</TEX> PVDF is used. The plateau region in the range of voltage lower than 1.25V becomes larger probably since the structure becomes less disordered by the addition of boron. When boron is added, the charge and discharge capacities decreased suddenly at the second cycle. This may be become only a part of Li are reversibly deintercalated and intercalated and a part of Li which are strongly combined with B are not deintercalated. The increases in charge and discharge capacities are considered to be resulted from the increase in the potential of Li in the boron-added carbons, caused by the strengthening of the chemical bond between the intercalated Li and the boron-carbon host since the boron acts as electron acceptor.

Lithium-Ion-Batteries (LIB) are promising energy storages for electric vehicles. However, to realize driving ranges equivalent to ones of combustion engine powered vehicles they need increased energy densities. The cells could be improved by targeted optimization of certain electrode manufacturing parameters, such as the calendering strenght [1]. However, calendering does not only affect several geometric parameters of the LIB, like electrode thicknesses and porosities, which can experimentally be measured. Calendering also affects other parameters, such as effective electronic and ionic conductivities and solid-liquid interfacial area [1, 2, 3]. The latter parameters are more complex to investigate, and exclusively experimental or exclusively simulative methods which are usually applied, are not sufficient to understand calendering impacts on battery behaviour. Therefore in this talk we introduce a method where experimental and simulative methods are combined in order to improve the understanding of calendering impact. For the experimental part inhouse made pouchcells were examined by measuring capacity, discharge curves at 1 C and electrochemical impedance spectroscopy. To investigate calendering impacts, these experiments were carried out for cells containing non-calendered NMC cathodes and ones containing cathodes calendered to a degree of 22% compaction of their original thickness, which is a moderate calendering strenght where no electrode damage by compression forces is expected, respectively. Additionally the samples' electronic conductivity was measured. To achieve consistency between experiments and simulation we implemented a pseudo 2D physico-chemical model [4, 5] which enables to simulate discharge curves and Li transport within solid and liquid during discharge process. This model was parameterized [6] and validated to achieve agreement between experimental and simulated discharge curves of pouchcells. Then impacts of the particular parameters affected by calendering (Figure 1 a) were simulated and evaluated stepwisely. With the aid of simulation-supported analysis we show to what extend each of the particular parameters described above contributes to the impact of calendering (Figure 1 b). Geometric parameters, being namely electrode thickness and porosity contribute relatively sparsely, whereas the most significant impact is caused as calendering improves effective ionic and electronic conductivities and solid-liquid interfacial area. Larger electrode thickness and porosity in non-calendered electrodes only cause a slight voltage loss, whereas lower effective ionic and electronic conductivities in non-calendered electrodes not only cause additional voltage losses but also lead to capacity losses due to kinetic limitations of electron and Li ion transport at 1 C discharge. Furthermore the effective solid-liquid interfacial area where electrons and Li ions are exchanged between electrolyte and active material particles appears to be reduced in non-calendered electrodes. Contributions of these effects can be quantified and corresponding cell internal correlations can be analyzed by additionally investigating simulated Li transport within electrolyte and active material. Especially for LIB, where electrodes are complex particle-pore networks with solid and liquid diffusion in addition of binder and carbon black, simulation-supported investigation is shown to yield physically sound correlations between calendering and battery performance. The method presented in this work is also a useful approach to achieve a targeted optimization of LIB manufacturing parameters in order to improve crucial performance quantities, such as the energy density of automotive batteries. Figure 1: Parameters affected by calendering and their particular impacts on electrochemical performance at 1 C [1] W. Haselrieder, S. Ivanov, D. K. Christen, H. Bockholt, A. Kwade, Impact of the Calendering Process on the Interfacial Structure and the Related Electrochemical Performance of Secondary Lithium-Ion Batteries, ECS Transactions 50 (26) (2013) 59-70. [2] G.-F. Yang, S.-K. Joo, Calendering effect on the electrochemical performances of the thick Li-ion battery electrodes using a three dimensional Ni alloy foam current collector, Electrochimica Acta 170 (2015) 263-268. [3] H. Zheng, L. Tan, G. Liu, X. Song, V.S. Battaglia Calendering effects on the physical and electrochemical properties of Li[Ni1/3Mn1/3Co1/3]O2cathode, Journal of Power Sources 208 (2012), 52-57. [4] J. Newman, W. Tiedemann Porous-Electrode Theory with Battery Applications, AlChE Journal 21(1) (1975) 25-41. [5] N. Legrand, S. Rael, B. Knosp, M. Hinaje, P. Desprez, F. Lapicque, Including double-layer capacitance in lithium-ion battery mathematical models, Journal of Power Sources 251 (2014) 370-378. [6] G. Lenze, N. Lin, U. Krewer, Analysis of Parameterization Steps for a Physico-Chemical Lithium-Ion-Battery Model, ModVal12, Freiburg, Germany, Mar. 31 – Apr. 1, 2015. Figure 1