Metalophilic Gel Polymer Electrolyte for in Situ Tailoring Cathode

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Metalophilic gel polymer electrolyte for in-situ tailoring cathode/electrolyte interface of high-nickel oxide cathode in quasi-solid-state Li-ion batteries Yanyun Sun, Yang-Yang Wang, Guoran Li, Sheng Liu, and Xue-Ping Gao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b02440 • Publication Date (Web): 04 Apr 2019 Downloaded from http://pubs.acs.org on April 5, 2019

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ACS Applied Materials & Interfaces

Metalophilic gel polymer electrolyte for in-situ tailoring cathode/electrolyte interface of high-nickel oxide cathode in quasi-solid-state Li-ion batteries Yan-Yun Sun‡, Yang-Yang Wang‡, Guo-Ran Li, Sheng Liu and Xue-Ping Gao* Institute of New Energy Material Chemistry, School of Materials Science and Engineering, Nankai University, Tianjin 300350, China. E-mail: [email protected], Tel/Fax: +86-22-23500876. ‡Yan-Yun Sun and Yang-Yang Wang contribute equally to this work. Keywords: lithium-ion batteries, quasi-solid-state battery, gel polymer electrolyte, high-Ni oxide cathode, electrode/electrolyte interface.

Abstract

High-Ni layered oxides are potential cathodes for high energy Li-ion batteries due to their large specific capacity advantage. However, the fast capacity fade by undesirable structural degradation in liquid electrolyte during long-term cycling is a stumbling block for the commercial application of high-Ni oxides. In this work, a functional gel polymer electrolyte, grafted by sodium alginate, is introduced to increase the stability of high-Ni oxide cathode at the levels of both the particle and electrode. An in-situ generated ion-conducting layer appears on the interface through the chemical interaction between transition-metal cations of the cathode and metalophilic reticulum group in sodium alginate. Such tailoring layer can not only enhance

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the interfacial compatibility on cathode/electrolyte interface, reducing the interfacial resistance, but also inhibit the HF corrosion, suppressing the dissolution of transition-metal cations and harmful gradient distribution of components through the oxide cathode at the electrode level. Meanwhile, detrimental microcracks in oxide microspheres and between primary crystallites are impressively inhibited at the particle level. The high-Ni oxide cathode with the metalophilic gel polymer electrolyte shows excellent cycle stability with large initial capacity of 204.9 mA h g−1 at 1.0 C rate, and high discharge capacity retention within 300 cycles at high temperature. 1. Introduction Li-ion batteries (LIBs) are promising chemical power sources for hybrid electric vehicles (HEVs) and electric vehicles (EVs). One of the key points of fabricating high energy LIBs is to develop cathode with high performance, such as large reversible capacity and excellent cycle stability. The high-Ni layered oxide cathode with high reversible capacity (>200 mA h g−1) and better cost advantage can effectively pave the way to the promotion of high energy LIBs.1-6 However, high-Ni oxides still suffer the detrimental structural evolution at levels of both the particle and electrode, accompanied with rapid capacity fade. The intrinsic Li+/Ni2+ mixing in the bulk is responsible for such structural deterioration, including the phase degradation from layered structure to inactive rock−salt phase or spinel, the formation of microcracks in secondary and 2


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primary particles, and the breakdown of conductive network. Besides, the side reactions at the particle/electrolyte interface further stimulate the structure instability, such as the HF attack accompanied with the dissolution of transition metals (TMs), the formation of solid electrolyte interface (SEI) associated with gradual decomposition of organic electrolyte at the solid/liquid interface.7-12 Particularly, at the electrode level, the cathode surface is more reactive due to different polarization in the porous electrode. The more serious destruction on the cathode surface is induced by high surface activity, leading to the fast capacity fade by the gradient spatial distribution of TMs across the porous electrode during cycling.13-20 Thus, it is significant to enhance the structure stability of high-Ni layered oxide cathode at both the particle and electrode level.

