Coordination Polymers Derived from 1,2,4,5 ... - ACS Publications

Article Cite This: Chem. Mater. 2019, 31, 5197−5205

pubs.acs.org/cm

Nickel(II) and Copper(II) Coordination Polymers Derived from 1,2,4,5-Tetraaminobenzene for Lithium-Ion Batteries Roman R. Kapaev,*,†,‡,§ Selina Olthof,∥ Ivan S. Zhidkov,⊥ Ernst Z. Kurmaev,⊥,# Keith J. Stevenson,† Klaus Meerholz,∥ and Pavel A. Troshin†,‡ †

Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Nobel str. 3, Moscow 143026, Russia Institute for Problems of Chemical Physics RAS, Acad. Semenov str. 1, Chernogolovka 142432, Russia § D.I. Mendeleev University of Chemical Technology of Russia, Miusskaya sq. 9, Moscow 125047, Russia ∥ Department of Chemistry, University of Cologne, Luxemburger str. 116, Köln 50939, Germany ⊥ Institute of Physics and Technology, Ural Federal University, Mira str. 19, Yekaterinburg 620002, Russia # M.N. Mikheev Institute of Metal Physics of Ural Branch of Russian Academy of Sciences, S. Kovalevskoi str. 18, Yekaterinburg 620108, Russia

Downloaded via BUFFALO STATE on July 24, 2019 at 03:37:02 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

‡

S Supporting Information *

ABSTRACT: Highly conductive electrochemically active materials are required for developing a new generation of ultrafast lithium-ion batteries (LIBs). Recently, a novel family of transition metal coordination polymers derived from arylamines exhibited conductivities of over 1 S cm−1. Low molecular weight analogues of these materials show rich and reversible electrochemical behavior. However, there are just very few reports on the application of such materials in LIBs. In this paper, linear nickel(II) and copper(II) coordination polymers derived from 1,2,4,5-tetraaminobenzene are reported and investigated as anode and cathode materials for LIBs. In the anode mode, both materials show ultrafast cycling behavior with impressive stability. Particularly, for the nickel-based material, a specific capacity of 83 mA h g−1 is reached at 20 A g−1 current density, and 79% of this capacity is retained after 20 000 cycles. Moreover, the copper-based polymer used as a cathode component shows a specific capacity of up to 262 mA h g−1 in the voltage range of 1.5−4.1 V vs Li/ Li+, which corresponds to the energy density of 616 W h kg−1.

■

a higher current rate of 500 mA g−1, the capacity reduced to ∼55 mA h g−1. HAB-derived MOFs were also successfully used in sodiumion batteries11 and supercapacitors,12 and were proposed for confining lithium polysulfides in Li−S batteries.13 The rapid development of materials science applications of HAB-based MOFs is hindered by poor availability and extreme air sensitivity of the starting benzenehexamine precursor, which readily decomposes or gets oxidized during the synthesis or while handling.14 Surprisingly, coordination compounds derived from other arylamines have not been investigated so far as active materials for metal-ion batteries. Unlike HAB, 1,2,4,5-tetraaminobenzene (TAB) is commercially available as a tetrachloride salt. This compound is more stable due to the fewer number of electron-donating NH2groups. TAB-derived one-dimensional (1D) coordination polymers have been successfully synthesized and used for memory devices and micro-supercapacitors.15,16 Here, we report that TAB-based nickel and copper linear coordination polymers can be used as ultrafast anode materials stable over

INTRODUCTION Lithium-ion batteries (LIBs) are very attractive because of their relatively high energy density and good cycling stability.1,2 However, traditional LIBs typically have a limited rate capability, i.e., they are not able to charge and discharge quickly (e.g., with the current rates above 5C) without severely sacrificing their capacity.3,4 One of the key requirements for LIBs operating at high current densities is high electronic and ionic conductivity of the redox-active electrode materials.3 Recently, highly conductive (σ > 1 S cm−1) metal−organic frameworks (MOFs) derived from arylamines and Ni(II) or Cu(II) salts have been reported.5−7 These planar structures, which possess transition metal cations bonded by four NHgroups, were shown to have metallic behavior.5 Low molecular weight analogues of these polymeric complexes show rich and reversible electrochemical behavior while accepting or donating up to two electrons per metal center.8,9 However, these coordination compounds have been studied only scarcely in lithium-ion batteries. Recently, Wada et al. investigated a Ni-based MOF synthesized from hexaaminobenzene (HAB) as a cathode material for LIBs.10 In the voltage range of 2.0−4.5 V vs Li/Li+, the material showed a specific capacity of 155 mA h g−1 at 10 mA g−1 current rate, corresponding to a specific energy density of 434 W h kg−1. At © 2019 American Chemical Society

