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Article

Highly Conductive Single-Ion Polymeric Electrolyte for Long-Cycle-Life Lithium Metal Batteries

by
Yuying Yang
,
Yabin Zhang
,
Yuxin Song
,
Tingbin Ma
,
Luqing Zhang
* and
Shuxiang Zhang
Shandong Provincial Key Laboratory of Fluorine Chemistry and Chemical Materials, School of Chemistry and Chemical Engineering, University of Jinan, Jinan 264200, China
*
Author to whom correspondence should be addressed.
Energies 2024, 17(14), 3398; https://doi.org/10.3390/en17143398
Submission received: 20 June 2024 / Revised: 3 July 2024 / Accepted: 5 July 2024 / Published: 11 July 2024
(This article belongs to the Section D2: Electrochem: Batteries, Fuel Cells, Capacitors)
Browse Figures
Versions Notes

Abstract

:
Considerable research has been conducted on single-ion conductive polymeric electrolytes with high lithium ion transference numbers. However, low ionic conductivity is a long-standing challenge for lithium metal batteries, hindering the development of extending their cycle life. In this study, we synthesized a novel fluorine-containing single-ion polymeric electrolyte, LiP(VDF-co-MAF)BB (Polyvinylidene fluoride trifluoromethyl acrylate lithium borate polymer; subsequently referred to as PPMBB), exhibiting a room temperature conductivity of 1.03 × 10−3 S/cm. This electrolyte demonstrates a high lithium ion transference number of 0.7901 and an extended electrochemical stability window of 5.5 V. Under a 2 C discharge rate, it manifests a remarkable discharge specific capacity of 146.8 mAh/g. Moreover, even after 364 cycles, the capacity retention remains at 76%. The single-ion polymeric gel electrolyte designed in this work provides a promising strategy for the prolonged cycling performance of lithium metal batteries.

1. Introduction

In response to the imminent challenges of the greenhouse effect and diminishing fossil energy resources [1,2,3,4], lithium metal batteries have gained significant attention due to their inherent advantages, including high energy density (3860 mAh·g−1), low electrochemical potential (−3.04 V vs. standard hydrogen electrode), long lifespan, and environmental safety. In recent years, they have been widely adopted in electronic portable devices, wearable devices, new energy vehicles, and so on [5,6,7]. Despite their notable merits, the commonly used liquid electrolytes face substantial difficulties, such as the uncontrolled growth of lithium dendrites, membrane penetration, liquid electrolyte leakage, short circuits, and even safety issues like fire and explosions [8,9,10]. Effectively addressing these challenging issues has sparked widespread enthusiasm within the scientific community.
Gel-polymer electrolytes are emerging as a promising substitute for organic liquid electrolytes, attracting considerable attention owing to their high safety, good interface compatibility, excellent flexibility and mechanical properties, and significant potential for achieving high energy density storage [11,12,13]. To fully exploit the advantages of gel-polymer electrolytes, several design factors must be considered, including the following: (1) high conductivity and lithium ion migration to enhance cycling performance [14,15]; (2) good interface compatibility to undesirable side reactions; (3) excellent mechanical properties to inhibit dendrite growth; and (4) stable performance to ensure safety under demanding operating conditions [16,17,18,19].
Current research primarily focuses on single-ion-conductive gel-polymer electrolytes, wherein anions are immobilized on the polymer main chain or within an inorganic framework. This encompasses diverse materials that include carboxylic acid-based polymers [20,21], sulfonic acid-based polymers [22,23,24], sulfone imide-based polymers [25,26,27,28], anion-centered around boron [29,30,31,32,33], and anion-centered around phosphorus [34]. For instance, Huo et al. reported an iodine-driven strategy, introducing negatively charged iodine groups onto the polymer chain to effectively attract Li+ ions, facilitating Li+ transport, and promoting the dissociation of Li salts. The resulting iodine-based single-ion-conductive polymer electrolyte (IPE) exhibited a high ionic conductivity of 0.93 mS·cm−1 and tLi+ up to 0.86 at 25 °C [35]. Wang et al. report a self-crosslinked polymer electrolyte (SCPE) for high-performance lithium batteries, prepared using a one-step method based on 3-methoxysilyl-terminated polypropylene glycol (SPPG, a liquid oligomer). It is worth noting that lithium bis(oxalate)borate (LiBOB) can react with SPPG to form a crosslinked structure via a curing reaction. This self-formed polymer electrolyte exhibits excellent properties, including high room temperature ionic conductivity (2.6 × 10−4 S/cm−1), a wide electrochemical window (4.7 V), and a high Li ion transference number (0.65) [36]. However, a significant challenge faced by single-ion-conductive gel-polymer electrolytes lies in balancing the high lithium ion migration number with the overall low ion conductivity.
Research has found that boron (B) atoms possess strong electron-withdrawing abilities, effectively dispersing charges, weakening coordination with anions, favoring charge dissociation, and exhibiting good thermal stability and electrochemical performance. Wang et al. discovered that the LiBF4 additive not only improved the stability of high-voltage NCM811 cathodes but also facilitated the dissolution of LiNO3 in carbonate electrolytes through its Lewis acidity. This aided the structural stability of LiNi0.8Co0.1Mn0.1O2 (NCM811) at 4.4 V, enabling fast charging [37]. Zhu et al. prepared a novel single-ion polymer electrolyte, LiPAAOB, using varying proportions of polypropylene, boric acid, lithium hydroxide, and oxalic acid. With an ambient temperature ionic conductivity of 2.3 × 10−6 S·cm−1 and a stable electrochemical window up to 7.0 V (vs. Li+/Li), as well as LiPVAOB using different ratios of polyvinyl alcohol, boric acid, lithium hydroxide, and oxalic acid with an ambient temperature ionic conductivity of 6.11 × 10−6 S·cm−1 and a stable electrochemical window up to 7.0 V (vs. Li+/Li), both were suitable for high-energy-density 5 V lithium-ion batteries [38,39].
Inspired by the examples mentioned above, researchers propose a novel single-ion-conductive gel-polymer electrolyte centered around boron atoms. The crosslinked network formed by crosslinking boron atoms and oxalic acid integrates the advantages of single-ion conduction while maintaining a high conductivity, enhancing cycling stability. The semi-crystalline nature of polyvinylidene fluoride-hexafluoropropylene (P(VDF-co-HFP)) increases the amorphous region, facilitating ion migration [40,41,42]. Therefore, we chose P(VDF-co-HFP) as the base, and in terms of thermal stability, it plays a significant role. LiDFOB was chosen as the lithium salt due to its excellent performance, high solubility, suitable temperature range, and structural similarity to the main components. Sun et al. found that LiDFOB’s oxalic acid portion combined with the simultaneously generated LiF resulted in uniformly distributed nanostructured LiF particles. The presence of nanostructured LiF in the solid electrolyte interface (SEI) led to a uniform diffusion field gradient on the lithium electrode, improving cycling performance [43].
Therefore, we employed a one-step method to prepare a novel single-ion polymer electrolyte, LiP(VDF-co-MAF)BB oxalate, with a physically crosslinked P(VDF-co-HFP) polymer matrix to enhance material mechanical properties. LiP(VDF-co-MAF)BB effectively restrains the movement of boron anions, enhances lithium ion migration, and introduces oxalic acid functional groups and existing EO functional groups, reduces material crystallinity, improves segment flexibility and promotes overall ion migration. The combination of these factors significantly enhances the overall performance of the material. The single-ion conductive gel electrolyte prepared in this study exhibits a migration number of 0.7901, and an ion conductivity of 1.03 × 10−3 S·cm−1 at room temperature. At a 2 C rate, it provides a high discharge specific capacity of 146.8 mAh/g. Moreover, after 364 cycles, the capacity retention reaches 76%.

