Ultrathin Manganese-Based MetalâOrganic Framework Nanosheets

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Ultrathin Manganese-Based Metal-Organic Framework Nanosheets: Low-Cost and Energy-Dense Lithium Storage Anodes with the Coexistence of Metal and Ligand Redox Activities Chao Li, Xiaoshi Hu, Wei Tong, Wensheng Yan, Xiaobing Lou, Ming Shen, and Bingwen Hu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b09363 • Publication Date (Web): 16 Aug 2017 Downloaded from http://pubs.acs.org on August 17, 2017

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Ultrathin Manganese-Based Metal-Organic Framework Nanosheets: Low-Cost and Energy-Dense Lithium Storage Anodes with the Coexistence of Metal and Ligand Redox Activities Chao Li,† Xiaoshi Hu,† Wei Tong,‡ Wensheng Yan,§ Xiaobing Lou,† Ming Shen,† and Bingwen Hu†,* †

State Key Laboratory of Precision Spectroscopy, Shanghai Key Laboratory of

Magnetic Resonance, Institute of Functional Materials, School of Physics and Materials Science, East China Normal University, Shanghai 200062, P. R. China ‡

Anhui Key Laboratory of Condensed Matter Physics at Extreme Conditions, High

Magnetic Field Laboratory, Chinese Academy of Sciences, Hefei 230031, P. R. China §

National Synchrotron Radiation Laboratory, University of Science and Technology of

China, Hefei 230029, P. R. China

Email: [email protected]


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ABSTRACT: We herein demonstrate the fabrication of Mn- and Ni-based ultrathin metal-organic framework nanosheets with the same coordination mode (defined as ‘Mn-UMOFNs’ and ‘Ni-UMOFNs’, respectively) through an expedient and versatile ultrasonic approach, and scrutinize their electrochemical properties as anode materials for rechargeable lithium battery for the first time. The obtained Mn-UMOFNs with structure advantages over Ni-UMOFNs (thinner nanosheets, smaller metal-ion radius, higher specific surface area) exhibit high reversible capacity (1187 mAh g-1 at 100 mA g-1 for 100 cycles), excellent rate capability (701 mAh g-1 even at 2 A g-1), rapid Li+ diffusion coefficient (2.48×10-9 cm2 s-1), and a reasonable charge/discharge profile with low average operating potential at 0.4 V. On the grounds of the low-cost and environmental benignity of Mn metals and terephthalic acid linkers, our Mn-UMOFNs show alluring promise as a low-cost high-energy anode material for futuristic LIBs. Furthermore, the lithiation/delithiation chemistry of Mn-UMOFNs was unequivocally studied by a combination of magnetic measurements, electron paramagnetic resonance and synchrotron-based soft X-ray spectroscopy (O K-edge and Mn L-edge) experiments, which substantiate that both the aromatic chelating ligands and the Mn2+ centers participate in lithium storage. KEYWORDS: metal-organic framework, manganese, ultrathin nanosheets, local environment, rehybridization.



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INTRODUCTION Despite the great success of rechargeable lithium-ion batteries (LIBs) in portable IT devices since commercialized by Sony in 1992 and now in electric vehicles (EVs) such as those from Tesla Motors, LIB technology still confronts formidable challenges regarding high-cost, constrained resource supply (lithium and cobalt), insufficient power density, limited lifespan, and safety obstacles.1-4 From the anode point of view, the currently commercialized graphite has safety issues as the low operational potential (< 0.1 V vs. Li+/Li) provokes the formation of dendritic lithium under high current rates.5,

6

On the other hand, Li4Ti5O12 spinel has ultra-stable

reversibility due to its well-known ‘zero-strain’ behavior, but it has slightly low capacity (175 mAh g-1) and relatively high operational potential (1.55 V vs. Li+/Li), leading to insufficient energy density.7, 8 Hence, pursuing prototype, low-cost, and environmental friendly anode materials with higher capacity and lower operational potential (but not close to the Li plating potential) is still of crucial importance to next-generation LIBs. Metal–organic frameworks (MOFs) are a class of porous crystalline materials based on a network structure consisting of metal ion nodes connected by electron-donating functional groups from organic building blocks.9,

10

Benefitting

from their large ion diffusion tunnels and abundant redox active constituents, MOFs are expected as appealing candidates to address the bottleneck problems faced by many Li-conversion and Li-alloying type electrodes, such as severe volume fluctuation, pile-up of solid-electrolyte interphase (SEI) layers, and sluggish ion


