单离子聚合物快离子导体

物 理 化 学 学 报
Acta Phys. -Chim. Sin. 2023, 39 (8), 2205012 (1 of 10)
Received: May 6, 2022; Revised: May 26, 2022; Accepted: May 27, 2022; Published online: June 9, 2022. *
Correspondingauthors.Emails:********************.cn(Y.S.);*******************.cn(L.C.).
This project was supported by the National Key Research and Development Program of China (2021YFB3800300) and the National Natural Science Foundation of China (21733012, 22179143).
国家重点研究发展项目(2021YFB3800300)和国家自然科学基金(21733012, 22179143)资助
© Editorial office of Acta Physico-Chimica Sinica
[Article] doi: 10.3866/PKU.WHXB202205012
A Single-Ion Polymer Superionic Conductor
Guoyong Xue 1,2, Jing Li 2, Junchao Chen 3, Daiqian Chen 2, Chenji Hu 2,3, Lingfei Tang 1,2, Bowen Chen 1,2, Ruowei Yi 2, Yanbin Shen 1,2,*, Liwei Chen 2,3,*
1 School of Nano-Tech and Nano-Bionics, University of Science and Technology of China, Hefei 230026, China.
2 i-Lab, CAS Center for Excellence in Nanoscience, Suzhou Institute of Nano-Tech and Nano-Bionics, Chinese Academy of Science,
Suzhou 215123, Jiangsu Province, China.
3
School of Chemistry and Chemical Engineering, Shanghai Jiaotong University, Shanghai 200240, China.
Abstract: All-solid-state batteries (ASSBs) have been considered a promising candidate for the next-generation electrochemical energy storage because of their high theoretical energy density and inherent safety. Lithium superionic conductors with high lithium-ion transference number and good processability are imperative for the development of practical ASSBs. However, the lithium superionic conductors currently available are predominantly limited to hard ceramics. Practical lithium superionic conductors employing flexible polymers are
yet to be realized. The rigid and brittle nature of inorganic ceramic electrolytes limits their application in high-performance ASSBs. In this study, we demonstrate a novel design of a ternary random copolymer single-ion superionic conductor (SISC) through the radical polymerization of three different organic monomers that uses an anion-trapping borate ester as a crosslinking agent to copolymerize with vinylene carbonate and methyl vinyl sulfone. The proposed SISC contains abundant solvation sites for lithium-ion transport and anion receptors to immobilize the corresponding anions. Furthermore, the copolymerization of the three different monomers results in a low crystallinity and low glass transition temperature, which facilitates superior chain segment motion and results in a small activation energy for lithium-ion transport. The ionic conductivity and lithium-ion transference number of the SISC are 1.29 mS·cm −1 and 0.94 at room temperature, respectively. The SISC exhibits versatile processability and favorable Young’s modulus (3.4 ± 0.4 GPa). The proposed SISC can be integrated into ASSBs through in situ polymerization, which facilitates the formation of suitable electrode/electrolyte contacts. Solid-state symmetric Li||Li cells employing in situ polymerized SISCs show excellent lithium stripping/plating reversibility for more than 1000 h at a current density of 0.25 mA·cm −2. This indicates that the interface between the SISC and lithium metal anode is electrochemically stable. The ASSBs that employ in situ polymerized SISCs coupled with a lithium metal anode and various cathodes, including LiFePO 4, LiCoO 2, and sulfurized polyacrylonitrile (SPAN), exhibit acceptable electrochemical stability, including high rate performance and good cyclability. In particular, the Li||LiFePO 4 ASSBs retained ~ 70% of the discharge capacity when the charge/discharge rate was increased from 1 to 8C . They also demonstrate long-term cycling stability (> 700 cycles at 0.5C rate) at room temperature. A capacity retention of 90% was achieved even at a high rate of 2C after 300 cycles at room temperature. Furthermore, the SISCs have been applied to Li||LiFePO 4 pouch cells and exhibit exceptional flexibility and safety. This work provides a novel design principle for the fabrication of polymer-based superionic conductors and is valuable for the development of practical ambient-temperature ASSBs.
