Integration of Metal-Organic Frameworks into a 3D Interconnected Network for Improved Ion Transport

Metal-organic frameworks (MOFs) have attracted intensive study as solid electrolytes (SEs) in recent years, especially on facilitating ion transport with functionalized channels. However, MOF particles work separately in SEs and numerous interfaces hinder the high-eciency ion transport, which lowers the performance of solid-state batteries (SSBs) especially at high C-rate. Herein, we constructed continuous ion pathways by integration of MOFs into a 3D interconnected network. Particle arrays of a newly developed MOF (Zr-BPDC-2SO 3 H) which has single ion transport ability were grown on the bacterial cellulose (BC) nanobers to provide a linear ion transport network. The interconnected MOFs network exhibits higher ionic conductivity of 7.88 × 10 − 4 S cm − 1 at 25 ℃ , single ion transport ability ( (cid:0) Li+ =0.88), wide electrochemical window up to 5.15 V, excellent interface compatibility and capability for supressing lithium dendrites. Most importantly, the SSB fabricated with the interconnected MOFs network shows more than 100% improved specic capacity than the SSB without integration and stable cycling performance at 3 C. This work demonstrates the effectiveness of integrated design and paves new way for developing high-performance SEs based on porous ion conductors.


Introduction
Lithium metal battery appears to be highly promising for energy storage owing to its high energy density [1][2][3] . However, the unsafety of liquid lithium metal battery hinders its applications because of the ammable organic liquid electrolytes and lithium dendrites as well. Additionally, it also suffers from the narrow electrochemical window and the relatively low Li + transference number (usually lower than 0.4) 4 .
Solid electrolytes (SEs) are therefore applied to overcome the drawbacks of liquid electrolytes and suppress the growth of lithium dendrites 5 . SEs can be mainly classi ed into three categories: organic polymeric SEs, inorganic SEs and composite ones. Organic polymeric SEs, such as poly(ethylene oxide) (PEO) and poly(vinylidene uoride-co-hexa uoropropylene) (PVDF-HFP) 6,7 , are featured with excellent exibility and low interfacial resistance. However, their comparatively low ionic conductivities (10 − 6 -10 − 9 S cm − 1 ) at room temperature restrain their applications, whereas inorganic SEs such as garnet-type Li 5 La 3 Zr 2 O 12 (LLZO) and Li 2 S-P 2 S 5 possess comparable ionic conductivity to that of liquid electrolytes 8,9 . Unfortunately, the inorganic SEs often induce poor interfacial contact between the electrodes and the SE, and lithium dendrites tend to grow along the grain boundaries 10,11 . One strategy to overcome the disadvantages of polymeric and inorganic SEs is to use composite of polymeric and inorganic SEs 12 . However, the composite still does not satisfy the requirements of SEs in term of ionic conductivity, stability, safety etc. Another strategy, therefore, is to develop new kind of SE with high ion conductivity and ion transference number, wide electrochemical windows, good interfacial contact, and chemical/thermal stability.
Porous materials, such as metal-organic frameworks (MOFs), are expected to address the abovementioned requirements by virtue of their talents in porosity, physiochemical designability and structure engineering [13][14][15][16][17][18][19][20] . For example, the numerous pores in MOFs are capable of storing guest molecules such as ionic liquids containing Li + (Li-IL), which can assist the transport of Li + 21,22 . Chemical designability affords MOFs with intrinsic ionic conductivity because of the electronegative chemical groups functionalized inside the pores, and the ability for single ion transport. However, how to enhance the intrinsic ionic conductivity and improve the interface between the MOFs particles as well as MOFs and electrodes still remains a big challenge. Although a lot efforts have been devoted to raise the ion conductivity of MOFs intrinsically 23,24 . For example, CuBTC, HSPE-1-8, P@CMOF and UiOLiTFSI with ion conductive groups have been reported to have ionic conductivity up to 10 − 4 S cm -1 with certain liquid electrolytes [25][26][27][28] . It should be also noted that the Li + transport is a multidimensional process in MOFbased electrolytes, involving the transport along the conductive channel in MOF particles, between different MOF particles and on the interface of MOFs and electrodes 29 . In previous reports, the MOF particles function separately and ions transport across the interface of separated particles in sequence, which lowers the speci c capacity and cycling performance of the SSBs. Therefore, to construct a transport pathway for Li + between MOF particles and reduce the interfacial resistance between MOFs and electrodes, as well as to understand the transport mechanism are of vital importance to realize the use of MOFs as SE in high-performance lithium metal batteries.