Consistent research efforts have been devoted to improve the structure stability of high-Ni oxide cathode materials, such as ion-doping in the bulk, and surface modification to the primary or secondary particles.21-23 Unfortunately, the interfacial instability of high-Ni oxide cathode is still a trouble in liquid electrolyte (LE). Recently, as compared with the conventional LE, gel polymer electrolytes (GPEs) show the advantage of slowing down the electrode/electrolyte interfacial corrosion for improving the electrochemical cycle stability. In the meantime, such electrode/electrolyte interface is flexible with lower interfacial resistance, as compared to the rigid interface in all-solid-state electrolyte system.24,25 Thus, GPE is the good solution to

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overcome the detrimental factors of high-Ni oxide cathode in the conventional LE.26 However, the aforementioned physical interaction on electrode/electrolyte interface is insufficient to resolve the interfacial degradation in the electrochemical reaction due to the inherent incompatibility on the quasi-solid/solid interface.27-29 Chemical interaction may effectively improve the electrode/electrolyte interfacial compatibility in quasi-solid-state battery. Based on this assumption , it is essential to graft special groups on GPEs in order to increase the interfacial interaction between GPEs and high-Ni oxide cathode. The eco-friendly sodium alginate (SA) is a potential material for LIBs owing to its good ionic conductivity, heat-resistance, and mechanical/electrochemical stability.30-35 More impressively, SA has special metalophilic reticulum groups, which refers to the α-L-guluronate (G) blocks along the linear alginate macromolecules, which are beneficial to immobilizing TMs by chemical coordination.31,36-40 In this respect, it would be a good way to introduce SA into GPEs for enhancing the interfacial compatibility and stability in quasi-solid-state system by a special metalophilic interaction on the cathode/electrolyte interface. In this work, SA is incorporated into poly(vinylidene fluoride-co-hexafluoropropylene)cellulose acetate (PHC) as a polymer matrix to fabricate a metalophilic GPE. The cathode/electrolyte interface stabilization mechanism at levels of both the particle and electrode is proposed based on the chemical interaction between the SA and TMs. Specifically, the chemical chelating interaction between TMs and metalophilic reticulum groups in SA, coupled with the flexible contact between the electrode and GPE, may lead to the in-situ formation of an ion conducting layer on GPE/cathode interface. As expected, the flexible thin layer on the interface with good compatibility could insure the desired electrochemical performance of highNi oxide cathode. 4


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2. Experimental Section 2.1 Preparation of the SA-PHC

The PHC polymer membrane was prepared via a simple electrospinning method. PVdF−HFP (1.05 g) and cellulose acetate (0.45 g) were dissolved in the mixture of N,Ndimethylacetamide (3 mL) and acetone (7 mL) under vigorously stirring. Then, the obtained solution (5 mL) was electrospun at a flow rate of 1 mL h-1 with 20 KV potential difference at 40 °C. The distance between the syringe nozzle and Al drum collector is kept at 15 cm. The collected PHC membrane is dried at 60 °C for 5 hours. At the same time, SA (1.5g) was dissolved in pure water (300 mL) under stirring. To fabricate the SA-PHC membrane, the asprepared PHC was dipped into the mixture of ethanol and water with a volume ratio of 1:5 for 1 min to reduce the surface tension and ensure the following interaction go on smoothly. After that, the obtained humid membrane transferred to the SA aqueous solution immediately and soaked for 6 h. Finally, the SA-PHC membrane was dried at 80 °C for 24 h. The loading of SA in SA-PHC is about 5.5 wt.%. The PHC and SA-PHC GPEs were prepared by soaking in the organic electrolyte with ethylene carbonate (EC) and dimethyl carbonate (DMC) as solvents (volume ratio: 3/7) and LiPF6 as Li salt (concentration:1M) for 2 h. 2.2 Materials characterization

The microstructure and composition of the materials were investigated by using fourier transforming infrared (FT-IR) spectra (BIO-RAD FTS-60) and SEM (JEOL-JSM7800F) equipped with energy-dispersive X-ray spectra (EDS). X-ray photoelectron spectra (XPS, Thermo Scientific ESCALAB 250Xi) was employed to detect the element state of the cathode