2. Experimental Section

2.1. Materials

Poly(vinylidene fluoride-co-hexafluoropropylene) (P(VDF-co-HFP)) with a molecular weight of 400,000 (supplier: SOLVAY 21510), Poly(vinylidene fluoride-co-methyl acrylate) (P(VDF-co-MAF)) synthesized (specific synthesis steps and characterization are in supporting information), and lithium carbonate (AR, 99.5%) were obtained from Shanghai Maclin Biochemical Technology Co., Ltd.; (Shanghai, China). Boric acid (GR (Hu trial), ≥99.8%), and oxalic acid (GR (Hu trial), ≥99.8%) were purchased from China National Pharmaceutical Group Chemical Reagent Co., Ltd.; (Shanghai, China). Acetonitrile (GR, ≥99.7%) was obtained from China National Pharmaceutical Group Chemical Reagent Co., Ltd.; (Shanghai, China). and purchased from Shanghai Maclin Biochemical Technology Co., Ltd.; (Shanghai, China). N,N-dimethylformamide (AR (Hu trial), ≥99.5%) and Dimethyl carbonate (AR (Hu trial), ≥99.5%) were obtained from China National Pharmaceutical Group Chemical Reagent Co., Ltd.; (Shanghai, China). The electrolyte (1 mol/L lithium difluorooxalate borate/ethylene carbonate + dimethyl carbonate called KLD-LD05) came from Keloude Experimental Equipment Technology Co., Ltd.; (Shenzhen, China). The lithium iron phosphate, acetylene black, and polyvinylidene fluoride were all purchased from Shenzhen Kejing Co., Ltd.; (Shenzhen, China), and N-methyl-2-pyrrolidone (99.9%) was bought from KRD Experimental Equipment Technology Co., Ltd.; (Shenzhen, China).