Key Words: All-solid-state lithium metal battery; Solid polymer electrolyte; Superionic conductor;
Single-ion conductor; In situ polymerization; Rate performance
单离子聚合物快离子导体
薛国勇1,2，李静2，陈俊超3，陈代前2，胡晨吉2,3，唐凌飞1,2，陈博文1,2，易若玮2，
沈炎宾1,2,*，陈立桅2,3,*
1中国科学技术大学纳米技术与纳米仿生学院，合肥 230026
2中国科学院苏州纳米技术与纳米仿生研究所，创新实验室卓越纳米科学中心，江苏苏州 215123
3上海交通大学化学化工学院，上海 200240
摘要：具有高锂离子迁移数和良好可加工性能的锂快离子导体对于全固态电池的发展非常重要。

然而，现有的锂快离子导体主要限制于硬质陶瓷，目前尚无柔性聚合物类型的锂快离子导体被报道。

在这个工作中，我们报告了一种通过三种不同有机单体的自由基聚合反应形成的三元无规共聚单离子快离子导体(SISC)。

该SISC中包含丰富的锂离子传输位点和具有阴离子锚定功能的阴离子受体。

此外，三种不同单体的共聚反应带来低结晶度和低玻璃化转变温度(T g)，有利于链段运动，从而获得小的锂离子传输的活化能(E a)。

电化学测试结果表明，该SISC的室温离子电导率和锂离子迁移数分别达到1.29 mS·cm−1和0.94。

将SISC与锂金属负极和多种正极(包括LiFePO4、LiCoO2和硫化聚丙烯腈(SPAN))原位聚合，组装得到的全固态电池具有良好的电化学稳定性。

其中，Li||LiFePO4全固态电池表现出高达8C的倍率性能和良好的循环寿命(在0.5C倍率下稳定循环> 700圈)。

这项工作提供了一种新颖的聚合物基快离子导体设计理念，对于发展高性能全固态电池具有重要意义。

关键词：全固态锂金属电池；聚合物固态电解质；超离子导体；单离子导体；原位聚合；倍率性能
中图分类号：O646
1 Introduction
Replacing flammable organic electrolytes in commercial Li-ion batteries with solid state electrolyte (SSE) to assemble all-solid-state batteries (ASSBs) is believed to be the ultimate solution for addressing the safety issue of Li-ion batteries 1,2. Ideally, SSEs should have a high room-temperature (RT) ionic conductivity and good contact with the electrode materials to ensure a fluent Li-ion transportation through the entire battery, a large lithium-ion transference number (t Li+) to avoid serious ionic concentration polarization during high power charge/discharge, and a wide electrochemical stability window to enable the use of high energy density electrode materials 2–4. However, it is a challenge to develop a SSE to meet all these requirements. Ceramic electrolytes usually have a high RT ionic conductivity, a large t Li+ close to unity, and a wide electrochemical stability window, but they often suffer from large electrode|electrolyte interfacial resistance and poor processibility due to their rigid and brittle nature, restricting their practical application in ASSBs 5,6. By contrast, solid polymer electrolytes (SPEs) are flexible and processable 7. Importantly, they might be prepared by in situ polymerization during battery assembly to ensure a small electrode|electrolyte interfacial resistance 3,8. Nevertheless, most SPEs suffer from low RT ionic conductivity (10−6–10−4S·cm−1) and small lithium-ion transference number (t Li+ < 0.5) 9–12. To obtain a decent electrochemical performance of ASSBs for practical applications, the RT ionic conductivity of the SPE should exceed 10−4 S·cm−1, and the t Li+ should be as close to unity as possible 12. Therefore, over the past few decades, extensive efforts have been devoted to enhancing the RT ionic conductivity and the t Li+ of SPEs 12,13.
The ionic conductivity of SPEs is mainly affected by the number of Li+ solvation sites, the dissociation ability of Li+ of the functional groups, and the polymer chain motion 14–18. There-fore, copolymerization or crosslinking of functional units 3,19,20, nano-filler doping 21–23, and addition of plasticizer 24–26 have been used to increase the ionic conductivity of SPEs 27. On the other hand, organic/inorganic hybrids 28, polyanions 29,30, and anion acceptor-containing polymers 31 have been used to inhibit the movement of anions and boost the Li+ dissolution, obtaining a high t Li+ of SPEs. Progresses have been made in the past years on either the ionic conductivity or the t Li+ of SPEs. For example, Lin et al. 32 developed a poly(vinyl ethylene carbonate) (PVEC) based polymer electrolyte via a polymerization process and obtained a superior RT ionic conductivity of 2.1 × 10−3 S·cm−1, but the t Li+ is at a relatively low value of 0.4. Shin et al. 33 reported a processable single-ion conducting borate polymer consisting of weakly coordinating borate anions connected through butenediol linkers, which shows exceptional selectivity for Li-ion conduction (t Li+ = 0.95), however, the RT ionic conductivity is limited at 1.5 × 10−4 S·cm−1 even in the presence of 30% (mass fraction) polar solvating plasticizer (propylene carbonate). Therefore, it is still an outstanding challenge to develop high quality SPEs with both high RT ionic conductivity (~ 10−3 S·cm−1) and high t Li+ approaching unity.
