Stabilization of Hexaaminobenzene in a 2D Conductive Metal

Article Cite This: J. Am. Chem. Soc. 2018, 140, 10315−10323

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Stabilization of Hexaaminobenzene in a 2D Conductive Metal− Organic Framework for High Power Sodium Storage Jihye Park,†,⊥ Minah Lee,†,⊥ Dawei Feng,† Zhehao Huang,§ Allison C. Hinckley,† Andrey Yakovenko,∥ Xiaodong Zou,§ Yi Cui,‡ and Zhenan Bao*,†

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Department of Chemical Engineering, and ‡Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, United States § Berzelii Centre EXSELENT on Porous Materials, Department of Materials and Environmental Chemistry, Stockholm University, Stockholm SE-106 91, Sweden ∥ X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois 60439, United States S Supporting Information *

ABSTRACT: Redox-active organic materials have gained growing attention as electrodes of rechargeable batteries. However, their key limitations are the low electronic conductivity and limited chemical and structural stability under redox conditions. Herein, we report a new cobalt-based 2D conductive metal−organic framework (MOF), Co-HAB, having stable, accessible, dense active sites for high-power energy storage device through conjugative coordination between a redox-active linker, hexaaminobenzene (HAB), and a Co(II) center. Given the exceptional capability of Co-HAB for stabilizing reactive HAB, a reversible three-electron redox reaction per HAB was successfully demonstrated for the first time, thereby presenting a promising new electrode material for sodium-ion storage. Specifically, through synthetic tunability of Co-HAB, the bulk electrical conductivity of 1.57 S cm−1 was achieved, enabling an extremely high rate capability, delivering 214 mAh g−1 within 7 min or 152 mAh g−1 in 45 s. Meanwhile, an almost linear increase of the areal capacity upon increasing active mass loading up to 9.6 mg cm−2 was obtained, demonstrating 2.6 mAh cm−2 with a trace amount of conducting agent.

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INTRODUCTION Organic compounds have drawn increasing interest for energy storage devices, including batteries and capacitors, because of their redox activity, ubiquity, and lighter weight than conventional transition metal electrodes.1,2 The synthetic tunability of organic compounds also enables versatile electrochemistry to store various ions (e.g., Li+, Na+, K+, Mg2+), endowing these materials with the flexibility to satisfy the design requirements for next-generation energy devices.3−5 However, the use of organic electrodes often faces several challenges, such as dissolution in electrolytes, degradation of electrochemical activity, and slow reaction kinetics, primarily due to the weak stability and low electrical conductivity of the electrode materials.6,7 To alleviate the solubility issue, polymerization and immobilization of active molecules have been extensively studied,8 but these methods are often synthetically challenging and suffer from limited accessibility and low density of active sites.9,10 To solve the low conductivity problem, composites with conductive nanostructures have been studied and shown to successfully increase the rate capability of the system.11,12 However, the addition of an inactive conductive agent decreases the relative loading of the active component per unit mass and volume, thus inevitably © 2018 American Chemical Society

limiting the overall energy density of the electrode. To overcome these limitations, a simple, stable system with dense active sites and relatively low impedance is highly desired to constitute energy-dense, high-power electrodes for advanced energy storage technologies. As a unique category of coordination polymers, metal− organic frameworks (MOFs) present one of the simplest yet most effective platforms to immobilize organic molecules through metal−ligand coordination (M−L) bonds.13 The precise arrangement of their building blocks enables crystalline and porous frameworks, providing access to built-in active sites for electrochemical reactions. So far, several MOFs have been reported as potential electrode materials for energy storage, including lithium-ion batteries14−24 and sodium-ion batteries.25 Although these MOFs show moderate activities and relatively stable cycling, most studies showed limited redox capabilities due to the sparse active sites and low electrical conductivity.21 More recently, conductive MOFs have drawn gowing attention for electrochemical applications,19,20,26 and a notable breakthrough was reported utilizing conductive MOFs Received: June 8, 2018 Published: July 25, 2018 10315