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and the membrane, thus identify the surface state of the electrode and GPEs. The thermal stability of the membranes were measured by differential scanning calorimetry (DSC, Mettler Toledo) with a sealed high-pressure DSC pan. The element space distributions in the cathode were analyzed by LIBS (ASI J200) within the area of 2.0×2.0 mm2. Transmission electron microscopy (TEM, JEOL-JEM2800) images and selected area electron diffraction (SAED) patterns were monitored by a charge-coupled device (CCD, Gatan). The stress-strain tests were implemented using a Zwick testing machine with a 10 mm min-1 crosshead speed. 2.3 Electrochemical measurements To prepare the working electrode, the slurry consisting of LiNi0.88Co0.09Al0.03O2 (NCA) (produced by Tianjin Lishen battery Co. Ltd.), conductive agent (Super P) and binder (PVdF) with a mass ratio of 8:1:1 was pasted onto Al foil. After drying at 110 °C overnight, 10 mm round electrodes loaded with active electrrode materials (7−8 mg cm-2) were punched out for assembling the coin-type cells (CR2032). Anode was Li foil with a dimension of 14×0.7 mm. Liquid electrolyte (LE) or the as-prepared functional GPEs was adopted as the electrolyte. A Celgard 2400 porous polypropylene layer was used as the separator for the cell with LE while the ones with GPEs needs no extra separator. The cycle stability and rate capability were detected by a LAND CT-2001A instrument (Wuhan, China) between 3.0−4.3 V (vs. Li/Li+) at room or high temperature (25 or 55 °C). Electrochemical workstation (CHI 600A) was utilized to detect the cyclic voltammograms (CVs) (scan rate: 0.1 mV s−1). Electrochemical impedance spectra (EIS) and the ionic conductivity were measured by electrochemical workstation (Zahner IM6ex) in the frequency range of 10 mHz to 100 kHz with 5 mV disturbance amplitude. The ionic conductivity was tested at 55 °C by the cell, where GPEs or Celgard were placed between

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two identical stainless steel electrodes. The following equation is adopted to calculate the ionic conductivity (σ, s cm-1), σ= L/RS

(1)

where L (cm) is the thickness of the GPE, S (cm-2) corresponds to the effective area of GPE. R (Ω) presents the bulk resistance of GPE, which is always represented by the intersection of the straight line in high frequency and real axis. 41 3. Results and discussion As a polymer matrix of GPEs, the poly(vinylidene fluoride-co-hexafluoropropylene)cellulose acetate (PHC) with high electrolyte uptake42-45 is prepared via a simple electrospinning method.46-48 The SA grafted PHC (SA-PHC) is obtained subsequently by a simple soaking and intermolecular reaction process between the carboxyl in SA and hydroxyl in PHC as shown in Figure 1. Both the PHC and SA-PHC possess 1D fiber morphology with a diameter of about 0.5 µm. There is no obvious difference in morphology except for the relatively uniform and small pore size for SA-PHC. SA is introduced into the polymer matrix evenly according to elemental mappings for selected area, where sodium is uniformly distributed on the PHC matrix. FT-IR spectra further reveals the successful grafting of SA on the PHC surface, as evidenced by the

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occurrence of the new characteristic peak at 1600 cm-1, as well as the peak shifting and intensity variation for some groups of PHC (Figure S1).

Figure 1 SEM images for (a) the as-prepared PHC and (c) SA-PHC membranes. (b) The schematic illustration of introduction mechanism of the SA into PHC. (d) The elemental mapping signals of O, C, F, and Na for SA-PHC membrane obtained by scanning the relative SEM image. To examine the intrinsic performance of polymer membranes, which are crucial for GPEs, the commercial Celgard separator (25 µm in thickness) and PHC, SA-PHC (~30 µm in thickness) are adopted to perform the following tests. As depicted in Figure 2a, after uptaking 10 L LE, PHC involves with the issue of serious shrinkage owing to unsatisfying mechanical strength that is unable to hold the shrinkage stress resulted by the surface tension and plasticizing effect of LE. Thus, unless the thickness of PHC is increased at the cost of energy density, the battery will be short-circuited after uptaking LE. As for the commercial Celgard separator, the wettability of electrolyte is not good enough to quickly wet the film. In contrast, SA-PHC can maintain its original size and shape after absorbing LE, benefiting from the better mechanical/flexible property and electrolyte retention proved by photos of different soaked 8