Herein, we have designed a novel ternary random copolymer
single-ion superionic conductor (SISC), dubbed as P(M-B-V)
(Scheme 1). The P(M-B-V) exhibits both high RT ionic conductivity (~ 10−3 S ·cm −1) and high t Li + approaching unity. It can either be prepared into a free-standing film to ensure mechanical strength or be prepared via in situ polymerization during battery assembly to a low resistant electrode|electrolyte interface, therefore resulting in ASSBs with excellent rate capability and stable cyclability.
2 Experimental section
2.1 Materials
2-hydroxyethyl methacrylate (HEMA; 97%), trimethyl borate
(TMB; 99%), and methyl vinyl sulfone (MVS; 95%) were
purchased from Adamas (Shanghai) and used as received.
Vinylene carbonate (VC; DoDoChem (Suzhou)), 2,2-azobisisobutyronitrile (AIBN; Aladdin (Shanghai)), bis(trifluoromethanesulfonyl) imide lithium (LiTFSI; DoDoChem), anhydrous acetonitrile (Adamas), TF cellulose
membrane (NKK (Japan)), LiFePO 4 (LFP; Sinlion Battery Tech,
Co., Ltd. (Suzhou)), LiCoO 2 (LCO; XTC Corp (Xiamen)), sulfur
powder (S; 99.9%, Sigma-Aldrich (Shanghai)), acrylonitrile
(AN; Aladdin), acetylene black (AB; Alfa Aesar (Tianjin)),
poly(vinyl difluoride) (PVDF; Aladdin), poly(ethylene oxide)
(PEO; M v = 400000, Sigma-Aldrich), N -methyl pyrrolidone
(NMP; 99.5%, Aladdin), and LiTFSI were dried under vacuum
at 100 °C for 48 h, and the TF cellulose film was dried under
vacuum at 80 °C for 24 h and stored in an argon-filled glove box
prior to use.
2.2 Synthesis of boron-based crosslinking agent
(BTETM) 34
A mixture of TM
B (0.52 g, 5 mmol) and HEMA (2.05 g, 15.75
mmol) was dissolved in 10 mL anhydrous acetonitrile, and the
solution was stirred at 50 °C for 3 h in an argon-filled glove box.
Subsequently, the solution was transferred to a reaction device
ventilated with argon, and it was continuously stirred at 70 °C
for another 3 h to remove methanol produced as a by-product.
Finally, the obtained raw product was distilled at 65 °C under
reduced pressure to remove unreacted TMB and residual
acetonitrile, and it was then dried under high vacuum conditions
at 30 °C for 48 h. The resultant dried (boranetriyltris(oxy))
tris(ethane-2,1-diyl) tris(2-methylacrylate) (BTETM) was quickly transferred to an argon-filled glove box for storage to prevent possible hydrolysis. The product was characterized by 1H nuclear magnetic resonance (1H-NMR) and 11B nuclear magnetic resonance (11B-NMR) spectroscopy (as shown in Fig. S1). 1H-NMR [400 MHz, CDCl 3, δ, TMS ref] of BTETM: 6.11 (vinyl, CH), 5.56 (vinyl, CH), 4.23 (CH 2―O ―C(O)), 4.01 ((CH 2)―OB), 1.94 (isobutyl, CH 3). 11B-NMR [400 MHz, CDCl 3, δ, TMS ref] of BTETM: 18.02 (B ―O). 2.3 Preparation of P(M-B-V) SISC A solution mixture containing MVS, BTETM, and VC with a certain mass ratio was evenly mixed under magnetic stirring at room temperature. Then, an appropriate amount of LiTFSI (20% (mass fraction) of total monomers) was dissolved in the above solution. After 2,2-azobisisobutyronitrile (AIBN) (1% (mass fraction) of total monomers) was added and stirred for 1.5 h, the
precursor solution was injected into a cellulose membrane (Φ 18
mm) and followed by heating at 70 °C for 20 h. All these processes were conducted in an argon-filled glove box. 2.4 Preparation of all-solid-state Li|P(M-B-V)|LFP (LCO) battery LFP (LCO), AB, and PVDF powders with a weight ratio of 80 : 10 : 10 were mixed in NMP solvent under vigorous magnetic stirring overnight. Then, the slurry was cast on an aluminum foil and dried at 80 °C for 12 h under a vacuum. The areal loading of the LFP (LCO) cathode was 1.15–1.50 mg ∙cm −2. Subsequently, the electrolyte precursor solution composed of MVS, BTETM, VC, LiTFSI, and AIBN was injected into the cellulose membrane (Φ 18 mm), which