DOI: 10.1021/jacs.8b06020 J. Am. Chem. Soc. 2018, 140, 10315−10323

Article

Journal of the American Chemical Society

Figure 1. (a) Synthetic scheme of Co-HAB. (b) Proposed three-electron reversible reaction in Co-HAB. (c) The calculated structure of Co-HAB in 3D. Na+ and PF6− ions are illustrated with size information for comparison. Hydrogens are omitted for clarity. (d) HRTEM image of Co-HAB along [001] showing a hexagonal pore packing. Inset: Fourier transform of the image. Scale bar = 10 nm (left). HRTEM image in the red box. Scale bar = 2 nm (top right). Lattice-averaged and symmetry-imposed image obtained from the HRTEM image. The eclipsed structure model of CoHAB is overlaid (bottom right). (e) Synchrotron PXRD patterns of Co-HAB and Pawley profile fit (λ = 0.45212 Å). (f) N2 sorption isotherm of Co-HAB.

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RESULTS AND DISCUSSION The imine functional group can undergo redox processes at the CN bond, and thus imine compounds can provide a high theoretical specific capacity due to their small mass unit.6 However, the imine species often undergo irreversible selfcondensation/polymerization or other side reactions during the redox processes, precluding reversible charge storage in energy devices.6,28,29 One possible solution to stabilize these reactive intermediates is to form a coordination complex with transition metals wherein the negative charge formed upon reduction can be delocalized through d-π conjugation.30 Particularly, when such d-π conjugation is extended to 2D structures, the system often yields an electrically conductive network, which can be classified as a conductive MOF.31−33 Combining electrical conductivity and redox activity in a MOF enables fast electron transport to the redox centers, and thus offers an ideal system that can fully utilize imine redox chemistry for rechargeable batteries with high power density. Bearing these features in mind, we selected hexaaminobenzene (HAB) as the organic linker to construct the 2D conductive MOF (Figure 1a). Because the fully oxidized form of HAB can be considered a dense assembly of six imine groups, it can, in theory, undergo up to six-electron redox reactions. Thus, HAB could have one of the highest densities of redox centers among all possible molecular structures, which satisfies the primary design principle for energy-dense electrode materials. The bridging metal also plays an important role in conductive MOFs. It governs the electronic structure,

for energy storage based on electrochemical double layer mechanism; however, these materials exhibited negligible redox activity,27 which limits their energy density. Integration of both high density of redox-active centers and high conductivity into a MOF-based platform is a promising albeit challenging route to fast and energy-dense electrodes. Herein, we report a 2D conductive MOF consisting of a redox-active hexaaminobenzene (HAB) and Co(II) ion node, Co-HAB. With systematic control of the synthetic conditions to increase particle size and crystallinity, our obtained Co-HAB exhibited bulk electrical conductivity as high as 1.57 S cm−1. Meanwhile, the Co-HAB displayed outstanding chemical stability in both aqueous and organic media, and thermal stability to 300 °C. We confirmed that the conductive Co-HAB stores nearly three Na+ and electrons per HAB in an organic electrolyte (theoretical specific capacity of 312 mAh g−1) while showing substantial pseudocapacitive contributions. With high active loading up to 90 wt % in the electrode, Co-HAB delivers 228 mAh g−1 at 1 A g−1 and 151 mAh g−1 at 12 A g−1, outperforming many other organic and inorganic electrode materials for high-power Na+ storage devices. To the best of our knowledge, Co-HAB is the first conductive MOF to demonstrate storage of sodium and, moreover, to do so rapidly and stably. Co-HAB is thus an excellent candidate for a highpower electrode material to build fast Na-ion batteries or hybrid capacitors. 10316


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Figure 2. (a) HRTEM images of Co-HAB-A10 (left) and Co-HAB-D (right). (b) SEM images of Co-HAB-A10 (left) and Co-HAB-D (right). (c) PXRD patterns of Co-HABs synthesized from different conditions. (d) Full width at half-maximum (FWHM) obtained from PXRD diffractions, which corresponds to [020] using a Gaussian peak-shape function. (e,f) Electrical conductivity measurement of Co-HAB in bulk as a function of particle size (size parmeter was set by multiplication of the short side and long side of the particles of each sample from SEM images) and fwhm, respectively. (g) N2 sorption isotherms and (h,i) PXRD patterns of Co-HAB after treatments in different chemical environments.