Received: April 6, 2019 Revised: June 24, 2019 Published: June 25, 2019 5197

DOI: 10.1021/acs.chemmater.9b01366 Chem. Mater. 2019, 31, 5197−5205

Article

Chemistry of Materials

CN-carbons in NiTIB is larger compared to the CN-carbons of the free DCD, which is expected for the ligand groups binding a cation. The chemical shift for the CH-carbons of NiTIB is in the higher field compared to the corresponding atoms of the free ligand, indicating the increase in the electron density at these sites. The MAS ssNMR spectrum of CuTIB provided no useful information since it appeared as a single broad line due to the presence of paramagnetic Cu2+ ions. Indeed, the electron spin resonance (ESR) spectrum revealed a broad signal with a gfactor value of 2.0685, which is matching that reported for similar Cu2+ coordination complexes (Figure S2).18,19 Fouriertransform IR (FTIR) spectra of NiTIB and CuTIB were similar (Figure 1b), which indicates that these materials share the same structural pattern. According to powder X-ray diffraction (XRD), NiTIB formed a crystalline structure, whereas CuTIB was amorphous (Figure 1c). The XRD pattern of NiTIB is in good agreement with the data reported for the TAB-based nickel 1D polymer.15 Trace amounts of Cu2O were observed in the CuTIB sample, which indicates that Cu2+ was partially reduced by TAB during the synthesis. The content of Cu2O was 2 ± 0.5 wt %, as determined by the XRD method of standard additions (Figure S3, see the Experimental Section for details). The elemental analysis of both materials revealed the discrepancy with the theoretical composition of (MC6H6N4)∞, which was more prominent in the case of CuTIB (Table S1). In particular, the nitrogen content was lower than expected, which was likely due to the partial hydrolysis of NH-groups. The discrepancy might also be due to the relatively low degree of polymerization, especially for CuTIB. No chlorine or sulfur was found in the synthesized materials. Scanning electron microscopy (SEM) was used to reveal the morphology of the solid powders of both coordination polymers. The NiTIB particles appeared as ∼200 nm long and 20−40 nm wide filaments (Figures 1d and S4), whereas

tens of thousands of cycles, as well as cathode materials with the energy density exceeding 600 W h kg−1.

■

RESULTS AND DISCUSSION Synthesis and Characterization. The synthesis of the coordination polymers was performed following the previously reported routes5,15 and is described in detail in the Experimental Section below. Briefly, TAB·4HCl and NiCl2 or CuSO4 aqueous solutions were mixed, concentrated aqueous ammonia was added, and the mixtures were heated in air, which caused the formation of dark blue (for Ni) or dark brown (for Cu) precipitates (Scheme 1). The resulting substances, denoted as NiTIB and CuTIB below, were virtually insoluble in water and any common organic solvents. Scheme 1. Synthesis of NiTIB and CuTIB

The solid-state magic angle spinning NMR (MAS ssNMR) spectrum of NiTIB showed two major 13C signals at 93.70 and 168.81 ppm (Figure 1a), which correspond to the CH- and CN-carbons in the molecular structure of the polymer, respectively. The observed chemical shift values agree well with those reported for the low molecular weight analogues of NiTIB produced in the reactions of an oxidized form of TAB (3,6-diimino-1,4-cyclohexadiene-1,4-diamine, DCD) and Ni(acac)2.17 The free DCD ligand shows three signals at 99.35, 147.3, and 165.15 ppm in 13C MAS ssNMR,17 indicating that the imino and amino groups are spectrally well distinguishable. In NiTIB, all NH-groups are equivalent, i.e., each NiTIB repeating unit can be represented as a superposition of two resonance structures (Figure S1). The chemical shift for the

Figure 1. Characterization of the synthesized materials. (a) MAS ssNMR spectrum of NiTIB; (b) FTIR spectra of the as-synthesized NiTIB and CuTIB; (c) XRD patterns for the as-synthesized NiTIB and CuTIB; (d) SEM images of NiTIB before (left image) and after (right image) ballmilling, scale bars: 200 nm; (e) SEM images of CuTIB before (left images) and after (right images) ball-milling, scale bars: 5 μm (upper images) and 200 nm (bottom images). 5198


Article


Figure 2. Electrochemical behavior of NiTIB in the voltage range of 0.8−2.0 V vs Li/Li+. (a) Charge−discharge curves for different cycles at 50 mA g−1; (b) cycling stability and Coulombic efficiency at 50 mA g−1; (c) CVs for different cycles at 0.1 mV s−1 scanning rate; (d) cycling stability for different cycling modes.

Anode Mode. At the first cycle in the 0.8−2.0 V range, NiTIB showed a large discharge capacity that was mostly irreversible (Figures 2a,b and S7a). Such behavior is typical for many anode materials, including carbons,21,22 SnO2,23,24 GeO2,25 silicon,26 SiO2,27 and organic polymers.28,29 The irreversible processes most likely include solid electrolyte interphase (SEI) formation and substitution of NH hydrogen atoms with lithium, which is discussed below. After 40 cycles, the discharge capacity was 193 mA h g−1 and the Coulombic efficiency was stabilized at around 100% (Figure 2b). Cyclic voltammetry (CV) of NiTIB in the anode mode (Figure 2c) confirmed the domination of irreversible processes at the initial discharge. Subsequent cycles revealed two reduction peaks at 1.15−1.20 and 1.28−1.30 V as well as two oxidation peaks at 1.40−1.42 and 1.48−1.50 V vs Li/Li+. The intensities of the 1.28−1.30 and 1.40−1.42 V peaks were increasing in the first five cycles and then stabilized. At high current rates, the capacity of NiTIB dropped at the first cycles and then recovered back after a number of cycles (Figure 2d). This effect was more pronounced with increasing the current rate. In the extreme case of 20 A g−1, the discharge capacity was 107 mA h g−1 at the first cycle, dropped to nearly zero for the next ∼3000 cycles, and then slowly increased and reached the second maximum (56 mA h g−1) after ca. 10 800 cycles (Figure 2d). The results obtained in the high current rate charging/ discharging experiments indicate the formation of a highly conductive phase during the initial lithiation. At high currents, only a small fraction of the initially nonconductive NiTIB is available for lithiation. According to our hypothesis, the initial transformation of NiTIB to some lithiated product increases the overall conductivity of the electrode, thus improving its capacity. The formation of a more conductive environment involves a continuously increased fraction of NiTIB in the redox process, thus converting the whole material to the active and conductive state. To support this hypothesis, the evolution of conductivity during the charging/discharging was monitored using staircase