2.2. Characterization and Measurements

Fourier-transform infrared spectroscopy (FTIR) was used, ranging from 4000 cm−1 to 400 cm−1 using a NICOLET iS10 instrument. (Thermo Fisher Scientific (China) Technology Co., Ltd.; Shanghai, China). 19F, 11B, and 7Li nuclear magnetic resonance spectra were obtained using German Bruker Avance III HD 500 MHz and 600 MHz spectrometers(Brooke (Beijing) Technology Co., Ltd.; Beijing, China) with DMF-d7 as the solvent. X-ray diffraction (XRD) analysis was performed using a German Bruker D8 Advance instrument (Bruker Corporation; Billerica; America) at angles from 5 to 70° with a scan speed of 5°/min.
The morphology of the membranes and the morphology of the Li anode after cycling were observed using scanning electron microscopy (SEM) (HITACHI 8100 model) (Hitachi Scientific Instruments (Beijing) Co., Ltd.; Beijing, China) with an accelerating voltage of 15 kV. Elemental distribution was determined by energy-dispersive X-ray spectroscopy (EDS) (Hitachi Scientific Instruments (Beijing) Co., Ltd.; Beijing, China). Tensile tests were conducted using an MTS Industrial Systems SANS instrument (Meters Industrial Systems (China) Co., Ltd.; Shenzhen, China) with a stretching speed of 1 mm/min. Thermogravimetric analysis (TGA) was carried out using a NETZSCH TG 209F3 instrument (NETZSCH Scientific Instruments Trading (Shanghai) Ltd.; Shanghai, China) in the temperature range of 35 to 600 °C at a heating rate of 10 °C/min under N2 atmosphere. Differential scanning calorimetry (DSC) tests were performed using a US TA Instruments DSC25 instrument (Watsch Technology (Shanghai) Co., Ltd.; Shanghai, China) in the temperature range of −80 to 300 °C at a heating rate of 10 °C/min under N2 atmosphere.

2.3. Electrochemical Performance Testing

Electrochemical impedance spectroscopy (EIS) was conducted using a Bio-Logic electrochemical workstation at frequencies ranging from 0.1 to 106 Hz and temperatures from 30 to 70 °C. Initially, electrolyte membranes with a diameter of 14 mm were assembled into CR2032 symmetrical blocking cells for testing. The ionic conductivity (σ) was calculated using the formula: σ = L / SR , where σ is the ionic conductivity, L is the membrane thickness, R is the electrolyte impedance obtained from the impedance spectrum, and S is the effective area of the electrode (S = 0.071 cm2 in this experiment).
The lithium ion transference number (tLi+) was measured using the galvanostatic intermittent titration technique (GITT) and EIS. The cells were assembled into Li symmetrical non-blocking CR2032 coin cells. The lithium ion transference number (tLi+) was calculated by the formula: t Li + = I S   ( Δ V   –   I 0 R 0 ) / I 0 ( Δ V   –   I S R S ) , where I0 and IS are the initial and steady-state currents, and R0, RS are the impedances measured by EIS. The polarization voltage applied to the cell was ΔV = 10 mV.
The electrochemical stability window was tested at room temperature using linear sweep voltammetry (LSV) with a scan rate of 0.1 mV/s and a working voltage range of 0.5–6 V. Oxidation stability was tested at room temperature using cyclic voltammetry (CV) with a test voltage range of 2.5–4.2 V and a scan rate of 0.1 mV/s for 4 cycles. The testing system used Li//LFP CR2032 cells.
To verify the electrochemical cycling stability of the Gel-Polymer Electrolytes (GPEs), Li//Li symmetrical cells and Li//LFP cells were assembled and tested using the Neware Battery Testing System (CT-3008). The electrode slurry was prepared with a mass ratio of LFP–acetylene black–PVDF at 8:1:1, and NMP was added dropwise, stirred to form a viscous slurry, coated on aluminum foil, and dried in a vacuum oven at 80 °C for 12 h. The mass of the active material on LFP was approximately 2.1 mg/cm3. At room temperature, the charge and discharge time for Li symmetrical cells were both one hour, and the voltage range for charge and discharge performance tests was 2.5–4.2 V. All cells were assembled in a glovebox (O2 ≤ 0.5 ppm, H2O ≤ 0.5 ppm).