was then placed on top of the LFP (LCO) cathode and pasted with Li foil (China Energy Lithium Co. Ltd. (Tianjin)). Finally, the three-layer sandwich structure was sealed in a 2032 coin cell and kept constantly at 70 °C for 20 h. All these processes were conducted in an argon-filled glove box. 2.5 Preparation of all-solid-state Li|P(M-B-V)|SPAN battery The sulfurized polyacrylonitrile (SPAN) cathode was prepared based on our previous work and relevant literature 35,36 with an areal loading of approximately 2.00 mg·cm −2. The battery assembly process was similar to the all-solid-state
Scheme 1 Schematics for P(M-B-V) prepared via free-radical random co-polymerization.
Li|P(M-B-V)|LFP (LCO) battery described above.
2.6 Preparation of all-solid-state Li|P(M-B-V)|LFP
pouch-type battery
The flexible Li||LFP pouch cell with LFP areal loading of ~4.0 mg·cm−2 was fabricated via in situ polymerization similar as above. The size of the cathode electrode was 2.50 cm × 2.00 cm.
2.7 Preparation of Li+ block battery
LFP and PVDF powders with a weight ratio of 80 : 10 were mixed in NMP solvent under vigorous magnetic stirring overnight. Then, the slurry was cast on aluminum foil and dried at 80 °C for 12 h under a vacuum. The areal loading of the LFP cathode was 6.16 mg·cm−2. Next, the electrolyte precursor solution composed of MVS, BTETM, VC, LiTFSI, and AIBN was added to the cathode, which was then placed on top of the cathode and pasted with stainless steel (SS). Finally, the three-layer sandwich structure was sealed in a 2032 coin cell and kept constantly at 70 °C for 20 h. All these processes were conducted in an argon-filled glove box.
2.8 Characterization methods
Nuclear magnetic resonance (NMR) measurements were performed on a Varian Mercury Plus 400 MHz spectrometer and Bruker 400 MHz spectrometer. The morphology and elemental distribution of solid polymer electrolyte (SPE) samples were analyzed through scanning electron microscopy (SEM; FEI Quanta 400 FEG) equipped with energy dispersive X-ray spectrometry (EDX). A Thermo Scientific Nicolet 6700 spectrometer was used to collect Fourier transform infrared (FTIR) spectra of the samples. The crystallinity of the samples was investigated using X-ray diffraction (XRD) measurements on a D8 diffractometer (Bruker) with Cu Kα radiation in the 2θrange of 20°–80°. Thermogravimetric analysis (TGA; TG/ DTA6300) was conducted under a nitrogen atmosphere at a heating rate of 10 °C∙min−1 in the temperature range of 30-600 °C. Differential scanning calorimetry (DSC; DSC6220) analysis was performed in the temperature range of 80 to 80 °C with a heating rate of 10 °C·min−1. The Raman spectra were acquired using a Renishaw in via Qontor confocal Raman microscope using a 785/532 nm wavelength laser. The mechanical properties of the samples were measured by a Universal Material Tester Instron 3365 in the range of 0–10 N. The Young’s modulus was measured by atomic force microscopy (AFM; Asylum Research Cypher S AFM) and AC160TS-R3 tip (Olympus). The Sneddon model was used to fit the force curves. Spots were randomly picked on a selected region, and the value of Young’s modulus within a specific range was measured.
2.9 Electrochemical measurements
The ionic conductivity of the SPE samples was measured by alternating current (AC) impedance on a symmetric SS|SPE|SS cell, where the SS acted as an ion blocking electrode. The symmetrical battery was tested by an electrochemical workstation (the Bio-logic VPM-300) in the frequency range of 0.1 Hz to 7 MHz with an amplitude of 10 mV. The ionic conductivity (σ) was calculated as follows 37:
σ = L
RS
(1) where L is the thickness of the electrolyte membrane, R is the electrolyte resistance, and S is the effective electrode surface area.