(A stands for aqueous) were obtained (Figure 1a). To validate the proposed 2D structure, density functional theory (DFT) calculations were carried out. First, a structural model with an eclipsed packing (Figure 1c) was generated, and the geometry was optimized by DFT calculations. The unit cell parameters of the model were further refined by applying Pawley fit against the synchrotron powder X-ray diffraction (PXRD, λ = 0.45212 Å, Figure 1e). The resulting hexagonal unit cell parameters (a = b = 13.361(3) Å and c = 3.082(1) Å, corresponding to the interlayer spacing) are in good agreement with the initial DFTcalculated structure (Table S1), as the d-spacings of d001 = 3.08 Å and d010 = 11.57 Å from PXRD match well with those of the initial eclipsed model (d001 = 3.19 Å and d010 = 11.36 Å). To further confirm the structure, high-resolution TEM (HRTEM) images of the Co-HAB-A and an optimized product, Co-HABD (discussed later), were obtained. The HRTEM images (Figure 1d) clearly show hexagonal pores, smaller than 1 nm with a honeycomb arrangement along [001] with d010 = 11.7 Å, which corresponds well with the eclipsed model (calculated d010 = 11.6 Å). Meanwhile, 1D channels along the rod-like crystals can also be observed, which match well with our eclipsed structural model. Furthermore, the N2 adsorption isotherm of Co-HAB-D shows a BET surface area of ∼240 m2 g−1 with subnanometer pores (Figure 1f), which cannot be achieved with the nearly nonporous staggered structure.41 This result, in corroboration with the well-defined pores seen by HRTEM, validates the 2D eclipsed honeycomb structure. During the synthetic optimization, we found that the stoichiometry of the base with respect to HAB plays a critical role in determining the Co-HAB growth. For instance, when less than 3 equiv (with respect to HAB) of NH4OH was used,

conductivity, and chemical and thermal stability of the MOF on the basis of orbital interaction with the linker.34,35 Previously reported conductive MOFs typically used Cu(II) and Ni(II) as the metal nodes due to these ions’ preferred D4h coordination symmetry when interacting with strong field ligands (e.g., −NH2, −SH).26,33,36−38 The D4h coordination often yields a (2,3)-connected honeycomb 2D lattice when connected to a hexadentate ligand with D3h symmetry. Such structure can lead to a highly conjugated system because of the effective π orbital overlap between the metal center and the ligand. Co(II) has been relatively rarely studied in bulk synthesis of conductive MOFs, presumably due to the synthetic challenges associated with the less preferred D4h coordination symmetry of Co(II) species. Nevertheless, we hypothesized that cobalt conductive MOF may be an interesting system to explore the design of anode materials because of its weaker electron affinity than Cu(II) and Ni(II).39 In the reported synthesis of conductive MOFs with amino group-containing linkers, the proposed mechanism includes the deprotonation and subsequent partial oxidation of the linkers.40 This is typically achieved through an excess amount of base (e.g., NH4OH), which is mixed with the metal salts prior to the addition of the linker.31,33 However, perhaps due to the stronger affinity of Co(II) to OH− as compared to Cu(II) and Ni(II), we found this synthetic route produced a considerable amount of Co(OH)2 as an impurity (Figure S1). To avoid the hydroxide formation, HAB and Co(II) salts were first mixed in water to allow complexation between them. NH4OH was subsequently added in open air, and the reaction was stirred until dark navy-colored precipitates of Co-HAB-A 10317


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Figure 3. (a) Voltage profiles of Co-HAB-D in a Na cell at a current of 50 mA g−1 in the voltage window of 0.5−3.0 V. (b) Cycle retention of CoHAB-D in the voltage window of 0.5−3.0 V. (c) Rate capability of Co-HAB-D in comparison with that of Co-HAB-A4 and Co-HAB-A10. (d) Discharge profiles of Co-HAB-D at various current densities. (e,f) Capacity retention of Co-HAB-D at high current densities.