3. Results and Discussion

3.1. Structural Characteristics of LiP(VDF-co-MAF)BB and PPMBBn

The single-ion conductive gel polymer network LiP(VDF-co-MAF)BB, centered around anionic moiety B, was prepared using a one-step method (Figure 1a). The specific preparation process and ratio are shown in Table S1. From the 11B NMR spectrum shown in Figure 1b, the absorption peak at δ = 6.45 ppm indicates the successful synthesis of a boron-centered crosslinked network in LiP(VDF-co-MAF)BB. The electrolyte membrane, as indicated by Figure S5, after immersion, undergoes a transition from white to transparent and shifts from a solid to a gel-like state, enhancing the compatibility at the interface. Furthermore, it exhibits outstanding flexibility, as demonstrated by its ability to recover to its original state even after bending. This flexibility allows the membrane to accommodate the deformation and expansion of the battery during charging and discharging processes, not only improving the battery’s cycling performance and lifespan but also effectively alleviating internal stresses and reducing the risk of thermal runaway. This, in turn, promotes the stability and safety of the battery system.
The infrared spectrum in Figure 1c reveals absorption peaks at 1404 cm−1, 1171 cm−1, and 879 cm−1, corresponding to the stretching vibrations of -CH2, asymmetric stretching of -CF2, and symmetric stretching of -CF2; there is an absorption peak associated with the α-crystalline phase absorption peak of P(VDF-co-HFP), respectively. While at 840 cm−1, there is an absorption peak associated with the β-crystalline phase of P(VDF-co-HFP) [44,45,46,47]. It is evident that with an increase in doping concentration, the absorption peak intensity of the α-phase decreases, accompanied by an increase in the intensity of the corresponding β-phase absorption peak.
In the infrared spectrum (Figure 1d), the peak observed at 675.4 cm−1 corresponds to the residual DMF molecules binding to Li+ ions, indicating the presence of (Li(DMF)x)+ complexes, while free DMF molecules are absent in the system [48,49]. The persistence of DMF is beneficial, as it facilitates the movement of polymer chain segments and promotes the dissociation of lithium salts in the system. Consequently, the associated activation energy decreases, leading to enhanced conductivity. Peaks at 1075 cm−1 and 840 cm−1 represent the α and β crystalline phases of P(VDF-co-HFP). As doping increases, the intensity of the α-phase absorption peak decreases, and the corresponding β-phase peak intensifies. This suggests reduced crystallinity, increased amorphous regions, and improved Li+ migration, reflected in an enhanced lithium ion transference number. To confirm the uniformity of element distribution, Energy Dispersive X-ray Spectroscopy (EDS) analysis (Figure 1e) and elemental proportion analysis (Figure S6) were conducted. The images show a uniform distribution of elements, indicating excellent dispersion of the electrolyte throughout the system. This uniformity reduces hindrances to lithium ion transport, promotes even lithium ion deposition, and diminishes unnecessary side reactions at the interface, consequently enhancing the subsequent cycling performance.
SEM, crucial for characterizing microstructures and their impact on performance, reveals an asymmetric membrane structure (Figure 2a). The membrane consists of a surface sponge layer with large pores and a central finger-like pore structure [49]. The increased pore size in Figure 2b (compared to PPMBB0, as shown in Figure S7) enhances liquid absorption, promoting ion migration and thereby improving ionic conductivity. Through contact angle testing (Figure S8), we verified the interface wettability between the electrolyte and the electrolyte membrane. The conductivity and lithium ion migration number of PPMBB10 in Figure 3a,b have a higher conductivity and ion migration number, so subsequent tests will mainly focus on PPMBB10. It shows that the contact angle of PPMBB10 (18.39°) is significantly lower than that of PPMBB0 (29.79°). The test results indicate that the contact angle formed by the electrolyte on the surface of the electrolyte membrane suggests a strong affinity between the two. The implies that the electrolyte membrane can effectively interact with the electrolyte, ensuring its optimal performance and stability within the battery system. This validation of interface compatibility provides crucial support for the reliable operation of the battery. The results are also consistent with the conclusions in the infrared image.
The crystallinity of polymers has a significant impact on the conductivity of gel electrolytes [50,51]. Research indicates that the migration of lithium ions primarily relies on the amorphous regions within the polymer. Lower polymer crystallinity results in larger amorphous regions, facilitating easier lithium ion migration and consequently enhancing conductivity. In Figure 2c, the XRD pattern of the prepared PPMBB0–12.5 shows peaks at 2θ = 18.3°, corresponding to the (020) crystal plane of P(VDF-co-HFP)’s α phase and 2θ = 20.4° corresponding to the (200) crystal plane of the β phase [52]. The proportional analysis in Figure S9 reveals an increasing presence of the β phase with higher doping levels of LiP(VDF-co-MAF)BB. This enhancement, coupled with the increased polarity of the electrolyte membrane, promotes the migration of Li+ ions, ultimately improving the system’s conductivity and enhancing battery cycle performance. DSC curves of PPMBB0 and PPMBB10 (Figure 2e) show no significant change in the crystalline melting peak, indicating that the incorporation of LiP(VDF-co-MAF)BB has not altered the performance of the matrix material, which suggests good compatibility with the polymer substrate P(VDF-co-HFP). TGA curves (Figure 2f) display similar thermal stability for PPMBB0 and PPMBB10, affirming their compatibility and uniformity.
The increase in the fracture elongation of gel-polymer electrolytes contributes to preventing dendrite penetration, thereby enhancing the safety performance of batteries. This property imparts greater toughness to the electrolyte, significantly reducing the risk of short circuits or other safety issues that may arise during battery operation due to dendrite formation. Mechanical properties analyzed in Figure 2d reveal a decrease in stress but an increase in strain with LiP(VDF-co-MAF)BB incorporation. This is attributed to enhanced interactions between flexible segments of P(VDF-co-MAF). Although the uneven distribution of PPMBB12.5 leads to a slight reduction in strain, PPMBB10 still meets practical application conditions [53]. The polymer effectively suppresses the penetration of lithium dendrite during lithium ion transport, thereby prolonging the battery lifespan.