The activation energy E a was calculated from the Arrhenius equation (2) 38:
σ = A exp(-E a RT) (2) where A represents the frequency factor, R is the molar gas constant, and T is the absolute temperature.
The Bruce Vincent Evans technique was used to evaluate the lithium ion transference number (t Li+) at room temperature, where a symmetric Li|SPE|Li cell was used to measure the direct current (DC) signal. The current and AC impedance of the battery before and after polarization were tested. According to the direct current polarization method, the t Li+ can be calculated by the following equation 39:
t Li+ =
I s(∆V - I0R0)
I0(∆V - I s R s)(3) where ΔV is constant DC polarization voltage (10 mV), I0 and R0 are current and interfacial resistance at pristine-state, respectively, I s and R s are current and interfacial resistance at steady-state, respectively.
Linear sweep voltammetry (LSV) was used to measure the electrochemical stability of the SPE. A SS|SPE|Li cell was assembled and scanned at a rate of 0.1 mV·s−1 from 0 V to 6 V. The long-term cycling stability between SPE and lithium metal was explored by the galvanostatic cycling test on a symmetric Li|SPE|Li cell with a constant current density of 0.25 mA·cm−2 at 25 °C. Further, the electrochemical performance of Li|P(M-B-V)|LFP (LCO) and Li|P(M-B-V)|SPAN cells were tested via a Neware BTS battery tester.
3 Results and discussion
The BTETM used for preparation of the P(M-B-V) was synthesized through a substitution reaction between trimethyl borate (TMB) and 2-hydroxyethyl methacrylate (HEMA) 34. 1H-NMR and 11B-NMR spectra (Fig. S1a,b, Supporting Information) confirm the structure of synthesized BTETM 34,40. The P(M-B-V) was prepared by heating a mixture containing BTETM, MVS, and VC with a certain mass ratio, a LiTFSI salt (20% (mass fraction)), and a 2,2-azobisisobutyronitrile (AIBN) (1% (mass fraction)) as an initiator at 70 °C for 20 h.
FTIR and 1H-NMR spectra prove the copolymerization of the BTETM, MVS, and VC (Fig. 1a,b, and Fig. S2). As shown in Fig. 1a, the absorption bands of ―C＝C― (~1565 cm−1) and C＝C―H (~ 3166 cm−1) from the VC molecule and the ―C＝C― (~1635 cm−1) from the MVS and the BTETM molecule disappear in the P(M-B-V) spectrum. Instead, a new band emerges at approximately 2976 cm−1, indicating the C＝C double bond from the reactants has transformed to the C―C single bond during the polymerization 41. The peaks at 1715 and
1725 cm −1 are ascribed to the stretching vibration band of the carbonyls in BTETM monomer and P(M-B-V), respectively 42. In addition, the stretching vibration band of the carbonyl in VC (at 1792–1825 cm −1) keeps the same before and after the copolymerization 41. These chemical structure changes during polymerization are also confirmed by 1H NMR spectrum. As shown in Fig. 1b and Fig. S2, the chemical shift of the H in the CH ＝CH at 7.15 (marked as a) of the VC molecule is shifted to 6.08 and 5.72 (a’’ and a’, respectively) after the polymerization, indicating the formation of ―C ―H. Besides, the vinyl groups (marked as c, d and f) of the MVS and the BTETM molecules disappear in P(M-B-V) 41,43.
The P(M-B-V) shows versatile processibility. As shown in Fig. S3, a free-standing film that is flexible and can be readily bent or folded is obtained by copolymerizing the P(M-B-V) monomers in a cellulose membrane framework. The P(M-B-V) in cellulose film shows an average Young’s modulus of 3.4 ± 0.4 GPa (Fig. 1c), which is higher than that of a poly(ethylene oxide) (PEO) SPE (0.5 ± 0.1 GPa, Fig. S4). The maximum tensile stress of the P(M-B-V) in cellulose film is 13.1 MPa, which is significantly greater than that of the cellulose film itself 10.7 MPa (Fig. 1d). Scanning electron microscopy (SEM) image and energy dispersive X-ray spectroscopy (EDX) characterization on the surface and cross-section of the P(M-B-V) in cellulose film (Fig. 1e and S5) show a uniform elemental distribution of the SPE throughout the cellulose membrane, indicating the favorable wettability of the reactants on the cellulose backbone, which is beneficial for Li + transportation through interconnected channels of cellulose. Last but not the least, thermogravimetric analysis (TGA) results show that the thermal degradation of P(M-B-V) occurs at ~ 266 °C (Fig. S6a,b), reflecting its excellent thermal stability.