(Figure 2c−f). Specifically, a value of 0.48 S cm−1 was measured for the Co-HAB-A10 sample, which was more than 6 times higher than that of the Co-HAB-A4 sample (Figure S6). Encouraged by this crystallinity−conductivity relationship, we further optimized the synthesis to obtain a more crystalline product for higher electrical conductivity. For the modified synthesis, N,N-dimethylformamide (DMF) was chosen as a cosolvent because of its poorer ability to solvate the reactants relative to water. It was thus expected to slow the MOF growth, allowing for a more crystalline product than the fully aqueous Co-HAB-A series. Indeed, the resulting product (CoHAB-D) synthesized from DMF/water mixture (1:1, v/v) showed much-improved crystallinity and a bulk conductivity of 1.57 S cm−1 (Figure 2e,f). It is worth noting that this is one of the highest values, comparable with other conductive HABand benzenehexathiolate (BHT)-based MOFs.32,38,40,41,46,47 Crystallites of Co-HAB-D exhibit an even larger size (∼75 nm, on the long side) than Co-HAB-A10 (Figure 2a,b). The conductivity of these materials is many orders of magnitude higher than that of the nanosheets of the previously reported 2D coordination polymers prepared from similar starting materials via a very different synthetic route.32 The correlation between crystallinity and conductivity provides a rationale for these findings, given the significantly higher crystallinity of our MOF materials relative to these coordination polymers. The high conductivity of Co-HAB represents one of the highest values among the reported conductive MOFs measured in pellets.26 Thus, we expect that Co-HAB would provide fast electron transport within the electrode, which is highly desirable to achieve high power energy devices. While conductive MOFs have drawn growing attention in various energy applications, studies of their chemical stability have been somewhat limited. However, the chemical stability of electrode materials is a critical parameter to evaluate their practical potential as it dictates their handling process, testing conditions, lifetime, and recyclability.48 Among the reported

an amorphous product, Co-HAB-A3, was obtained with a low yield (Figure 2c). The filtrate of this reaction was strongly colored dark blue, which is indicative of unconverted HAB linkers (Figure S2). The yield and crystallinity of the products improved progressively with increasing base amount up to 10 equiv, Co-HAB-A10 (Figure 2c). Scanning electron microscopy (SEM) shows that the better crystalline products tend to exhibit larger rod-like crystallites (Figure S3). Another structural confirmation was obtained using X-ray photoelectron spectroscopy (XPS). The high-resolution Co 2p spectra display satellite peaks in both the Co 2P3/2 and the Co 2P1/2 regions, typical features of the cobalt ion in the +2 oxidation state in square planar geometry.42,43 We also noticed a small shoulder at 781.1 and 793.0 eV, which we attribute to structural disorder, a deviation from the ideal coordination of the neutral complex in the MOF. The disordered feature in the spectrum decreases as the amount of the base increases, which coincides with the improving crystallinity of these materials. The high-resolution N 1s spectra can be deconvoluted into two peaks; the binding energies at 398.3 and 400.0 eV are attributed to the quinoid imine (N−) and benzoid amine (−NH−),44 consistent with the expected complex structure45 in Co-HAB (Figure 5e,f). Having confirmed the physical structure, we further examined the electronic structure and electrical conductivity of Co-HAB using UV−vis−NIR spectroscopy and the fourpoint probe conductivity measurement, respectively. The CoHAB film on a quartz substrate shows a broad absorption band in the near-IR region, which corresponds to an absorption edge of 0.87 eV, suggesting a high degree of conjugation arising from the sufficient orbital overlap between HAB and Co(II) (Figure S5). The electrical conductivity was measured on a pressed pellet of Co-HAB. We hypothesized that the crystallinity and particle size would correlate with the bulk conductivity as we expect the conductivity to be maximized within a crystal domain. As expected, the electrical conductivity increases with the improved crystallinity and larger particle size 10318