CuTIB formed irregular agglomerates up to 5 μm in size (Figures 1e and S4). The conductivity of NiTIB and CuTIB measured for pressed pellets using impedance spectroscopy was below 10−8 S cm−1 at room temperature. However, one would expect higher conductivities for such coordination compounds considering the literature data, e.g., metal complexes with TAB as a ligand were shown to have efficient π-electron delocalization.17 On the one hand, low conductivity of NiTIB and CuTIB might be explained by the presence of high concentrations of defects, which are known to nullify the intrinsic properties of the materials. For example, disordered HAB-derived MOFs show insulating behavior, although they are conductive in highquality crystals.5 On the other hand, the one-dimensional nature of the polymers reported here can also impair the conductivity compared to the two-dimensional MOFs that were shown to be highly conductive. Unlike HAB-derived MOFs, which show metallic behavior,5 pristine NiTIB and CuTIB are semiconductors, which is suggested by their UV− vis spectra, indicating relatively large optical bandgaps of 1.19 eV for NiTIB15 and 1.61 eV for CuTIB (Figure S5). Electrochemical Behavior. One of the main techniques for enhancing the electrochemical performance is nanosizing, which alleviates ion and electron transport and increases specific capacity and rate capability for the materials with low conductivity.20 For this reason, NiTIB and CuTIB were subjected to ball-milling, which led to a decrease in the particle size, most notably for CuTIB (Figure 1d,e and Figure S4). No substantial changes in the FTIR spectra of the ball-milled samples were observed compared to the as-synthesized ones (Figure S6), which indicates that no considerable chemical decomposition occurred under the ball-milling conditions. For the investigation of the designed coordination polymers as battery materials, we assembled half-cells using lithium as the counter-electrode and 1 M LiPF6 solution in organic carbonates as the electrolyte. The cell characterization was performed within the 0.8−2.0 V (anode mode) and 1.5−4.1 or 2.0−3.8 V potential ranges (cathode modes). 5199


Article


element, which is related to the diffusion within the particles and might be described using the Warburg impedance formula31

potentiostatic electrochemical impedance spectroscopy (SPEIS). In the initial state (open-circuit voltage 2.94 V), the Nyquist plot of a NiTIB half-cell contained a highfrequency semicircle, which is attributed to the charge transfer at the lithium metal anode, and a segment of a much larger low-frequency semicircle associated with NiTIB high charge transfer resistance (Figure 3a).30 Upon a slow potential

Z W = AW (1 − i)ω−1/2

(1)

where ZW stands for the impedance, i is the unit imaginary number, ω is the radial frequency, and AW is the Warburg coefficient, which is inversely proportional to the square root of the diffusion coefficient.32 Since |ZW| increase at a fixed ω is equivalent to the AW growth, we might conclude that the diffusion coefficient is reducing upon delithiation. A reverse sweep from 2.0 to 0.8 V leads to the Warburg element shrinking, indicating that the diffusion characteristics were restored. To improve the performance at high current rates, we carried out the first discharge at 200 mA g−1. Subsequent cycling at 20 A g−1 showed that the capacity depression was substantially relieved. The second capacity maximum at 83 mA h g−1 appeared after 4300 cycles (Figure 2d) instead of 10 800 cycles when no preconditioning at low current density was applied. Another trend for NiTIB was the improvement in cycling stability with increasing the current rate (Figure 2d). At 20 A g−1, the discharge capacity was 48 mA h g−1 after 20 000 cycles, which corresponded to 86% of the maximal capacity value. If the first discharge preconditioning was performed at 200 mA g−1, 59 mA h g−1 retained after 20 000 cycles at 20 A g−1, which was 71% of the maximal value observed after 4300 cycles. Such stability characteristics are among the best ever reported for high-rate LIBs. No short-circuiting of the cells was observed even after 20 000 cycles, which is an indicator of lithium dendrite growth suppression. This suppression is likely due to the heat-induced self-healing of lithium, as recently shown by Li et al.33 We supposed that one of the sources of irreversible discharge capacity at the first several cycles was substitution of the NH hydrogen atoms with lithium since these atoms are relatively mobile and acidic. Such substitution was proposed

Figure 3. Impedance spectra of (a, b) NiTIB and (c, d) CuTIB cells during the initial discharging and charging.

decrease, especially below 1.6 V, the radius of this lowfrequency semicircle reduced dramatically, indicating the drop in the active material charge transfer resistance (Figure 3a). During the subsequent potential sweep from 0.8 to 2.0 V, a linear region with a 45° angle became more pronounced in the Nyquist plots (Figure 3b). This is referred to as the Warburg

Figure 4. Electrochemical behavior of CuTIB in the voltage range of 0.8−2.0 V vs Li/Li+. (a) Charge−discharge curves for different cycles at 50 mA g−1; (b) cycling stability and Coulombic efficiency at 50 mA g−1; (c) CVs for different cycles at 0.1 mV s−1 scanning rate; (d) cycling stability for different cycling modes. 5200