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membranes after standing for different time in Figure S2. Both the SA-PHC and PHC can be maintained as moist gels after resting for 3 min, while most of the liquid electrolyte evaporates on the Celgard only after 10 seconds. The enhanced mechanical strength of the SA-PHC is further validated by the stress-strain curves for the membranes before and after absorbing LE (Figure 2b). As illustrated, although the mechanical property of the wet membranes are poor as compared with dry ones owing to plasticizing effect of liquid electrolyte, the wet SA-PHC still outperforms PHC, even dry PHC on the mechanical property. The higher tensile strength of 8.9 MPa and relative strain of 102.8% are obtained for the wet SA-PHC, presenting superior toughness to prevent fracture during application. In contrast, the wet PHC has a little better strain (114.3%) but particularly poor tensile strength (2.5 MPa). The SA-PHC not only shows improved mechanical stability, but also realizes enhanced thermal stability. From Figure 2a, we can find that after heat treatment at 150 °C for 30min, the SA-PHC can still keep its original dimension unlike other two membranes, indicating superior thermal stability. DSC analysis further proves the excellent thermal stability of the SA-PHC (Figure 2c) with higher initial exothermic peak at 151.3 °C as compared to PHC and Celgard (149.1 and 132.8 °C, respectively). The enhanced mechanical and thermal stability of SA-PHC benefit from the intermolecular reaction between SA and PHC as depicted in Figure 1, which can strengthen the structure of polymer matrix. Besides, the good thermal and mechanical properties of SA itself also contribute to the enhancement of the mechanical and thermal stability of SA-PHC. Moreover, the polymer matrix stability after electrochemical test for 300 cycles at 55 °C is detected by SEM images of the cycled membranes (Figure 2e-f). As seen, the PHC suffers severe conglutination coupled with elimination of the porous structure, resulted by the dissolved TMs which can weaken the crystallinity and mechanical strength of the membrane during

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cycling.43 Surprisingly, the SA-PHC can maintain its porous structure for uptaking LE, suggesting markedly enhanced structure stability of the polymer matrix with help of the chemical interaction between SA and PHC evidenced by the peak shift in Figure S1. In detail, there is a strong interaction between the hydroxyl in PHC and LE, which can lead to high electrolyte uptake, but aggravate the plasticizing effect and mechanical property of the polymer matrix. The intermolecular interaction between SA and PHC can not only ensure homogenous introduction of the SA with superior mechanical property, but also weaken the interaction between PHC and LE. Furthermore, the crosslinking effect on SA-PHC by the inevitable dissolved Ni2+ during electrochemical cycling also contributes to the improved structure stability of SA-PHC. This mechanism is confirmed by XPS for C 1s for different membranes before and after 300 cycles (Figure S3). Impressively, the C 1s peaks in the PHC is obviously changed after cycling originating from the structure evolution of the polymer matrix during electrochemical cycling, while the SA-PHC can effectively overcome such problem as proved by the similar characteristics on C 1s core level before and after electrochemical test. Additionally, the SAPHC also presents the prominent ionic conductivity, which is critical for the performance of LIBs (Figure 2d, Table S1). In general, the SA-PHC performs better ionic conductivity, mechanical property, thermal stability, and structure stability during electrochemical cycling, which are promising in the application of LIBs.

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Figure 2 Characteristics of the membranes: (a) the pictures of the Celgard, PHC and SA-PHC membranes before application, after uptaking liquid electrolyte (10 µL), and after heat treatment (150 °C for 30min). (b) Stress-strain curves of PHC and SA-PHC membranes before and after uptaking liquid electrolyte. (c) DSC profiles and (d) AC impedance plots of the Celgard, PHC and SA-PHC membranes. SEM images for the (e) PHC and (f) SA-PHC membranes after 300 cycles at 55 °C between 3.0−4.3 V. Cyclic voltammograms (CVs) of the cathodes with different electrolytes are carried out at 25 °C. As indicated in Figure 3a-c, all the samples show three pairs of typical redox peaks,

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which represent three kinds of phase transitions of LiNi0.88Co0.09Al0.03O2 (NCA) as reported.49 As compared to the cathode with LE, the redox peaks of the cathode with PHC move slightly to higher potential with larger potential difference between the oxidation/reduction peaks, and the subsequent peaks vary significantly with increasing cycle number. All these phenomena imply the increased polarization and worse reversibility, which are attributed to the structure instability of PHC or cathode, and large electrode/electrolyte interfacial impedance. Impressively, the cathode with SA-PHC presents relatively better peak coincidence between different cycles. The potential deviation between pairs of redox peaks is smaller, which is indicative of the improved interfacial stability and electrochemical reversibility by grafting SA on PHC. Meanwhile, the stronger and more sharp redox peaks reflect the higher electrochemical reactivity of the cathode with SA-PHC. The good LE wettability and high ionic conductivity of SA-PHC are the main reason for such improvement of the electrochemical activity of the cathode. The electrochemical charge/discharge performance of the high-Ni layered oxide cathodes using LE and GPEs are evaluated with 2032-type coin cells. Long-term cycle performance at 55 °C is displayed in Figure 3d. The cell with SA-PHC approaches a higher initial discharge capacity of 204.9 mA h g−1 at 1.0 C rate and superior capacity retention rate of 68.33% within 300 cycles, considerably outperforming the cells with LE (195.5 mA h g−1 and 36.93%) and with PHC (198.1 mA h g−1 and 49.97%). Further, the decline in discharge midpoint potential of the cathode is notably slowed down by using SA-PHC according to Figure 3e. Figure S4 presents the charge-discharge curves, cycle and rate perfomance of NCA at 25 °C. Similarly, the cell with SA-PHC also exhibits optimized capacity retention and potential stability when operated at 25 °C. And the high-rate performance is enhanced with larger capacity of 159.3 mA h g-1 at 5 C for the cathode with SA-PHC. Therefore, the cell with SA-PHC displays remarkably improved electrochemical