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Figure 4. (a) Voltage profiles comparison of Co-HAB-D electrodes with various active mass loadings. (b) Areal capacity of Co-HAB-D electrodes as a function of active mass loading. The red line indicates the theoretical values. (c) Rate capability of Co-HAB-D with active mass loading of 4.4 mg cm−2. (d) Areal capacity comparison of Co-HAB-D with representative anode materials for fast Na-ion storage in a previous study.38,50−59

g−1 was obtained, corresponding to 2.8 Na+ per HAB (Figure 3a). Considering partially oxidized HAB in the neutral CoHAB platform can theoretically store 3 Na+ (312 mAh g−1), this result demonstrates nearly full utilization of redox activity of HAB was realized in the MOF for the first time. After the formation of the solid/electrolyte interface (SEI) layer in the first discharge, negligible changes in the voltage profiles were observed during repeated cycles, which suggests that the electrochemical reaction is highly reproducible. Figure 3b shows the cycle stability of the Co-HAB, wherein a capacity of 226 mAh g−1 at a current rate of 500 mA g−1 was maintained with a Coulombic efficiency close to 100% for more than 50 cycles (Figure 3b). We noticed that further reduction of CoHAB makes the MOF substantially less stable, as indicated by poor cycle retention. According to the voltage profiles obtained with a lower discharge cutoff potential at 0.05 V versus Na+/ Na, only 34% of the specific capacity was retained after 40 cycles in a Na cell (Figure S9). These results imply that the ultralow potential during the discharge may further reduce Co(II), causing irreversible redox processes, likely involving the destruction of the framework (Scheme S1c). To maintain the structural stability and the cycle performance, the voltage window was thus fixed to be above 0.5 V versus Na+/Na. To ascertain the impact of electron transport in Co-HAB on electrochemical energy storage, we compared the rate performances of Co-HAB-D with Co-HAB-A4 and CoHAB-A10 in Na cells (Figure 3c). Notably, the Co-HAB-D electrode yielded much higher specific capacities at the elevated current densities of 0.5 and 1 A g−1, which can be

MOFs, Co-HAB exhibits superior chemical stability both in aqueous and in organic media. PXRD and N2 sorption isotherms indicate Co-HAB maintains its framework integrity after treatments with concentrated NH4OH, boiling water, and 0.1 M acidic and basic solutions of HCl and KOH, respectively (Figure 2g,h). The chemical stability tests of Co-HAB were also carried out in the organic electrolytes [1 M LiPF6 in ethylene carbonate/diethyl carbonate (EC/DEC) and 1 M NaPF6 in diethylene glycol dimethyl ether (DEGDME)] that are typically used for electrochemical studies and showed wellretained crystallinity of Co-HAB (Figure 2i). In addition to its remarkable chemical stability, Co-HAB also exhibits an outstanding thermal stability; its PXRD patterns and N2 adsorption isotherms remain nearly unaltered after heat treatments at different temperatures (100−300 °C) (Figure S8). Such exceptional stability can be attributed to the robust coordination bonds between HAB and Co(II) as well as the strong chelating effect resulting from the hexadenticity on the smallest benzene unit.49 These features allow Co-HAB to be coupled with a wide range of electrolytes and to run at elevated temperatures, enabling rigorous optimizations of the operating conditions in energy devices. With its optimized conductivity and exceptional chemical/ thermal stability, Co-HAB-D was tested as an electrode material to store sodium ions. To evaluate the redox capability and cycle stability of Co-HAB, we tested electrodes composed of 90 wt % Co-HAB as an active material in half-cells versus sodium metal. In a voltage range of 0.5−3.0 V versus Na+/Na at 50 mA g−1, a reversible specific capacity as high as 291 mA 10319


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Figure 5. (a) Cyclic voltammograms of Co-HAB at different sweep rates. (b) b-Value determination of the cathodic currents assuming that the current obeys a power-law relationship. The inset shows peak currents versus scan rate in logarithm and the fitted lines at 1.25 V (red) and 0.5 V (blue) during reduction. (c) Capacitive and diffusion currents contributed to the charge storage of Co-HAB at the rate of 0.375 mV s−1. (d) A Selfdischarge profile of Co-HAB in a half-cell versus Na metal. The inset shows voltage profiles before and after storage for the self-discharge test. (e,f) Ex situ XPS spectra of N 1s and Co 2p regions during discharge and recharge processes in a potential window of 0.5−3.0 V.