Article

Chemistry of Materials previously for NH-containing organic anodes28 and it is naturally expected considering the well-known reactions of NH compounds (e.g., iPr2NH) with lithium forming N-lithiated derivatives (e.g., iPr2NLi commercially available as LDA) and releasing molecular hydrogen. To support this hypothesis, the electrodes were treated with a large excess of lithium diisopropylamide (LDA), which is a strong non-nucleophilic base.34 Indeed, cyclic voltammograms showed that two sets of well-resolved reduction/oxidation peaks with nearly equal intensities appeared already after the first reduction for the LDA-treated sample, in contrast to the untreated NiTIB, where similar CV profiles were observed only after several cycles (Figure S8). This observation confirms that N−H groups of the coordination polymer are most likely transformed into N− Li at the initial cycles, which can be also considered as the origin of the improved electronic and ionic conductivity of the electrode. Similar to NiTIB, the Coulombic efficiency for CuTIB stabilized at ∼100% after 40 cycles (Figure 4). The discharge capacity at the 40th cycle was 98 mA h g−1, which was two times lower than that for NiTIB at the same current density. Copper(I) oxide, which was detected as a minor impurity in CuTIB samples by XRD, is known to be an anode material that shows a plateau at 1.0−1.2 V vs Li/Li+ at the first discharge in the galvanostatic mode and a sharp peak in the CV, which is attributed to Cu2O transformation into copper.35,36 According to the CV, no pronounced oxidation/reduction peaks were observed for CuTIB (Figure 4c). The data reported herein show no signatures of the Cu2O reduction, indicating that its presence might be neglected when considering the electrochemical properties of CuTIB. Overall, the electrochemical behavior of CuTIB was well matching the aforementioned results obtained for NiTIB, which implies similar redox chemistry for both coordination polymers. Particularly, the capacity fade in the first cycles was more pronounced at increased current rates. Impedance spectroscopy showed the formation of a conductive phase upon the initial lithiation when the potential was decreased to 0.8 V (Figure 3c). Subsequent charging to 2.0 V led to an increase in the low-frequency charge transfer resistance and the growth of Warburg coefficient (Figure 3d), which is equivalent to the diffusion coefficient decrease. It is notable that CuTIB showed better cycling stability at relatively low current rates (5 A g−1 and less) compared to NiTIB (Figure 4d). CuTIB-based electrodes maintained a discharge capacity of 55 mA h g−1 after 20 000 cycles at 5 A g−1, which was 2.75 times higher than the corresponding value for NiTIB. This result points to wide opportunities to tailor such fundamental properties of the designed coordination polymers as the specific capacity and cycling stability by modifying their chemical structures via using different metal ions and, potentially, also a variety of organic ligands. Cathode Modes. The capacity of CuTIB in the 1.5−4.1 V range was up to 262 mA h g−1 at 50 mA g−1, which corresponded to a specific energy density of 616 W h kg−1 (Figures 5a,b and S9). These values are higher than those typically reported for classical materials, such as LiCoO2 or LiFePO4 (∼560 W h kg−1),37 and are larger than those for the previously reported hexaaminobenzene-based MOF (155 mA h g−1/434 W h kg−1 at 10 mA g−1).10 At the same time, the capacity of the material dropped after 25 cycles in the galvanostatic mode to about 25 mA h g−1. Similar to the anode mode, increasing the current rate while

Figure 5. Electrochemical behavior of CuTIB in the cathode modes. Charge−discharge curves for different cycles at 50 mA g−1 and specific discharge capacities in (a, b) 1.5−4.1 V and (c, d) 2.0−3.8 V ranges.

charging/discharging the battery stabilized the electrochemical performance of the material. At 1 A g−1, the maximal capacity was 145 mA h g−1, and 50 mA h g−1 was retained after 150 cycles (Figure 5). To improve the cycling stability at lower current rates, we narrowed the operating voltage range to 2.0−3.8 V vs Li/Li+ to suppress the irreversible structural transformations and parasitic reactions occurring at low or high voltages. Indeed, the capacity after 150 cycles at 50 mA g−1 increased to 50 mA h g−1 (Figure 5c,d). The maximal capacity of 155 mA h g−1 was observed at 50 mA g−1, which corresponds to the energy density of 400 W h kg−1. NiTIB demonstrated lower capacities in the cathode mode compared to CuTIB (Figure S10). At 50 mA g−1 in the 1.5− 4.1 V range, the capacity reached 155 mA h g−1, corresponding to 337 W h kg−1. Limiting the voltage range to 2.0−3.8 V considerably improved the NiTIB cycling stability. After 100 cycles at 200 mA g−1, the capacity was 40 mA h g−1, corresponding to 101 W h kg−1. The maximal capacity and energy density values of 98 mA h g−1 and 250 W h kg−1, respectively, were obtained for NiTIB-based cathodes at 50 mA g−1 in the 2.0−3.8 V range. XPS Studies. X-ray photoelectron spectroscopy (XPS) of the CuTIB electrode in the initial state showed a set of peaks typical for Cu(II), with the Cu 2p3/2 signal at 934.4 eV and strong satellites at 938−948 eV (Figure 6a).38 An additional peak at 932.3 eV corresponds to Cu(I),38 which is attributed to Cu2O impurity detected by XRD (Figure 1c). The content of Cu(II) decreased during lithiation, which is evident from the spectra of the electrodes discharged to 1.5 and 0.8 V vs Li/Li+ (Figure 6a). For the electrode discharged to 0.8 V, a single peak with no apparent satellites was observed, which indicates that the reduction of Cu(II) to Cu(I) was virtually complete. The N 1s peak of CuTIB shifted from 399.6 eV in the initial state to 399.2 eV for the electrode discharged to 0.8 V; the same shift is observed for the sample discharged to 1.5 V (Figure 6b). Hence, both copper and the ligand participate in the electrochemical processes during lithiation. For the CuTIB electrode charged to 4.1 V, the N 1s peak appears at 400.3 eV (Figure 6b), which is higher compared to 5201