The RT ionic conductivity of the P(M-B-V) with different monomer ratios was investigated and optimized (Fig. S7a, b and Fig. S8a,b, Table S1). The results reveal that the P(M-B-V) exhibits the highest RT ionic conductivity of 1.29 × 10−3 S ∙cm −1 when the mass ratio of BTETM : MVS : VC = 2 : 1 : 7 (the molar ratio of BTETM : MVS : VC = 1 : 1.8 : 16, Fig. 2a). This ionic conductivity value is much greater than that of P(VC) (6.91 × 10−6 S·cm −1, Fig. S9) and PEO SPEs (~10−5 S·cm −1) 44. The temperature dependence of the ionic conductivity follows the Arrhenius relationship (Fig. S8a,b). The lowest activation energy for Li + diffusion 0.15 eV occurs at the optimized BTETM: MVS: VC ratio of 2:1:7, which is also much lower than that of the P(VC) (0.53 eV , Fig. S10) and the PEO SPEs (~0.60 eV) 44. X-ray diffraction (XRD) and differential scanning calorimetry (DSC) characterization reveal that the high RT ionic conductivity and low activation energy of the P(M-B-V) could
Fig. 1 (a) FTIR spectra of the VC, MVS, BTETM, and P(M-B-V). (b) 1
H-NMR spectra of the VC, MVS, BTETM, and P(M-B-V). (c) AFM Young’s modulus mapping of the P(M-B-V). (d) Stress-strain curves of the cellulose membrane and the P(M-B-V) in cellulose film.
(e) Top-view SEM image of P(M-B-V) and the corresponding EDX elemental mapping (B, F, and S).
be attributed to its amorphous nature and low T g . As shown in Fig. 2b, the P(M-B-V) shows an amorphous structure due to the random copolymerization of three monomers. As a result, the T g of the P(M-B-V) appears ~ 6.36 °C (Fig. 2c) lower than that of the P(VC) (~ 13.58 °C, Fig. S11), indicating that P(M-B-V) may have superior chain segment mobility and lower activation energy for Li + diffusion 45. Furthermore, the moderate interactions between the Li + with the C ＝O, C ―O ―C, and S ＝O groups in P(M-B-V) might also be important for improving the ionic conductivity 41,43,46.
Impressively, the t Li + of the P(M-B-V) with the optimized monomer ratio is close to unity (t Li + = 0.94, Fig. 2d and Table S1), which is much higher than that of the P(VC) (t Li + = 0.48, Fig. S12a and Table S2) and the commercial liquid electrolyte (t Li + = 0.36, Fig. S12b and Table S2). This could be attributed to that the BTETM monomer can act as an anion-acceptor in the random copolymer. Specifically, the sp 2 hybridized boron atom of the BTETM can serve as a Lewis acid, which possesses a vacant orbital to interact with TFSI - ions via Lewis acid-base interaction, thereby limiting the movement of the TFSI - ions and enhancing the t Li + 34. Evidently, the t Li + value of a commercial liquid electrolyte is also significantly improved from 0.36 to 0.60 when 20% (mass fraction) of BTETM is added (Fig. S12c and Table S2). We found that the t Li + of a commercial liquid electrolyte could also increase from 0.36 to 0.48 when 10% (mass fraction) of MVS is added (Fig. S12d and Table S2), which could be due to an increase of free Li + originated from the strong interaction between the LiTFSI and the S ＝O group of the MVS 46,47. It is worth noting that, the RT ionic conductivity and t Li + of the P(M-B-V) are superior to those reported in recent works, as shown in Fig. 2e.
Raman spectroscopy characterization shows that the complete dissociation of the LiTFSI salt in the copolymer contributes to the superior electrochemical performance of the P(M-B-V) SISC. Raman spectra in Fig. 2f show that characteristic peaks of Li +-coordinated TFSI - and Li +-uncoordinated TFSI - are observed at 748 cm −1 in LiTFSI solid and 742 cm −1 in LiTFSI dissolved in liquid VC, respectively 8,48. The spectra of LiTFSI in other solids indicate that LiTFSI is completely dissociated in the P(BTETM), P(MVS), and the optimized P(M-B-V) but not in P(VC), corroborating that BTETM and MVS molecules dissociate LiTFSI salt and contribute to the high ionic conductivity and t Li + in P(M-B-V).