of carbon black. For the 4.4 mg cm−2 electrode, the areal capacity of 0.85 mAh cm−2 can be achieved at the current density of 4.4 mA cm−2, further supporting the exceptional high-rate performance of Co-HAB-D (Figure 4c). We compared the areal performance of Co-HAB-D with other sodium electrodes, primarily anodes showing high rate performances.38,50−58 Among the reported sodium anodes delivering the capacity within 20 min to date, Co-HAB-D demonstrated the highest areal capacity (Figure 4d). In previous studies, the areal capacities were barely assessed at high rates because their high rate capabilities were attained predominantly by using nanostructured electrodes with small mass loadings and limited packing densities (Table S3). In contrast, our MOF platform, exhibiting high electrical conductivity and ordered subnano 1D channels, affords high mass loading electrodes to have proper electrical and structural integrity. Considering its areal capacity, cycle stability, rate capability, and high active loading, Co-HAB represents a superb electrode material for ultrafast Na-ion batteries and Naion hybrid capacitors. To understand the electrochemical processes for Na-ion storage, cyclic voltammetry (CV) was carried out at sweep rates ranging from 0.375 to 3 mV s−1. The cyclic voltammogram displays two quasi-reversible reduction/oxidation processes (Figure 5a). As Co-HAB can store nearly 3 Na+ per HAB, which corresponds to 2 Na+ per coordination unit during the galvanostatic cycling, these features represent successive reduction by two electrons, yielding a dianionic coordination unit in Co-HAB (Scheme S1b). The resulting negative charges on Co-HAB can be balanced with incoming Na+ ions, thereby sustaining a reversible reaction mechanism in the cation host for energy storage.

attributed to (i) ideal crystallographic pathways for Na+ diffusion and (ii) a higher electrical conductivity as compared to that of Co-HAB-A electrodes. Figure 3d shows the discharge profiles of Co-HAB-D at various current densities from 0.1 to 12 A g−1. Notably, Co-HAB-D still demonstrates a specific capacity of 214 and 152 mAh g−1 even at the extremely high current densities of 2 and 12 A g−1, respectively. As demonstrated by its stable capacity retention at high rates (Figure 3e,f), Co-HAB-D has outstanding tolerance to repeated fast sodiation/desodiation processes. It should be emphasized that such high rate capability obtained with 90 wt % active loading was primarily due to the improved electrical conductivity of Co-HAB-D. We further confirmed that carbonfree electrode can still exhibit 76% of the theoretical capacity (Figure S10). Considering most organic compounds require a large amount of conductive additive to compensate for the lack of intrinsic electrical conductivity,11,12 our results highlight the advantage of the intrinsic conductivity in the redox-active MOF platform. We further tested the electrochemical properties of CoHAB-D electrodes having higher mass loadings to increase the areal capacity, which is one of critical metrics for practical applications. As shown in Figure 4a and b, the areal capacity can be increased almost linearly from 0.35 to 2.6 mAh cm−2 with the increase of active mass loading from 1.3 to 9.6 mg cm−2. The voltage profile of the electrode with the highest active mass loading of 9.6 mg cm−2 is almost identical to that of 1.3 mg cm−2 at the same rate of 50 mA g−1. More importantly, 90% of specific capacity of the 1.3 mg cm−2 electrode was still exhibited in the 9.6 mg cm−2 electrode (Figure S11). The areal capacity of 2.6 mAh cm−2 with negligible polarization confirms remarkable charge transport kinetics in the Co-HAB-D electrode, which only contains 5% 10320


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and electrochemical conditions, which is beyond the scope of this Article. Nonetheless, the foregoing results demonstrate our concept of employing coordination chemistry between HAB and metal centers to construct conductive MOFs can be an effective and general approach to better exploit electrochemical activity of organic molecules for Na+ storage.