Article


In other words, the CuTIB polymer backbone is truly ambipolar and, therefore, can accommodate both negative and positive charges at low and high cell potentials, respectively. XPS revealed no obvious changes for the NiTIB electrode discharged to 1.5 V vs Li/Li+. For the sample discharged to 0.8 V, the XPS signals for both nickel and nitrogen were too weak to be analyzed, which is likely due to the SEI formation at the active material surface. After charging to 4.1 V, the Ni binding energy increased compared to the initial state with the Ni 2p3/2 peak shifting from 856.2 to 856.7 eV. This shift should be related to the positive charging of the polymer backbone compensated by PF6− anions as revealed for CuTIB. Similar behavior was observed for the Ni hexaaminobenzene-based MOF charged above 4 V vs Li/Li+ in a lithium-ion battery.10 Proposed Charge−Discharge Mechanisms. Based on the data provided in this paper, the charge−discharge mechanisms can be proposed for NiTIB and CuTIB as outlined in Scheme 2. The observations supporting the proposed scheme are summarized below. 1. Both polymers get p-doped at high potentials and ndoped at low potentials, as follows from XPS and from charge−discharge capacities in anode and cathode modes. 2. At low potentials, NH-protons get substituted with Li, as follows from the irreversible capacities at the first cycles and from the fact that the polymer treated with a strong base (LDA) behaves similar to the cycled polymer. 3. In the anode mode, NiTIB reversibly accepts 2Li+, as follows from the reversible discharge capacity values and CV profiles with two distinct redox peaks in the 1.1−1.5 V vs Li/Li+ range. 4. In the anode mode, CuII in CuTIB reduces to CuI, as follows from XPS; this reduction is observed at a relatively high potential (1.5 V vs Li/Li+), thus, it might be suggested that CuI is present in all states of the anode mode (0.8−2.0 V vs Li/Li+), and hence the backbone stays negatively charged.

Figure 6. X-ray photoelectron spectra of the electrodes in the initial state, charged (4.1 V) state and discharged (0.8 and 1.5 V) states featuring the core levels of (a) Cu 2p and (b) N 1s (b) of the CuTIB as well as (c) the Ni 2p core level of NiTIB. Solid lines show fit of the experimental points. Contributions of individual Cu(II) and Cu(I) peaks are shown as dotted lines in (a).

the initial state. It indicates that CuTIB can undergo oxidation at high potentials and form relatively stable cation species similar to the polymeric aromatic amines.38 The positive charge on the polymer backbone is balanced in that case by PF6− anions. While discharging from 4.1 to 1.5 V, the reduction of the positively charged backbone is followed by its lithiation at lower potentials.

■

CONCLUSIONS We have shown that easily accessible transition metal coordination polymers derived from the 1,2,4,5-tetraaminobenzene ligand represent highly promising anode and cathode materials for lithium-ion batteries. Cycling the materials in the

Scheme 2. Proposed Charge−Discharge Mechanisms of NiTIB and CuTIB

5202


Article

Chemistry of Materials cathode modes (1.5−4.1 or 2.0−3.8 V vs Li/Li+) reveals moderate to inferior stability, although high capacity (up to 262 mA h g−1) and energy density (up to 616 W h kg−1) values can be reached. Further improvements might be achieved via a rational structural design of new coordination polymers probably involving a broader range of metal ions and ligands. The designed coordination polymers exhibit remarkably stable cycling in the voltage range of 0.8−2.0 V vs Li/Li+. The capacity depression observed at the beginning of the cell cycling at high current rates can be mitigated by applying preconditioning at lower currents to generate more conductive lithiated phase of MOFs. This conductive state ensures ultrafast redox reaction kinetics, making it possible to approach impressive specific capacities at ultrahigh current densities (up to 83 mA h g−1 at 20 A g−1). It should be emphasized that the electrodes based on the designed MOFs can withstand tens of thousands of cycles at high current rates (>100C) without severe capacity losses, which allows one to consider them as advanced materials for a new generation of ultrahigh rate batteries.

■

balls. The materials were milled with isopropanol for 30 min at 1000 rpm. The solvent was then vacuum-evaporated. Electrochemical Characterization. To prepare the working electrodes, active materials, Super P and poly(vinylidene difluoride) with the mass ratio of 4:5:1 were thoroughly mixed with Nmethylpyrrolidone to form a homogeneous slurry, which was then tape-casted onto Al foil, dried, and calendered at room temperature. The composite mass loading was around 1.2 mg cm−2. CR2032-type coin cells were assembled in an Ar-filled glove box. A lithium disk was used as the counter-electrode, 1 M LiPF6 solution in a 1:1 v/v mixture of ethylene carbonate and dimethyl carbonate (EC/DMC, anhydrous, Sigma-Aldrich) was used as the electrolyte, and glass fiber (Whatman GF/A Glass microfiber filters, GE Healthcare) was used as the separator. For all cells, cycling was started with discharging unless stated explicitly. Galvanostatic measurements were performed using Neware BTS-4000 stations. CVs were measured with a BioLogic VMP3 or Elins P-20X8 potentiostat. Specific capacity was determined based on the active material mass. Energy density was calculated by integration of the specific capacity vs voltage discharge curves. Impedance Spectroscopy. Two-probe impedance spectroscopy for pelletized pristine materials was carried out at room temperature using Elins P-40X with an FRA-24M module in the 500 kHz−100 mHz frequency range. SPEIS measurements were performed with a BioLogic VMP3 station in the 1 MHz−10 mHz frequency range. The voltage amplitude was 10 mV. The first sweep to 0.8 V was done in 40 steps, subsequent sweeps from 0.8 to 2.0 V and backward were done in 20 steps. Before each impedance spectrum was measured, the cells were conditioned for 15 min; the spectra were recorded with drift correction. Electrode Treatment with LDA. To prepare the LDA solution, 200 μL of 2.5 M n-butyllithium in hexane and 200 μL of dry diisopropylamine were mixed in 5 mL of dry THF at −10 °C under an Ar atmosphere. Electrodes with a total active material mass of 2.3 mg were introduced into the resulting solution for 3 h, extensively washed with hexane, and dried in an Ar-filled glove box. XPS Electrode Preparation and Measurements. The electrodes were discharged to 1.5 or 0.8 V vs Li/Li+ at 50 mA g−1. For the charged electrodes, discharging to 1.5 V was carried out before charging to 4.1 V at 50 mA g−1. The cells were disassembled in an Arfilled glove box, the electrodes were extensively washed with anhydrous dimethyl carbonate and dried in argon at room temperature. XPS measurements were performed using a PHI XPS Versaprobe 5000 spectrometer (ULVAC Physical Electronics) or a Phoibos 100 (Specs) hemispherical analyzer. An Al or Mg source was used, and the energy resolution was 0.5 or 0.8 eV. A dual-channel neutralizer was used to compensate for the local charging of the sample under study due to the loss of photoelectrons. Samples showed varying degrees of charging, which was compensated by calibrating the spectra to the C 1s peak at 284.8 eV.