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transport kinetics.11, 12 Hence, it is not accidental that several MOFs were evaluated either as a cathode13-15 or as an anode material for LIBs.16-22 These stimulating studies demonstrate the potential of MOFs as LIB anodes, but the model materials reported hitherto are still insufficient to be applied in futuristic LIBs because they usually display smooth charge/discharge profiles with high average operating voltage, which is not able to sustain high energy density in a commercial full battery. Furthermore, the lithiation/delithiation chemistry of most MOFs-based anode materials is based on either the redox active metal-ions16, 17 or the Li+ intercalation into organic moiety,18-21 which provides a limited theoretical capacity. Hence, it is still very urgent to fabricate MOFs-type anodes with the coexistence of metal and ligand redox activities to boost their lithium storage capability. In the present study, manganese- and nikel-based ultrathin metal–organic framework nanosheets (defined as ‘Mn-UMOFNs’ and ‘Ni-UMOFNs’, respectively) were garnered with 1,4-benzenedicarboxylic acid (H2BDC) liker via an ingenious ultrasonic route and investigated as anode materials in Li-ion coin cells for the first time. Ultrathinning MOFs into two-dimensional (2D) nanosheets should be an effective approach to shorten ion diffusion distance, allow superior electron transport, and create high percentages of coordinatively unsaturated redox-active metal sites, thus improving their rate capabilities.23-26 As preconceived, the more impressive Mn-UMOFNs electrode exhibits a low discharge plateau at ~0.4 V, a high reversible capacity of 1187 mAh g-1, along with superior rate capability and long-term cyclic performance, outperforming all previous reported Mn-based MOFs anodes. It should



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be noted here that the H2BDC ligand is available in abundance from the metabolites of aromatic hydrocarbon oxidation and the recycling of polyethylene terephthalate,27 and Mn is one of the most-desired transition metals for electrodes because it is globally available, inexpensive, and environmental friendly, demonstrating the alluring promise of Mn-UMOFNs to substitute graphite and Li4Ti5O12 as futuristic low-cost energy-dense anode material for LIBs. Furthermore, the detailed lithiation/delithiation mechanism of Mn-UMOFNs was proposed by a combination of magnetic

measurements,

electron

paramagnetic

resonance

(EPR)

and

synchrotron-based soft X-ray spectroscopy (sXAS) experiments. The results substantiate that both the metal centers and the aromatic BDC2- chelating ligands take part in lithium storage, providing a high theoretical capacity of 1392 mAh g-1.

RESULTS AND DISCUSSION Characterizations of Mn-UMOFNs and Ni-UMOFNs. The physical and structural characterizations of the as-prepared Mn-UMOFNs and Ni-UMOFNs are demonstrated in Figure 1 and Figure 2. Figure 1a, 1b present the scanning electron microscope (SEM) images of the as-prepared Mn-UMOFNs sample, which exhibits densely packed and randomly arranged ultrathin nanosheets morphology. Transmission electron microscope (TEM) image in Figure 1d shows that the ultrathin nanosheets are intimately interconnected to form a network microstructure. In comparison, Ni-UMOFNs show significantly different ‘wafer biscuit’-like nanosheet structure, which are also densely accumulated and intimately interconnected (Figure 2a, 2b, 2d). Atomic force microscopy (AFM) analyses were performed to acquire more accurate


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data about the thickness of Mn-UMOFNs and Ni-UMOFNs. As shown in Figure 1c and Figure 2c, the thicknesses of the two samples are in the range of 1.8-7.0 (mostly < 3)

and

8−10

nm,

respectively,

further

disclosing the

ultrathin

lamellar

nanoarchitecture of the obtained Mn-UMOFNs and Ni-UMOFNs. The crystal structures of Mn-UMOFNs and Ni-UMOFNs were investigated via powder X-ray diffraction (PXRD), as displayed in Figure 1f and Figure 2f. The PXRD patterns of Mn-UMOFNs suggest that the as-prepared Mn-UMOFNs are isostructural to the previously reported Mn2(OH)2BDC MOF (space group: C2/m, a= 19.9174(1)Å, b= 3.3617(1)Å, c= 6.3270(1)Å, β= 96.224(1)).28 Ni-UMOFNs have the same coordination modes with Mn-UMOFNs, and the diffraction peaks for Ni-UMOFNs could be unequivocally indexed to the monoclinic Ni2(OH)2BDC MOF (space group: C2/m, a= 19.8413(4)Å, b= 3.3181(1)Å, c= 6.6282(1)Å, β= 96.55(1), V= 409.98 Å3 and Dx= 2.297 g cm-3).29 The crystal structures of Mn-UMOFNs and Ni-UMOFNs are layered and they contain two octahedrally coordinated metallic sites (noted as M1 and M2), as depicted in Figure 1g.28-30 Each M1 is connected to four μ3-OH and two O-carboxylate atoms, whereas each M2 is connected to two μ3-OH and four O-carboxylate atoms. These octahedra are edge/corner connected in the (200) crystallographic plane to form 2D metal layers separated by BDC2- linkers. Subsequent composition analyses by X-ray photoelectron spectroscopy (XPS) conform that the as-prepared Mn-UMOFNs are composed of Mn, O, and C, while Ni-UMOFNs are composed of Ni, O, and C without other impurity elements (Figure S1a, 1b). The high resolution Mn 2p and Ni 2p XPS spectra in Figure S1c and S1d