Hereby, we propose a novel strategy to fabricate interconnected structure of MOFs to eliminate the interface between MOF particles, as well as build a linear pathway for ion transport as illustrated in Fig. 1.
To realize the e cient ion transport within MOF, a MOF functionalized with high-density electronegative groups (Zr-BPDC-2SO 3 H) was rstly designed, in which the functional groups can facilitate the ion transport along pore channels as shown in Fig. 1b. More importantly, the nanosized ion conductive MOF particles with interconnected structure were in-situ synthesized on the network of a bacterial cellulose (BC) skeleton (Fig. 1b). Furthermore, a exible interconnected network of conductive MOF can be obtained by virtue of the network of BC, which will provide integrated three-dimensional pathways for ion transport. Compared with the SE fabricated with conventional mixing strategy (MOFs/PVDF-HFP), the interconnected MOFs network exhibits excellent single ion transport ability ( Li + =0.88), much lower interfacial resistance (74 Ω), wider electrochemical windows (5.15V) and outstanding effect for suppressing the lithium dendrites. Most importantly, the interconnected MOF-based SE delivers a capacity as high as 119 mA h g − 1 at 3 C (more than 100% higher than that of the SSB without integration) and demonstrates a great improvement of cycling performance at high C-rate.
The large amounts of -SO 3 H groups are expected to serve as ion hopping sites because of its electronegativity and possible coulomb force toward metal ions, which can then enhance the ability of and 80% Zr-BPDC-2SO 3 Li + . The Zr-BPDC-2SO 3 Li + /PVDF-HFP membrane exhibits good exibility and mechanical strength ( Supplementary Fig. 5a). t Li + was calculated from the current-time curve and the ac impedance spectra before and after polarization using Li|Li symmetric cells referring to the Evans method at room temperature 32 . As shown in Supplementary Fig. 6 To apply MOFs as SEs in solid batteries, mixing MOFs with polymers is the most typical way to fabricate a exible SE membrane 23,33 . However, in this case, there is hardly chemical contact between MOF particles and thus the ion transport through the interface between MOF particles becomes challenging due to the lack of conductive medium. To solve this problem, in this work, the ion conductive MOF was insitu synthesized on the framework of BC nano bers and an interconnected MOFs network was obtained. BC was chosen here for its well-arranged linear chain and 3D structure. Both the peaks of Zr-BPDC-2SO 3 H and BC nano ber can be observed in the PXRD pattern of interconnected MOFs network ( Supplementary  Fig. 8), suggesting that the structure of MOF particles in the network remains the same as the separately synthesized Zr-BPDC-2SO 3 H crystals. As shown in Fig. 3a, BC serves as a template, and Zr-BPDC-2SO 3 H completely covers the nano bers of BC in the form of particle arrays without intervals, which is expected to be exempt from the long-distance ion transport and high interfacial resistance between the physically contacted individual particles in the mixture of MOF/polymer. Moreover, nodes can be observed between MOF arrays in the enlarged view, which results in an interconnected 3D network. The mass ratio of MOF and BC nano ber in the interconnected MOFs network was determined to be 7:3. The cross-sectional view in Fig. 3b shows that the interconnected MOFs network has a uniform thickness of about 86.9 µm. The ion transport ability of the interconnected MOFs network was also investigated. The interconnected MOFs network shows an Li + conductivity of 7.88 × 10 − 4 S cm − 1 with less than 17.63 wt% PC in pores ( Supplementary Fig. 9b, d), which is higher than that of the pellet of Zr-BPDC-2SO 3 H powders (2.65 × 10 − 4 S cm − 1 , 27.62 wt% PC in pores, Supplementary Fig. 9a, c). Moreover, the ion transference number of the interconnected MOF networks is as high as 0.88, indicating that the network inherits the excellent single ion transport ability of Zr-BPDC-2SO 3 H (Fig. 3d) and the transference number is higher than that of previously reported MOFs based SEs ( Supplementary Fig. 7a, Supplementary Table 1 Fig. 10). A smooth surface formed by uniform nanosized MOF chains can be observed in the interconnected MOFs network, which is supposed to bene t the interfacial contact with the electrodes. Whereas, numerous potholes can be observed in Zr-BPDC-2SO 3 Li + /PVDF-HFP, and aggregation of particles leads to uneven surfaces for Zr-BPDC- Fig. 5e, f), resulting in the higher interfacial resistance.