EXPERIMENTAL SECTION

Materials Synthesis. In a typical synthesis, nickel chloride (1.36 mmol) or copper sulfate (1.36 mmol) dissolved in 20 mL of water was added to 1,2,4,5-tetraaminobenzene tetrahydrochloride (1.36 mmol) solution in 200 mL of water. Concentrated aqueous ammonia (4.5 mL) was added and the mixture was stirred in air at 65−70 °C for 1.5 h. The solvent was removed in a vacuum and the residue was extensively washed with water and then with acetone using a Soxhlet extractor. The resulting product was oven-dried at 100 °C for 1 h. Characterization. FTIR spectra were measured using a Perkin Elmer Spectrum BX system with KBr-based pellets. The UV−vis− NIR absorption spectrum was recorded in a nitrogen atmosphere using an AvaSpec-2048-2 UV−VIS fiber-optic spectrometer integrated with an MBraun glove box. XRD patterns were recorded using a Bruker D8 ADVANCE diffractometer with Cu Kα1 radiation. To determine the content of Cu2O in the CuTIB sample, the fixed amounts of Cu2O were added to the sample, and the integral intensity of the most intensive Cu2O peak (hkl = 111) was measured. The mass of the CuTIB sample was fixed for all XRD measurements, and the measurement conditions were constant. The Cu2O content was calculated by extrapolating the resulting linear dependence to a zero signal (standard addition method). Chemical composition was analyzed by express gravimetry using an Elementar vario MICRO cube. The ESR spectrum of CuTIB was recorded using an Adani CMS8400 spectrometer. The ssNMR spectroscopy experiments were recorded on a Bruker AVANCE III spectrometer operating at 101 MHz for 13C nuclei, using a 3.2 mm MAS probe at 22 ± 1 °C. The chemical shifts were referenced to tetramethylsilane at 0 ppm by adjusting the CH2 signal of solid adamantane spinning at 8 kHz to 38.48 ppm. Conventional cross-polarization experiments were used for the 13C NMR spectra acquisition with the spinning rate of 16 kHz. For the crosspolarization, a ramped RF field from 70 to 100% was applied on 1 H, and the 13C channel RF field was matched to obtain the optimal signal. The duration of 1H π/2 pulse was 2.5 μs and the contact time was 250 μs. SPINAL-64 decoupling was applied during the 20 ms acquisition with an RF amplitude of 100 kHz. The delay between the scans was set to 5 s and the number of scans was 8192. SEM measurements were carried out using a Hitachi SU8000 fieldemission scanning electron microscope (FE-SEM). Images were acquired in a secondary electron mode at 20 kV accelerating voltage and at a working distance of 8−10 mm. Ball-Milling. Milling was performed using a Fritsch Pulverisette 7 premium line system equipped with zirconia jars and 1 mm zirconia

■

ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.9b01366. Elemental composition of NiTIB and CuTIB, resonance structures for the NiTIB repeating unit, the ESR spectrum of CuTIB, data for determining the Cu2O content in CuTIB, SEM images, UV−vis−NIR spectrum for CuTIB, FTIR spectra before and after ball-milling, and extended electrochemical data (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Roman R. Kapaev: 0000-0002-2753-8666 5203