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performance at both normal and tough conditions, which is exceptionally relevant to the cathode/electrolyte interfacial stability.

-1

4.00 V 1st 3.75 V 2nd 4.18 V 3rd

0.4 0.0

Potential (V vs. Li/Li )

Discharge capacity/ mAh g

-1

-0.4

3.65 V 3.93 V4.10 V

3.2 3.6 4.0 4.4 + Potential (V vs. Li/Li )

250

4.08 V

0.4

1st 2nd 3rd

0.4

3.81 V 4.24 V

0.0

0.0

3.70 V 3.97 V4.13 V

-0.4

0.8

25 °C -1

(b) PHC

Current (A g )

0.8

25 °C

Current (A g )

(a) LE

-1

Current (A g )

0.8

+

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60



-0.4

(c) SA-PHC

25 °C

Lower potential 3.97 V 1st 2nd 3rd

3.750 V 4.21 V

3.69 V 3.97 V 4.15 V Higher potential


55 °C

(d)

200 150 100

LE PHC SA-PHC

50 0

0

4.4 (e) 4.2 4.0 3.8 3.6 3.4 3.2 3.0 0

100

Cycle number

200

300

55 °C LE-1st LE-300th PHC-1st PHC-300th SA-PHC-1st SA-PHC-300th

50

100 150 200 -1 Capacity/ mAh g

250

300

Figure 3 Cyclic voltammograms (CVs) of the cathodes with (a) LE, (b) PHC and (c) SA-PHC at 25 °C. (d) The cycle performance and (e) charge−discharge curves of the cathodes with the LE, PHC and SA-PHC in selected cycles at 55 °C (charge: 0.5 C, discharge: 1.0 C).

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Figure 4 SEM images of random selected microspheres on the cathodes cycled with (a) LE, (b) PHC and (c) SA-PHC after different cycles (55 °C, 3.0−4.3 V). As high-temperature structural stability during cycling are critical problems for high-Ni oxide cathode,50-52 in the following section, we focus on the cathode characteristics at both the particle and electrode levels cycled under high operating temperature. At the particle level, microcracks, which are formed because of the volume change during cycling, are indeed the tumbling block in the commercialization of high-Ni oxide cathode materials.53-55 SEM observation of the random selected microspheres is used to probe the effect of different electrolyte system on the microcrack formation at the particle level. Referring to Figure 4, there appears to be microcracks in the secondary spheres on the cathode with LE only after 100 cycles. And the cracks further evolve in the primary crystallites (Figure S5 exhibits the SEM crosssections and enlarged versions of NCA microspheres after different cycles.), finally leading to the entire breakage of secondary spheres. These phenomena may result in the collapse of the

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conductive network at the electrode level as displayed in the SEM images of the cathodes after 300 cycles (Figure S6). The more fresh surface is exposed to the electrolyte for taking part in detrimental side reactions, causing the extra consumption of LE and the formation of highresistive layer. As a comparison, the cathodes from the cells with GPEs present gentle cracking feature even after 300 cycles, especially, the cathode with SA-PHC, benefiting from the stable cathode/electrolyte interface. Phase transition is one of the main cause for intragranular cracks in high-Ni oxide cathode, thus, HR-TEM images and SAED patterns are collected to confirm the function of SA-PHC in restraining the detrimental phase transition and microcracks at the particle level. It is apparent that the cathode cycled with SA-PHC maintains its original crystallographic structure that can be _