2SO 3 Li + /PVDF-HFP and Zr-BPDC-2SO 3 H/BC (Supplementary
Critical current density (CCD) is determined by the current density at voltage drop during the step increased galvanostatic test and higher CCD represents better capability for suppressing dendrites formation. To evaluate the CCD of different SEs, the current densities of galvanostatic test were step increased and the holding time for one cycle was 1 h using a symmetric Li|SE|Li cell. As shown in Fig. 4a, the voltage increases as the increasing of current density until the short circuit of the cells. Speci cally, the Li|MOF-based network|Li shows the lowest voltage at the same current density due to its least interfacial resistance (Fig. 3f). The CCD of the Li|Zr-BPDC-2SO 3  mAh cm − 2 , as proved in Fig. 4c. All this results certi cate that the interconnected MOFs network has remarkably improved capacity for suppressing the growth of lithium dendrites, which matches well with its better interfacial compatibility and electrochemical stability than other compared SEs ( Fig. 3e and f).
To interview how the different SEs in uence the deposition of Li + on lithium anode, the surface morphologies of lithium plates after 50 cycles of Li plating/stripping at current areal capacity of 0.10 mAh cm − 2 were investigated. Compared with the origin lithium plate ( Supplementary Fig. 11a1 and a2), the surface of lithium plate in the Li|interconnected MOFs network|Li symmetric cell is at and no lithium dendrites are observed ( Supplementary Fig. 11d1 and d2). However, lots of protuberance can be observed on the uneven surface of lithium plates in the symmetric cells with BC and Zr-BPDC-2SO 3 Li + /PVDF-HFP as SEs (Supplementary Fig. 11b1, b2, c1 and c2). The better capability of interconnected MOFs network for optimizing lithium deposition can be ascribed to the smooth and homogeneous surface of interconnected MOFs network. Its homogeneous surface created an even potential energy surface by -SO 3 − which can interact with Li + (proved by XPS results in Supplementary Fig. 4) and thus prohibited the inhomogeneous deposition of Li + at interface 34,35 , resulting in better cycling performance in Li-Li symetric cells.
Finally, Li metal SSBs were fabricated and the widely used commercial LiFePO 4 , super P and PVDF were mixed as the cathode to investigate the in uence of different SEs for battery performance. As shown in Fig. 5a and b, the average discharge capacities of SSB with interconnected MOFs network as SE at 0.2, 0.5, 1, 2, 3 and 5 C are 159, 145, 143, 140, 119 and 108 mA h g − 1 , respectively. In comparison, the Zr-BPDC-2SO 3 Li + /PVDF-HFP and Zr-BPDC-2SO 3 H/BC based SSBs have similar discharge capacities at low C-rates. However, their discharge capacities decline severely at high C-rates. For instance, the discharge capacity of the SSB with interconnected MOFs network as SE at 3 C is 100% and 200% higher than that of SSBs with Zr-BPDC-2SO 3 Li + /PVDF-HFP and Zr-BPDC-2SO 3 H/BC as SEs. The cycling performance of the SSB fabricated with interconnected MOFs network at 1 C is exhibited in Fig. 5c and d, which retains a stable discharge speci c capacity of around 140 mA h g − 1 after 500 cycles without decay. The cycling performances of SSBs with different SEs at 3 C are compared in Fig. 5e. The SSB fabricated with interconnected MOFs network remains a speci c capacity of 119 mA h g − 1 at 3 C after 600 cycles with a decay rate of 0.02% per cycle (Fig. 5f). However, the SSBs fabricated with BPDC-2SO 3 Li + /PVDF-HFP and Zr-BPDC-2SO 3 H/BC show low and fast decayed speci c capacity, indicating their disability for running at high C-rate. Moreover, the unchanged PXRD of interconnected MOFs network after cycling at 3 C proves its electrochemical stability during cycling ( Supplementary Fig. 8). Compared with the reported SSBs based on MOFs, the SSB fabricated with interconnected MOFs network shows excellent rate performance and cycling performance especially at high C-rate over 1 C (Supplementary Fig. 7b, Supplementary Table 2). The much better performance of interconnected MOFs network can be ascribed to its integrated linear channels for ion transport, optimized interfacial compatibility and ability for suppressing lithium dendrites certi cated by above-mentioned experiments.

Discussion
In summary, it has been proved that a 3D interconnected ion conductive network, which is constructed by newly developed particle arrays of single-ion conductive MOF ( were mixed in N-methyl pyrrolidone (NMP) at a mass ratio of 8:1:1 to obtain the cathode mixture. And then the cathode mixture was subsequently coated on aluminium foil. The prepared electrode lms were dried at 60 ℃ for 24 h under vacuum before cell fabrication. The aluminium foil was punched into circles with a diameter of 12 mm. The loading of the active material is around 1.5-2 mg cm -2 .

Declarations
Competing interests The authors declare no competing interests.