3.2. Electrochemical Performance

To assess the membrane’s practical performance, impedance tests determined the relationship between conductivity and temperature (Figure 3a). As temperature rises, conductivity gradually increases due to accelerated segmental motion, facilitating Li+ migration [54,55]. PPMBB10 exhibits significantly higher conductivity (1.03 × 10−3 S/cm at 25 °C) than PPMBB0 (0.08 × 10−3 S/cm at 25 °C), indicating improved conductivity with LiP(VDF-co-MAF)BB incorporation. However, a higher Li content does not necessarily lead to better performance, as excessive Li content increases viscosity and ion interactions, slowing down the migration of Li+.
Fitting the Arrhenius equation (σ = Ae(−Ea/RT)) to the conductivity can further explore the activation energy (Figure 3b) [56,57]. PPMBB0 exhibits an activation energy of 0.18 eV, while PPMBB10 displays a lower activation energy of 0.12 eV, suggesting that the incorporation of LiP(VDF-co-MAF)BB reduces the ion migration barrier, promoting segmental motion and macroscopically enhances conductivity.
The lithium ion transference numbers for PPMBB0 and PPMBB10 (Figure S10 and Figure 3c) are 0.3370 and 0.7901. Lithium ion transference numbers for PPMBB2.5, PPMBB5, PPMBB10, PPMBB12.5 are 0.37, 0.87, 0.9175, 0.7715 (Figure S11), respectively; on one hand, the increase in the amorphous region facilitates the migration of ions, while on the other hand, the crosslinked network centered around B effectively inhibits the movement of anions. This is advantageous for the easier migration of Li+ between electrodes, reducing concentration polarization, promoting uniform Li+ deposition, and enhancing cycling stability [58].
The electrochemical stability window at room temperature (Figure 3d) reveals a significant improvement for PPMB2.5–12.5, with windows of 5.13 V, 4.78 V, 5.14 V, 5.5 V, and 4.79 V, respectively. These values are higher than the original PPMBB0 window of 4.74 V, indicating enhanced oxidative stability. The introduction of LiP(VDF-co-MAF)BB improves the electrochemical stability window, reduces secondary reactions on the lithium metal surface, minimizes non-active lithium generation, and thereby enhances the electrochemical stability and lifespan of the battery.

3.3. Anode Interface Stability

Figure 4a illustrates the discharge capacity and Coulombic efficiency of Li/PPMBB0-GPEs/LiFeO4 and Li/PPMBB10-GPEs/LiFeO4 at different rates of room temperature. At the same current density, Li/PPMBB10-GPEs/LiFeO4 consistently exhibits higher discharge capacity than Li/PPMBB0-GPEs/LiFeO4. The initial discharge specific capacity of Li/PPMBB10-GPEs/LiFeO4 is 160.16 mAh/g, and after high-rate cycling at 1 C and 2 C, the discharge capacity remains at 136.37 mAh/g, maintaining a Coulombic efficiency of around 100%. Upon returning to 0.1 C, the capacity recovers to its initial value. In contrast, Li/PPMBB0-GPEs/LiFeO4 experiences significant capacity decay at 2 C, dropping to 95.98 mAh/g, with fluctuating Coulombic efficiency. The corresponding charge–discharge curves in Figure 4b,c reveal Li/PPMBB10-GPEs/LiFeO4 at a nearly constant discharge specific capacity, minimal overpotential, and stable charge–discharge platforms at different current densities. This behavior is attributed to the improved ionic conductivity and lithium ion transference number of PPMBB10, reducing ion migration barriers, suppressing interface side reactions, and lowering overpotential, thus enhancing cycling stability.
Due to PPMBB10 having high conductivity and a high transference number, the cycling stability of the cell at rates of 1 C and 2 C within the voltage range of 2.5–4.2 V was studied (Figure 5a). At room temperature, Li/PPMBB10-GPEs/LiFeO4 exhibits an initial discharge specific capacity of 146.8 mAh/g at 1 C, maintaining stable performance over 1000 cycles with a capacity retention rate of 76% and a Coulombic efficiency close to 100%. At 2 C, Li/PPMBB10-GPEs/LiFeO4 shows an initial discharge specific capacity of 151.5 mAh/g, maintaining stability over 364 cycles with a capacity retention rate of 84% and a Coulombic efficiency also close to 100%. In comparison, Li/PPMBB0-GPEs/LiFeO4 shows higher initial discharge capacities at both 1 C and 2 C but exhibits significant capacity decay and fluctuating Coulombic efficiency with cycling (Figure S12). This indicates uneven lithium ion deposition, likely due to hindered lithium ion transport caused by freely moving anions in the system. The charge–discharge curves also support the superior cycling stability of PPMBB10.
To further confirm the role of the polymer electrolyte PPMBB membrane in interface deposition stability, SEM images of the lithium metal surface after cycling are presented in Figure 6a,c. The unmodified PPMBB0 exhibits uneven lithium deposition with large particles, leading to a reduced active lithium content, severe capacity decay, and a non-uniform interface. In contrast, PPMBB10 effectively inhibits lithium ion migration, promoting uniform lithium ion deposition and forming a stable deposition interface. In contrast, as seen in Figure 6b,d, the modified polymer electrolyte membrane enhances interface stability.
In the Li 1s spectrum (Figure 6e), the presence of LiF (56.1 eV), LiCO3 (55.4 eV), and LiO2 (58.9 eV) is observed. In the B 1s spectrum (Figure 6f), B-O (191.8 eV) and B-F (192.2 eV) are present, along with LiF (684.5 eV) and B-F (685.4 eV) appearing in the F 1s spectrum (Figure 6h) [59]. These results suggest the involvement of lithium salt LiDFOB in constructing an SEI layer containing inorganic LiF on the surface of the lithium metal. This effectively isolates electrons, providing protection at the interface, enhancing the dynamics of lithium ion transport, and reducing side reactions at the lithium metal interface. In the C 1s spectrum (Figure 6g), peaks such as C-C (284.8 eV), C-O (286.5 eV), and C=O (288 eV) mainly arise from solvent decomposition. The appearance of these peaks signifies the presence of an organic–inorganic hybrid SEI, reducing the loss of active lithium metal, promoting interface stability, and thus increasing cycling stability.