matched with space group of R-3m (Figure 5a-b). The blocking on phase transition can be attributed to the metalophilic interaction between SA and TMs. The TMs tend to exist as the high valence state when interacted with electronegative COO- groups in SA to maintain electrical neutrality. In this way, the production of Ni2+ during cycling, which is the origin of phase transition, can be inhibited to a great extent. Figure S7 presents the XPS spectra of C 1s, Ni 2p and Co 2p for the cathodes cycled with different electrolyte. And the above deduction is evidenced by the peak shifting to higher binding energy on Ni 2p core level for the cathode with SA-PHC. In comparison, the structure evolution appears near the surface layer of the cathode cycled with PHC with the thickness of ~6 nm. SAED confirms the existence of spinel and rocksalt structure on the surface layer. And the halo rings with an intensity fluctuation implies poor crystallinity (Figure 5c). More significantly, the surface layer (~25 nm in thickness) appears for the cathode cycled with LE, reflecting more serious lattice distortion during cycling (Figure 5d). In the magnified image (Figure 5e), many edge dislocations can be identified, which are

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nucleation sites of intragranular cracks in primary crystallites.56 Such dislocations could be intrinsic features for the disintegration of primary crystallites (Figure 5f), the pulverization of secondary microspheres (Figure 5g), the crush of conductive network, and the final failure of battery.

Figure 5 HRTEM images and SAED patterns obtained from the region marked by red squares for (a) pristine NCA, and NCA cycled with (b) SA-PHC, (c) PHC and (d) LE after 300 cycles (55 °C, 3.0−4.3 V). (e) The magnified images for site C in (d), and (f) the magnified image and (g) complete picture for the cross section of NCA cycled with LE.

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The function mechanism of the above-mentioned metalophilic interaction between the SAPHC and high-Ni oxide cathode is illustrated by Figure 6a. And vivid representation of the cathode morphology is given by SEM to provide an intuitive understanding of the function of SA-PHC at the electrode level. It should be noted that all the cathodes were soaked and washed three times with dimethyl carbonate (DMC) before analysis. As seen in the cross-sections for the cathodes after 300 cycles at 55 °C (Figure 6b-d), there is no layer formed on the cathode surface from the cell with LE. While the cathode with PHC presents fragmentary layer on the surface owing to the physical contact between the flexible GPE and cathode surface. Differently, a compact layer of 0.8~1.2 µm appears on the cathode surface from the cell with SA-PHC. It benefits from the interfacial chemical interaction and the cross-linking effect of dissolved TMs on SA for enhancing the stability of the special interface layer. In detail, there is a strong interfacial interaction between the TMs on the cathode surface and the metalophilic group in SA as mentioned above, which may result in the formation of the surface layer. Further, the inevitable dissolved TMs can promote the crosslinking of the SA which can enhance the mechanical stability of the surface layer.

37,39

It is demonstrated from the elemental mappings

(Figure 6e-g) that the interface layer is mainly composed of SA-PHC, which has superior property as verified from Figure 1. It is further verified from the strong C 1s signals on the cycled cathodes that the formation of the interface layer is mainly originated to SA-PHC (Figure S7, Figure S8 provides the C 1s spectra of the pristine cathode). As compared with the cathode cycled with LE, the peaks at 289.16 and 286.51 eV are more significant, which are indicative of C=O and C-O bonds in SA. It means the good incorporating of SA-PHC on the cathode/electrolyte interface. This in-situ formed layer can not only provide the interfacial compatibility for effectively decreasing the harmful interfacial resistance in quasi-solid-state

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battery, but also protect the cathode surface from corrosion in cathode-liquid reaction for enhancing the interfacial stability at levels of both the particle and electrode as desired above. Interestingly, the superior interfacial compatibility and electrolyte retention based on SA-PHC can promote the effective use of LE in GPE. Consequently, a small amount of LE in GPE is enough to fulfil the requirement of LIBs with higher-energy density. This feature is confirmed by less liquid utilization per unit active materials for the cell with SA-PHC (Table S2).