Article


(12) Feng, D.; Lei, T.; Lukatskaya, M. R.; Park, J.; Huang, Z.; Lee, M.; Shaw, L.; Chen, S.; Yakovenko, A. A.; Kulkarni, A.; Xiao, J.; Fredrickson, K.; Tok, J. B.; Zou, X.; Cui, Y.; Bao, Z. Robust and conductive two-dimensional metal−organic frameworks with exceptionally high volumetric and areal capacitance. Nat. Energy 2018, 3, 30−36. (13) Gao, G.; Zheng, F.; Pan, F.; Wang, L.-W. Theoretical Investigation of 2D Conductive Microporous Coordination Polymers as Li−S Battery Cathode with Ultrahigh Energy Density. Adv. Energy Mater. 2018, 8, No. 1801823. (14) Mahmood, J.; Kim, D.; Jeon, I.-Y.; Lah, M. S.; Baek, J.-B. Scalable Synthesis of Pure and Stable Hexaaminobenzene Trihydrochloride. Synlett 2013, 24, 246−248. (15) Cheng, X.-F.; Shi, E.-B.; Hou, X.; Shu, J.; He, J.-H.; Li, H.; Xu, Q.-F.; Li, N.-J.; Chen, D.-Y.; Lu, J.-M. 1D π-d Conjugated Coordination Polymers for Multilevel Memory of Long-Term and High-Temperature Stability. Adv. Electron. Mater. 2017, 3, No. 1700107. (16) Hua, W.; Xiu, J.; Xiu, F.; Zhang, Z.; Liu, J.; Lai, L.; Huang, W. Micro-supercapacitors based on oriented coordination polymer thin films for AC line-filtering. RSC Adv. 2018, 8, 30624−30628. (17) Audi, H.; Chen, Z.; Charaf-Eddin, A.; D’Aléo, A.; Canard, G.; Jacquemin, D.; Siri, O. Extendable nickel complex tapes that reach NIR absorptions. Chem. Commun. 2014, 50, 15140−15143. (18) Veber, S. L.; Fedin, M. V.; Fokin, S. V.; Sagdeev, R. Z.; Ovcharenko, V. I.; Bagryanskaya, E. G. EPR Study of Ligand Effects in Spin Triads of Bis(o-Semiquinonato)Copper(II) Complexes. Appl. Magn. Reson. 2010, 37, 693. (19) Chaudhuri, P.; Verani, C. N.; Bill, E.; Bothe, E.; Weyhermüller, T.; Wieghardt, K. Electronic Structure of Bis(oiminobenzosemiquinonato)metal Complexes (Cu, Ni, Pd). The Art of Establishing Physical Oxidation States in Transition-Metal Complexes Containing Radical Ligands. J. Am. Chem. Soc. 2001, 123, 2213−2223. (20) Wagemaker, M.; Mulder, F. M. Properties and Promises of Nanosized Insertion Materials for Li-Ion Batteries. Acc. Chem. Res. 2013, 46, 1206−1215. (21) Hu, J.; Li, H.; Huang, X. Electrochemical behavior and microstructure variation of hard carbon nano-spherules as anode material for Li-ion batteries. Solid State Ionics 2007, 178, 265−271. (22) Wang, Z.-L.; Xu, D.; Wang, H.-G.; Wu, Z.; Zhang, X.-B. In Situ Fabrication of Porous Graphene Electrodes for High-Performance Energy Storage. ACS Nano 2013, 7, 2422−2430. (23) Jahel, A.; Ghimbeu, C. M.; Monconduit, L.; Vix-Guterl, C. Confined Ultrasmall SnO2 Particles in Micro/Mesoporous Carbon as an Extremely Long Cycle-Life Anode Material for Li-Ion Batteries. Adv. Energy Mater. 2014, 4, No. 1400025. (24) Gurunathan, P.; Ette, P. M.; Ramesha, K. Synthesis of Hierarchically Porous SnO2 Microspheres and Performance Evaluation as Li-Ion Battery Anode by Using Different Binders. ACS Appl. Mater. Interfaces 2014, 6, 16556−16564. (25) McNulty, D.; Geaney, H.; Buckley, D.; O’Dwyer, C. High capacity binder-free nanocrystalline GeO2 inverse opal anodes for Liion batteries with long cycle life and stable cell voltage. Nano Energy 2018, 43, 11−21. (26) Li, X.; Gu, M.; Hu, S.; Kennard, R.; Yan, P.; Chen, X.; Wang, C.; Sailor, M. J.; Zhang, J.-G.; Liu, J. Mesoporous silicon sponge as an anti-pulverization structure for high-performance lithium-ion battery anodes. Nat. Commun. 2014, 5, No. 4105. (27) Meng, J.; Cao, Y.; Suo, Y.; Liu, Y.; Zhang, J.; Zheng, X. Facile Fabrication of 3D SiO2@Graphene Aerogel Composites as Anode Material for Lithium Ion Batteries. Electrochim. Acta 2015, 176, 1001−1009. (28) Wu, J.; Rui, X.; Long, G.; Chen, W.; Yan, Q.; Zhang, Q. Pushing Up Lithium Storage through Nanostructured Polyazaacene Analogues as Anode. Angew. Chem., Int. Ed. 2015, 54, 7354−7358. (29) Wu, J.; Rui, X.; Wang, C.; Pei, W.-B.; Lau, R.; Yan, Q.; Zhang, Q. Nanostructured Conjugated Ladder Polymers for Stable and Fast

Ivan S. Zhidkov: 0000-0001-8727-4730 Keith J. Stevenson: 0000-0002-1082-2871 Pavel A. Troshin: 0000-0001-9957-4140 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We thank Dr Sergey Vasiliev (IPCP RAS, Chernogolovka) for measuring solid-state NMR spectra, Dr Aleksey Galushko (IOC RAS, Moscow, Russia) for SEM experiments, Dr Rina Takazova and Prof. Anastasia Buyanovskaya (INEOS RAS, Moscow, Russia) for elemental analysis, and Olga Kraevaya (IPCP RAS, Chernogolovka) for measuring the ESR spectrum. We also thank Prof. Valery Traven for his continuous support of this work. This work was financially supported by the Russian Science Foundation (Grant No. 16-13-00111). XPS measurements were supported by the Ministry of Education and Science of Russia (task 3.7270.2017/8.9) and FASO (Theme “Electron” No. AAAA-A18-118020190098-5).