3.4. 7Li NMR Analysis

Figure 7a shows 7Li NMR reflects the solvation structure’s impact on Li+ coordination environment. Comparing 1 M LiDFOB and 1M LiDFOB/EC: DMC, the addition of a solvent results in a lower electron cloud density around Li+ due to solvation, causing a shift towards lower fields. Figure 7b demonstrates that, compared to 1M LiDFOB/EC: DMC, the addition of Li(PVDF-co-MAF)BB in 1M LiDFOB/EC:DMC causes a shift to lower fields. This shift is primarily attributed to the sp3 hybridization of boron, which being electron-deficient in the outer layer, strongly adsorbs positively charged lithium ions. This lowers the electron cloud density around Li+, causing a shift towards lower fields. Moreover, the interaction between the polymer and Li+ promotes lithium salt dissociation, releasing more mobile electrons, ultimately increasing the system’s conductivity.

4. Conclusions

In this study, a boron-centered single-ion-conductive gel-polymer electrolyte was successfully prepared by blending LiP(VDF-co-MAF)BB with P(VDF-co-HFP). The sp3 hybridization of boron atoms, characterized by a deficiency of electrons in the outermost orbital and the presence of lone pair electrons in oxygen, led to strong interactions between boron and oxygen. Consequently, the movement of boron anions in the system was restrained, facilitating the migration of Li+ ions and enhancing the transference number (tLi+). Additionally, the incorporation of fluorine-containing polymers increased the charge delocalization of borosulfate anions. Through Lewis acid-base interactions, a strong binding affinity with DFOB- promoted the dissociation of lithium salts, increasing the concentration of free lithium ions and improving the overall conductivity of the system. While maintaining high conductivity, the transference number was also elevated, promoting uniform lithium deposition at the interface and reducing the generation of dead lithium and lithium dendrites, thus enhancing the cyclic performance of the battery.
The prepared PPMBB electrolyte membrane exhibited remarkable room temperature conductivity (1.03 × 10−3 S/cm) and a high lithium ion transference number (0.7901). In Li//LFP cells, it delivered a substantial discharge capacity of 146.8 mAh/g at a 2 C rate. After 364 cycles, the capacity retention remained at 76%. These results highlight the potential of boron-centered single-ion-conductive gel-polymer electrolytes for future developments in gel-polymer electrolyte research.