Figure 6 (a) Scheme illustration of the functional mechanism of the SA-PHC. The crosssectional SEM images of the cathodes cycled with (b) LE, (c) PHC and (d) SA-PHC after 300th

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cycle (55 °C, 3.0−4.3 V) and the elemental mapping signals of (e) C, (f) F, (g) Ni obtained by scanning the SEM image on (d). The inhomogeneous mass transfer through the porous cathode always occurs in parallel with the electrochemical polarization at the electrode level. On one hand, the more intimate contact of the cathode surface and electrolyte gives the surface priority to the electrochemical reaction. On the other hand, the detrimental parasitic reactions occur on the electrolyte/cathode interface at the same time, such as the dissolution of TMs resulted by the HF erosion.18 Thus, laser induced breakdown spectroscopy (LIBS), which can characterize the spatial distributions of TMs through the depth of the electrodes, is utilized to get an insight into the function of the cathode surface layer induced by SA-PHC. As demonstrated in Figure 7, the distribution of Ni and Co in depth direction of the cathode with first five laser pulses is lively presented (the depth of one laser pulse is ~3µm). The gradient from red to purple corresponds to a gradual decrease in the element content. As predicted, there is a fresh gradient distribution that the concentration of TMs decreases gradually from inside to outside through the porous cathode cycled with LE only after 100 cycles. Especially, the Co concentration changes dramatically, causing damage to the structure stability of NCA. Relative to the electrode with LE, the distribution of TMs through the porous cathode with GPEs is relatively uniform. In details, the Co dissolution near the cathode surface with PHC is restricted to a certain degree within 100 cycles benefiting from the less amount of LE, however, the harmful concentration gradient become obvious with deep cycling to 200 cycles. Amazingly, the flat gradient of the constituent elements for the cathode with SAPHC is shown even after 300 cycles, reflecting the remarkable influence of the in-situ formed protection layer triggered by the interaction between the metalophilic group in SA-PHC and TMs in high-Ni oxide cathode.

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Figure 7 Distribution of (a) Ni and (b) Co with the initial 5 laser pulses in LIBS across the cathodes cycled with LE, PHC and SA-PHC within different cycles at 55 °C between 3.0−4.3 V. The gradient from red to purple like rainbow reflects the gradual decrease in the element content.

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Electrochemical impedance spectra (EIS) are applied to shed light on the electrochemical function of the SA-PHC on the interfacial layer in high-Ni layered oxide cathodes. As presented in Figure 8d-e, the cell with LE exhibits a dramatically high interfacial charge-transfer resistance (Rct from the second semicircle in Figure 8a-c)57, resulted by the formation of the thick SEI film at the cathode surface. The cathode with PHC also shows sharp increase in Rct with cycling as a result of the poor interfacial compatibility of PHC on the cathode in the quasi-solid-state system. Uniquely, the low interfacial charge-transfer resistance is always accompanied with cycling for the cathode with SA-PHC, due to the good interfacial stability and compatibility from the in-situ metalophilic interaction between the cathode and electrolyte.

Figure 8 EIS spectra for the cell with (a) LE, (b) PHC and (c) SA-PHC after selected cycles at the state of charge (4.3 V) at 55 °C and the fitted data of Rsf (d) and Rct (e) acquired by the equivalent circuit (Inset of (a)). In the equivalent circuit, Rs stands for the combined internal resistance of cell components, Rsf represents the migrating resistance of Li ions through the

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surface layer, Rct refers to the interfacial charge-transfer resistance, Wo corresponds to Warburg impedance, CPE1 and CPE2 are the constant phase elements.

4. Conclusions In summary, an ion-conducting layer on the cathode/electrolyte interface for high-Ni layered oxide cathode is in-situ formed based on the functional GPE grafted by SA (SA-PHC). It is confirmed that the chemical interaction is acted on the TMs in cathode and special metalophilic reticulum structure in SA, which is responsible for manipulating the interfacial compatibility, ionic conductivity, and electrochemical activity of the cathode at the levels of both the particle and electrode in quasi-solid-state LIBs. As to the particle level, the phase transition and unfavorable microcracks in oxide microspheres and between primary crystallites can be effectively restrained. At the electrode level, such ion-conducting layer can not only enhance the interfacial compatibility, but also restrain the serious dissolution loss of transition metal cations, insuring the gentle gradient distribution of constituent elements through the depth of the porous electrode in quasi-solid-state system. As a result, the high-Ni oxide cathode with metalophilic SA-PHC presents satisfactory electrochemical performance at both the common and harsh conditions. This innovative work brings new insight into a smart in-situ interface manipulated strategy of promising high-Ni layered oxide cathodes in quasi-solid state battery. Supporting Information Ionic conductivity, electrolyte/active materials, FTIR spectra, and C1s XPS spectra of all the samples, the charge−discharge curves and cross-sections of the cathode with different electrolytes.

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Acknowledgments

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Graphical abstract

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Metalophilic Gel Polymer Electrolyte for in Situ Tailoring Cathode

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