Linear sweep voltammetry (LSV) measurements show a reasonable electrochemical stability window of the P(M-B-V). The LSV curve of the asymmetrical Li|P(M-B-V)|SS cell (Fig. S13) shows that P(M-B-V) is electrochemically stable in the range of 0.25–4.68 V (vs. Li +/Li), indicating that it can be coupled with high-voltage cathodes.
Symmetric Li|P(M-B-V)|Li cells were assembled to investigate the lithium plating/stripping behavior at the Li|P(M-B-V) interface. Fig. S14 shows the cycling performance of the symmetric Li|P(M-B-V)|Li and Li|P(VC)|Li cells at a current density of 0.25 mA ∙cm −2. The Li|P(M-B-V)|Li cell presents
Fig. 2 (a) Nyquist plot of P(M-B-V) measured at RT. (b) XRD patterns of the P(VC), P(BTETM), P(MVS), and P(M-B-V) samples.
(c) DSC curve of the P(M-B-V). (d) The current-time plot of a Li|P(M-B-V)|Li symmetric cell and Nyquist plots before and after polarization (inset). (e) The comparison of RT ionic conductivity and Li + transference number of the P(M-B-V) SISC against
recent reports. (f) Raman spectra of pure solid LiTFSI and LiTFSI dissolved in VC, P(VC), P(MVS), P(BTETM), and P(M-B-V).
excellent long-term stability beyond 1000 h with flat lithium plating/stripping curves and a small polarization potential of ~ 11 mV , while the Li|P(VC)|Li cells exhibit a high polarization potential of ~ 73 mV and are short circuited after cycling for 50 h, indicating that the interface between the P(M-B-V) SISC and the Li anode is electrochemically stable.
To further explore the application of the P(M-B-V) SISC in full batteries, it was in situ polymerized to assemble ASSBs using Li metal anodes coupled with different cathodes (LiFePO 4 (LFP), LiCoO 2 (LCO), and sulfurized polyacrylonitrile (SPAN)). Before the assembly of the ASSB, a solid LFP cathode without electron-conducting agent acetylene black was prepared with in situ polymerized P(M-B-V) and its RT ionic conductivity was measured to be 9.16 × 10−4 S·cm −1 (Fig. S15), indicating that in situ polymerized P(M-B-V) SISC can effectively construct Li + conducting network in solid state cathodes.
Li|P(M-B-V)|LFP batteries exhibit excellent rate capacity due to favorable interfacial contact, excellent ionic conductivity, and large t Li + of the P(M-B-V) SISC (Fig. 3a, b). Even under a high
discharging current density of 2.24 mA·cm −2 (8C , 1C = 170 mAh·g −1, charging at 0.5C ), the discharge capacity is still maintained at 96.48 mAh·g −1, which is comparable to that of commercial liquid electrolyte (Fig. S16). In addition, the battery exhibits long cycle life. A capacity retention of 90% and an average voltage of ~ 3.34 V were obtained after 630 cycles at 0.5C (Fig. 3c). The corresponding charge-discharge curves show a small and stable polarization during cycling (Fig. 3d). Similar results were obtained when P(M-B-V) SISC-based Li||LFP batteries were cycled at 0.7C (Fig. S17and S18). Such rate and cycling performance of Li|P(M-B-V)|LFP cells compares more favorably to those of Li|P(VC)|LFP cells (Fig. S17). Even at a high rate of 2C , a high initial discharge capacity of 128.66 mAh·g −1 and a capacity retention of 90% after 300 cycles were obtained (Fig. 3e, f). In addition, Li|P(M-B-V)|LFP cells with cathode areal loading as high as 4.6 mg ∙cm −2 have also been assembled. As shown in Fig. S19, Li|P(M-B-V)|LFP cell with high cathode areal loading also exhibits good cycle stability. A capacity retention of 86% is obtained after 140 cycles at 0.2C .
Fig. 3 (a, b) Rate performance and representative charge-discharge voltage profiles of the Li|P(M-B-V)|LFP battery charged at 0.5C and discharged at different C rates. Room temperature cycling performance and voltage profile evolution of the Li|P(M-B-V)|LFP battery at (c, d) 0.5C and (e, f) 2C
.。