Supplementary Materials

The following supporting information can be downloaded at: https://www.mdpi.com/article/10.3390/en17143398/s1, Figure S1: Preparation process diagram of P(VDF-co-MAF); Figure S2: 19F NMR spectrum of polymer P(VDF-co-MAF); Figure S3: Infrared spectra of monomer MAF and polymer P(VDF-co-MAF); Figure S4: Gel Permeation Chromatography Spectra of Polymer P(VDF-co-MAF); Table S1: Specific Casting Liquid Formula for LiP(VDF-co-MAF)BB/P(VDF-co-HFP) Blended Electrolyte Membranes; Figure S5: (a–d) are before and after soaking the electrolyte membrane in electrolyte, as well as the morphology after bending and twisting; Figure S6: The total distribution of elements of the PPMBB10 electrolyte membrane; Figure S7: SEM images of cross-section and local magnification of PPMBB0 electrolyte membrane; Figure S8: Contact angle between PPMBB0 and PPMBB10 electrolyte membranes; Figure S9: Electrolyte membrane PPMBB0-12.5 crystal form proportion; Figure S10: Lithium ion migration number of PPMBB0 electrolyte membrane; Figure S11: Lithium ion migration numbers of electrolyte membrane PPMBB2.5, PPMBB5, PPMBB7.5, and PPMBB12.5; Figure S12: (a) Li/PPMB0-GPEs/Li at constant magnification (1 C), cyclic performance test chart; (b) Li/PPMBB0-GPEs/Li at constant magnification (2 C), cyclic performance test chart.

Author Contributions

Y.Y., L.Z. and S.Z.; Methodology, Y.Y., Y.Z. and S.Z.; Validation, Y.Y.; Formal analysis, Y.Y., Y.Z. and Y.S.; Investigation, Y.Y., Y.Z. and Y.S.; Resources, T.M. and L.Z.; Data curation, Y.Y. and T.M.; Writing—original draft, Y.Y., Y.S. and L.Z.; Writing—review & editing, Y.Z., T.M. and S.Z.; Visualization, L.Z.; Supervision, L.Z.; Project administration, L.Z.; Funding acquisition, L.Z. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by National Key Research and Development Program of China (No. 2022YFE0208200), University of Jinan Disciplinary Cross-Convergence Construction Project 2023 (XKJC-202302 and XKJC-202312).

Data Availability Statement

Data will be made available upon request.