■

REFERENCES

(1) Thackeray, M. M.; Wolverton, C.; Isaacs, E. D. Electrical energy storage for transportationapproaching the limits of, and going beyond, lithium-ion batteries. Energy Environ. Sci. 2012, 5, 7854− 7863. (2) Armand, M.; Tarascon, J.-M. Building better batteries. Nature 2008, 451, 652. (3) Tang, Y.; Zhang, Y.; Li, W.; Ma, B.; Chen, X. Rational material design for ultrafast rechargeable lithium-ion batteries. Chem. Soc. Rev. 2015, 44, 5926−5940. (4) Simon, P.; Gogotsi, Y.; Dunn, B. Where Do Batteries End and Supercapacitors Begin? Science 2014, 343, 1210−1211. (5) Dou, J.-H.; Sun, L.; Ge, Y.; Li, W.; Hendon, C. H.; Li, J.; Gul, S.; Yano, J.; Stach, E. A.; Dincă, M. Signature of Metallic Behavior in the Metal−Organic Frameworks M3(hexaiminobenzene)2 (M = Ni, Cu). J. Am. Chem. Soc. 2017, 139, 13608−13611. (6) Sheberla, D.; Sun, L.; Blood-Forsythe, M. A.; Er, S.; Wade, C. R.; Brozek, C. K.; Aspuru-Guzik, A.; Dincă, M. High Electrical Conductivity in Ni3(2,3,6,7,10,11-hexaiminotriphenylene)2, a Semiconducting Metal−Organic Graphene Analogue. J. Am. Chem. Soc. 2014, 136, 8859−8862. (7) Sun, L.; Campbell, M. G.; Mircea, D. Electrically Conductive Porous Metal−Organic Frameworks. Angew. Chem., Int. Ed. 2016, 55, 3566−3579. (8) Noro, S.-i.; Chang, H.-C.; Takenobu, T.; Murayama, Y.; Kanbara, T.; Aoyama, T.; Sassa, T.; Wada, T.; Tanaka, D.; Kitagawa, S.; Iwasa, Y.; Akutagawa, T.; Nakamura, T. Metal−Organic Thin-Film Transistor (MOTFT) Based on a Bis(o-diiminobenzosemiquinonate) Nickel(II) Complex. J. Am. Chem. Soc. 2005, 127, 10012−10013. (9) Herebian, D.; Bothe, E.; Neese, F.; Weyhermü ller, T.; Wieghardt, K. Molecular and Electronic Structures of Bis-(odiiminobenzosemiquinonato)metal(II) Complexes (Ni, Pd, Pt), Their Monocations and -Anions, and of Dimeric Dications Containing Weak Metal−Metal Bonds. J. Am. Chem. Soc. 2003, 125, 9116−9128. (10) Wada, K.; Ken, S.; Sono, S.; Hiroshi, N. Multielectron-Transferbased Rechargeable Energy Storage of Two-Dimensional Coordination Frameworks with Non-Innocent Ligands. Angew. Chem., Int. Ed. 2018, 57, 8886−8890. (11) Park, J.; Lee, M.; Feng, D.; Huang, Z.; Hinckley, A. C.; Yakovenko, A.; Zou, X.; Cui, Y.; Bao, Z. Stabilization of Hexaaminobenzene in a 2D Conductive Metal-Organic Framework for High Power Sodium Storage. J. Am. Chem. Soc. 2018, 140, 10315− 10323. 5204


Article

Chemistry of Materials Lithium Storage Anodes with High-Capacity. Adv. Energy Mater. 2015, 5, No. 1402189. (30) Nakahara, K.; Iriyama, J.; Iwasa, S.; Suguro, M.; Satoh, M.; Cairns, E. J. High-rate capable organic radical cathodes for lithium rechargeable batteries. J. Power Sources 2007, 165, 870−873. (31) Song, J. H.; Bazant, M. Z. Effects of Nanoparticle Geometry and Size Distribution on Diffusion Impedance of Battery Electrodes. J. Electrochem. Soc. 2013, 160, A15−A24. (32) Rodrigues, S.; Munichandraiah, N.; Shukla, A. K. A review of state-of-charge indication of batteries by means of a.c. impedance measurements. J. Power Sources 2000, 87, 12−20. (33) Li, L.; Basu, S.; Wang, Y.; Chen, Z.; Hundekar, P.; Wang, B.; Shi, J.; Shi, Y.; Narayanan, S.; Koratkar, N. Self-heating−induced healing of lithium dendrites. Science 2018, 359, 1513−1516. (34) Bhaw-Luximon, A.; Jhurry, D.; Spassky, N.; Pensec, S.; Belleney, J. Anionic polymerization of d,l-lactide initiated by lithium diisopropylamide. Polymer 2001, 42, 9651−9656. (35) Zhang, C. Q.; Tu, J. P.; Huang, X. H.; Yuan, Y. F.; Chen, X. T.; Mao, F. Preparation and electrochemical performances of cubic shape Cu2O as anode material for lithium ion batteries. J. Alloys Compd. 2007, 441, 52−56. (36) Fu, L. J.; Gao, J.; Zhang, T.; Cao, Q.; Yang, L. C.; Wu, Y. P.; Holze, R.; Wu, H. Q. Preparation of Cu2O particles with different morphologies and their application in lithium ion batteries. J. Power Sources 2007, 174, 1197−1200. (37) Nitta, N.; Wu, F. X.; Lee, J. T.; Yushin, G. Li-ion battery materials: present and future. Mater. Today 2015, 18, 252−264. (38) Biesinger, M. C.; Lau, L. W. M.; Gerson, A. R.; Smart, R. S. C. Resolving surface chemical states in XPS analysis of first row transition metals, oxides and hydroxides: Sc, Ti, V, Cu and Zn. Appl. Surf. Sci. 2010, 257, 887−898.

5205


Coordination Polymers Derived from 1,2,4,5 ... - ACS Publications

Recommend Documents