We further performed quantitative analyses on the chargestorage behaviors in Co-HAB, which showed substantial pseudocapacitive contributions. By plotting the peak currents (i) and sweep rates (v) in logarithm, we can determine the bvalue of the cathodic currents (Figure 5b) with an assumption that the current obeys a power-law relationship, i = avb (a and b are adjustable values).60 Both b-values for cathodic peaks at 1.25 and 0.5 V are 0.74, which confirms pseudocapacitive charge storage in Co-HAB. By separating the diffusioncontrolled process where i varies as v1/2, and the capacitive process where i varies as v,61 the capacitive contribution was calculated to be 47% out of the total current, at a sweep rate of 0.375 mV s−1 (Figure 5c). Given the pseudocapacitive storage mechanism, the self-discharge behavior of Co-HAB was examined (Figure 5d). The time-dependent voltage change on an open circuit was recorded with a fully charged cell for 10 days. The result shows that Co-HAB exhibits remarkably low self-discharge rate of less than 0.003 V h−1. This is significantly lower than most other capacitors that typically lose one-half of the maximum voltage within a few hours.62 Furthermore, after 10 days of the storage, more than 85% of original capacity was still confirmed, thus proving its outstanding stability in the fully charged state. In support of the redox mechanism, detailed studies were implemented by ex situ XPS, which provided evidence for the evolution of the oxidation state of Co-HAB in response to changes in the state of charge of the electrode. Both the cobalt node and the HAB of the cycled MOF electrode were monitored to understand the origin of the redox activity. In the N 1s spectra (Figure 5e), reversible changes of the peak intensities at 398.3 and 400.0 eV were observed during the discharge and recharge processes. This represents the reversible reduction of the quinoid imine moiety in Co-HAB to the corresponding benzoid amine species during the discharge/charge cycle. In contrast, no obvious change was observed in Co 2p spectra (Figure 5f), indicating cobalt remains divalent in the cycled electrode during the charge/ discharge. This result suggests that the reduction mainly occurs on the HAB instead of the Co(II). This is in accordance with observations of MII(o-phenyldiamine)2 complexes (M = Ni, Pd, Pt),30 which concluded that the reduction from neutral to dianionic species is invariably centered on the ligand. On the basis of these results and the previous findings, we propose that the redox activity of Co-HAB is centered on the HAB linkers, while the cobalt serves as the bridge to immobilize HAB and imparts high electrical conductivity to the 2D framework. To demonstrate the broader feasibility of using the conductive MOF platform with tunable metals for sodium storage, the electrochemical property of Ni-HAB was also examined (Figure S12). Under conditions identical to those used for Co-HAB, Ni-HAB exhibited a voltage profile similar to that of Co-HAB with a reversible capacity of 295 mAh g−1, corresponding to 2.8 Na+ per HAB. The cycle stability, however, was marginally worse than that of Co-HAB, which would require further investigation. During the course of our study, we noticed that reversible Li+ storage in Ni-HAB has been recently reported.16 A reversible capacity of approximately 100 mAh g−1 was demonstrated in the voltage range of 2−3.5 V vs Li+/Li, which corresponds to one Li+ storage per HAB. In their study, the redox reaction was observed to be centered on Ni atoms by XPS as opposed to our observation with Co-HAB. This contrasting observation to ours in the redox chemistry can be potentially due to dissimilar materials

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CONCLUSION We successfully address the stability and conductivity challenges associated with redox-active organic compounds by designing a 2D conductive MOF, Co-HAB. Systematic control of the synthetic conditions allows the highly crystalline Co-HAB to combine multiple features of an ideal electrode, including high intrinsic conductivity, high density of redoxactive sites, porosity, and excellent chemical/thermal stability. We first time demonstrated that Co-HAB can store three electrons and Na+ per HAB in organic electrolytes, thereby exhibiting a specific capacity as high as 291 mAh g−1 with a stable cycling life. In addition, Co-HAB shows a remarkable rate capability, yielding 152 mAh g−1 within 45 s, conferred by its high intrinsic conductivity and porosity. Given the chemical tunability of both the organic and the inorganic building blocks in this MOF system, we envision that the configuration of other redox-active organic molecules in a conductive framework should afford additional high-power electrode materials.