Conflicts of Interest

The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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Figure 1. Preparation of electrolyte membrane PPMBB. (a) Structural diagram of the gel-polymer electrolyte PPMBB. (b) 11B NMR of LiP(VDF-co-MAF)BB. (c) Infrared spectra of the electrolyte membrane PPMBB and (d) local infrared spectra of the PPMBB membrane. (e) Surface EDS diagram of electrolyte membrane PPMBB10 (Li; B; C; F).
Figure 1. Preparation of electrolyte membrane PPMBB. (a) Structural diagram of the gel-polymer electrolyte PPMBB. (b) 11B NMR of LiP(VDF-co-MAF)BB. (c) Infrared spectra of the electrolyte membrane PPMBB and (d) local infrared spectra of the PPMBB membrane. (e) Surface EDS diagram of electrolyte membrane PPMBB10 (Li; B; C; F).
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Figure 2. Morphology and physicochemical properties of electrolyte membrane PPMBB. (a) Cross-sectional SEM image of electrolyte membrane PPMBB10. (b) Local SEM image of electrolyte membrane PPMBB10. (c) XRD image of electrolyte membrane PPMBB0–12.5 membrane. (d) Stress–strain curves of electrolyte membrane PPMBB0–12.5. (e) DSC curves of PPMBB0 and PPMBB10 electrolyte membranes. (f) TGA plots of electrolyte membrane PPMBB0 and PPMBB10.
Figure 2. Morphology and physicochemical properties of electrolyte membrane PPMBB. (a) Cross-sectional SEM image of electrolyte membrane PPMBB10. (b) Local SEM image of electrolyte membrane PPMBB10. (c) XRD image of electrolyte membrane PPMBB0–12.5 membrane. (d) Stress–strain curves of electrolyte membrane PPMBB0–12.5. (e) DSC curves of PPMBB0 and PPMBB10 electrolyte membranes. (f) TGA plots of electrolyte membrane PPMBB0 and PPMBB10.
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Figure 3. Electrochemical performance characteristics of electrolyte membrane PPMBB. (a) Graph of the conductivity of the electrolyte membrane PPMBB0–12.5 as a function of temperature. (b) Arrhenius conductivity plots for PPMBB0 and PPMBB12.5. (c) Lithium ion mobility number plot for PPMBB10. (d) Electrochemical window plot of PPMBB0–12.5.
Figure 3. Electrochemical performance characteristics of electrolyte membrane PPMBB. (a) Graph of the conductivity of the electrolyte membrane PPMBB0–12.5 as a function of temperature. (b) Arrhenius conductivity plots for PPMBB0 and PPMBB12.5. (c) Lithium ion mobility number plot for PPMBB10. (d) Electrochemical window plot of PPMBB0–12.5.
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Figure 4. Half-cell rate cycling stability of electrolyte membrane PPMBB. (a) Li/PPMBB0-GPEs/Li and Li/PPMBB10-GPEs/LFP at different magnifications (0.1 C, 0.2 C, 0.5 C, 1 C, 2 C) cyclic performance test chart. (b,c) are the charge–discharge curves of Li/PPMBB0-GPEs/LFP and Li/PPMBB10-GPEs/LFP at different magnifications (0.1 C, 0.2 C, 0.5 C, 1 C, 2 C), respectively.
Figure 4. Half-cell rate cycling stability of electrolyte membrane PPMBB. (a) Li/PPMBB0-GPEs/Li and Li/PPMBB10-GPEs/LFP at different magnifications (0.1 C, 0.2 C, 0.5 C, 1 C, 2 C) cyclic performance test chart. (b,c) are the charge–discharge curves of Li/PPMBB0-GPEs/LFP and Li/PPMBB10-GPEs/LFP at different magnifications (0.1 C, 0.2 C, 0.5 C, 1 C, 2 C), respectively.
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Figure 5. Half-cell rate cycling stability of electrolyte membrane PPMBB. (a) Li/PPMBB10-GPEs/Li at constant magnification (1 C), cyclic performance test chart. (b) Charge–discharge curves of Li/PPMBB10-GPEs/Li at constant magnification (1 C) cycles. (c) Li/PPMBB10-GPEs/Li at constant magnification (2 C), cyclic performance test chart. (d) Charge–discharge curves of Li/PPMBB10-GPEs/Li at constant rate (2 C) cycles.
Figure 5. Half-cell rate cycling stability of electrolyte membrane PPMBB. (a) Li/PPMBB10-GPEs/Li at constant magnification (1 C), cyclic performance test chart. (b) Charge–discharge curves of Li/PPMBB10-GPEs/Li at constant magnification (1 C) cycles. (c) Li/PPMBB10-GPEs/Li at constant magnification (2 C), cyclic performance test chart. (d) Charge–discharge curves of Li/PPMBB10-GPEs/Li at constant rate (2 C) cycles.
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Figure 6. Surface topography after cycles. (a,b) are SEM images of Li/PPMBB0-GPEs/LFP at 2 C for 200 cycles of lithium metal surface (left) and cross-section, respectively. (c,d) SEM images of Li/PPMBB0-GPEs/LFP at 2 C magnification after 200 cycles of lithium metal surface (left) and cross-section (right), respectively. (e–h) are XPS spectra of Li/PPMBB10-GPEs/LFP on the surface of lithium metal after 200 cycles at 2 c magnification (Li 1 s, B 1 s, C 1 s, F 1 s).
Figure 6. Surface topography after cycles. (a,b) are SEM images of Li/PPMBB0-GPEs/LFP at 2 C for 200 cycles of lithium metal surface (left) and cross-section, respectively. (c,d) SEM images of Li/PPMBB0-GPEs/LFP at 2 C magnification after 200 cycles of lithium metal surface (left) and cross-section (right), respectively. (e–h) are XPS spectra of Li/PPMBB10-GPEs/LFP on the surface of lithium metal after 200 cycles at 2 c magnification (Li 1 s, B 1 s, C 1 s, F 1 s).
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Figure 7. Analysis of the mechanism of action of the electrolyte membrane. (a) NMR plots of 7Li of 1M LiDFOB and 1M LiDFOB/EC + DMC. (b) 7Li NMR plot of LiP(VDF-co-MAF)BB, LiP(VDF-co-MAF)BB/EC+DMC, and LiP(VDF-co-MAF)BB+1M LiDFOB/EC+DMC.
Figure 7. Analysis of the mechanism of action of the electrolyte membrane. (a) NMR plots of 7Li of 1M LiDFOB and 1M LiDFOB/EC + DMC. (b) 7Li NMR plot of LiP(VDF-co-MAF)BB, LiP(VDF-co-MAF)BB/EC+DMC, and LiP(VDF-co-MAF)BB+1M LiDFOB/EC+DMC.
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Yang, Y.; Zhang, Y.; Song, Y.; Ma, T.; Zhang, L.; Zhang, S. Highly Conductive Single-Ion Polymeric Electrolyte for Long-Cycle-Life Lithium Metal Batteries. Energies 2024, 17, 3398. https://doi.org/10.3390/en17143398

AMA Style

Yang Y, Zhang Y, Song Y, Ma T, Zhang L, Zhang S. Highly Conductive Single-Ion Polymeric Electrolyte for Long-Cycle-Life Lithium Metal Batteries. Energies. 2024; 17(14):3398. https://doi.org/10.3390/en17143398

Chicago/Turabian Style

Yang, Yuying, Yabin Zhang, Yuxin Song, Tingbin Ma, Luqing Zhang, and Shuxiang Zhang. 2024. "Highly Conductive Single-Ion Polymeric Electrolyte for Long-Cycle-Life Lithium Metal Batteries" Energies 17, no. 14: 3398. https://doi.org/10.3390/en17143398

APA Style

Yang, Y., Zhang, Y., Song, Y., Ma, T., Zhang, L., & Zhang, S. (2024). Highly Conductive Single-Ion Polymeric Electrolyte for Long-Cycle-Life Lithium Metal Batteries. Energies, 17(14), 3398. https://doi.org/10.3390/en17143398

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