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EXPERIMENTAL SECTION

Synthesis of Co-HAB-A Series by Varying the Amount of NH4OH. A solution of 21.0 mg (0.072 mmol) of cobalt nitrate hexahydrate [Co(NO3)2·6H2O] was dissolved in 6 mL of water. Ten milligrams of HAB·3HCl (0.036 mmol) in 5 mL of H2O was then added into the cobalt nitrate solution in air with stirring (300 rpm). 60 μL of 6 M aqueous ammonium hydroxide (NH4OH) (0.36 mmol, 10 equiv to HAB) was then subsequently added to yield Co-HABA10. The mixture was stirred for an hour at rt in air. The resulting dark navy solids were filtered and washed with copious amounts of water and acetone, and dried under vacuum at 80 °C for 3 h unless otherwise noted. Likewise, the reactions with different stoichiometry of NH4OH (3−8 equiv) were denoted with a number that corresponds with base equivalent used in the synthesis (e.g., CoHAB-A3−Co-HAB-A8). Synthesis of Co-HAB-D Using Mixed Solvent. A solution of 105 mg (0.36 mmol) of cobalt nitrate hexahydrate [Co(NO3)2· 6H2O] was dissolved in a mixture of 25 mL of DMF and 20 mL of water in a round-bottom flask. The solution was preheated at 75 °C in the oil bath for 15 min. 50 mg of HAB·3HCl (0.18 mmol) in 5 mL of H2O was then added to the cobalt nitrate solution under stirring in open air. 180 μL of 6 M aqueous ammonium hydroxide (NH4OH) (1.08 mmol, 6 equiv to HAB) was added immediately. The reaction mixture was stirred for 2 h at 75 °C. The resulting dark navy solids were filtered and washed with copious amounts of water and acetone, and dried under vacuum at 80 °C for 3 h. MOF Electrode Preparation. To prepare the working electrodes, Co-HAB powders were mixed with carbon black conductive additive (Super P, TIMCAL) and polytetrafluoroethylene (PVDF) binder in the mass ratio of 90:5:5. Subsequently, N-methyl-2-pyrrolidone (NMP) was added to the mixture and stirred for 1 h. The resulting slurry was coated on Cu foil and dried at 70 °C overnight under vacuum. The dried samples were cut into 1 cm2 circular disks to afford a mass of 1−1.5 mg for typical electrodes. To prepare high mass loading electrodes, Co-HAB powders were mixed with polytetrafluoroethylene (PTFE) binder instead of PVDF in the same mass ratio of 90:5:5. Glass microfiber filters (Whatman, GF/D) were used as the separator. 1 M NaPF6 in diethylene glycol dimethyl ether served as the electrolytes. Na metal was used as the anode for the Na cells. The resulting coin cells were loaded into a battery tester (Arbin 10321


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Journal of the American Chemical Society Instruments) for galvanostatic cycles and into a VSP300 potentiostat (Biologic) for cyclic voltammetry.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/jacs.8b06020. Characterization details and additional experimental data (PDF)

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AUTHOR INFORMATION

Corresponding Author

*[email protected] ORCID

Jihye Park: 0000-0002-8644-2103 Xiaodong Zou: 0000-0001-6748-6656 Yi Cui: 0000-0002-6103-6352 Zhenan Bao: 0000-0002-0972-1715 Author Contributions ⊥

J.P. and M.L. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge the support from the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy through the Advanced Battery Materials Research (BMR) Program. J.P. acknowledges support from the Dreyfus Foundation Postdoctoral Fellowship for Environmental Chemistry. M.L. acknowledges partial support by the Postdoctoral Fellowship from the National Research Foundation of Korea under grant no. NRF-2017R1A6A3A03007053. D.F. acknowledges support from the U.S. Department of Energy, Office of Sciences, Office of Basic Energy Sciences, to the SUNCAT Center for Interface Science and Catalysis. A.C.H. acknowledges support from the National Science Foundation Graduate Research Fellowship under grant no. DGE-1147474. The structural characterization by HRTEM and PXRD was supported by the Knut & Alice Wallenberg Foundation through the project grant 3DEMNATUR and the Swedish Research Council (VR) through the MATsynCELL project of the Röntgen-Ångström Cluster.

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