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demonstrate the existence of Mn2+ and Ni2+, respectively. The C 1s peaks in Figure S1g and S1h are deconvoluted into three peaks at 284.70 (284.70), 286.20 (285.78), and 288.80 (288.51) eV, which can be assigned to C=C, C-O, and O-C=O, respectively. Moreover, the coordinated μ3-OH group in Mn-UMOFNs and Ni-UMOFNs can be clearly characterized in O 1s XPS spectra (Figure S1e, 1f) and Fourier transform infrared spectroscopy (FT-IR) spectra (Figure S2). Beyond that, the element mapping of Mn-UMOFNs from STEM image (Figure 1e) demonstrates uniform distribution of Mn, O, and C elements throughout entire Mn-UMOFNs surface, further proving their uniform composition and the pure phase result drawn by PXRD. Similarly, the uniform composition of Ni-UMOFNs can be proved from the STEM-EDS mapping images in Figure 2e. Then N2 adsorption-desorption isotherms were applied to study the porosity of the obtained Mn-UMOFNs and Ni-UMOFNs, as shown in Figure S3a, 3b. The measured Brunauer–Emmett–Teller (BET) specific surface areas of Mn-UMOFNs and Ni-UMOFNs are 32.65 and 15.04 m2 g-1, respectively. In addition, the corresponding Barrett–Joyner–Halenda (BJH) pore size distributions also present that both Mn-UMOFNs and Ni-UMOFNs have mesopores with a narrow size distribution of 3-25 nm. On the basis of the above observations and analyses, ultrathin Mn-based and Ni-based MOF nanosheets with the same coordination modes were successfully fabricated via a facile ultrasonic approach. There are three points to note: (1) The layered structures of Mn-UMOFNs and Ni-UMOFNs with nanometer thicknesses can provide a short ion/electron transfer path to the internal electroactive sites and high


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surface area for interfacial lithium storage, thus resulting in rapid Li+ diffusion and charge transfer kinetics.31 (2) Ultrathinning of Mn-UMOFNs and Ni-UMOFNs might create coordinatively unsaturated metal sites on the exposed surfaces due to the partially terminated BDC2- linking with surface metal cations, giving rise to enhanced metal redox activity.25, 26 (3) The obtained Mn-UMOFNs have thinner nanosheets and higher specific surface area than Ni-UMOFNs. Hence, Mn-UMOFNs would be expected to provide more coordinatively unsaturated metal sites when compared with Ni-UMOFNs. Besides, the larger specific surface area of Mn-UMOFNs should favor the impregnation and storage of electrolyte ions and the accommodation of volume change during repeated cycling, thus achieving ameliorative rate capability.32, 33 Electrochemical performance evaluation of Mn-UMOFNs and Ni-UMOFNs as the anode materials in LIBs. As a proof-of-concept, the electrochemical performances of Mn-UMOFNs and Ni-UMOFNs in LIBs were evaluated using CR2032 coin cells with lithium disks as the counter and reference electrode. It should be mentioned here that to our knowledge, the monoclinic Mn2(OH)2BDC and Ni2(OH)2BDC MOFs have never been reported as anode materials in LIBs before. Figure 3a present the galvanostatic discharge-charge (GDC) potential profiles between 0.01 and 3 V (vs. Li+/Li) of the obtained Mn-UMOFNs and Ni-UMOFNs electrodes in the initial cycle. The Mn-UMOFNs electrode delivers a significantly lower discharge plateau in comparison to that of Ni-UMOFNs, implying a more favorable anodic characteristic. After electrochemical activation for 10 cycles, the discharge plateau of Mn-UMOFNs drops down to ~0.4 V (vs. Li+/Li) and maintains around this potential during the



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subsequent cycles (Figure 3b). Besides, ~72% of the reversible capacity is delivered at lower than 0.4 V, which is ideal for anode materials, because this feature can effectively prevent Li metal dendrite formation and sustain high energy density of a full battery.34 Cyclic voltammetry (CV) curves for the Mn-UMOFNs electrode at a scan rate of 0.2 mV s−1 (Figure S4a, 4b) show cathodic peaks at much lower potentials both in the 1st scan and the 2nd scan when compared to the Ni-UMOFNs electrode, coinciding well with the GDC profile analyses. The cycling performances of as-fabricated Mn-UMOFNs and Ni-UMOFNs electrodes were firstly evaluated at a current density of 100 mA g-1, as displayed in Figure 3c. The Mn-UMOFNs electrode exhibits superior reversibility and cyclability and a reversible capacity as high as 1187 mAh g-1 is retained after 100 galvanostatic charging/discharging cycles, with Coulombic efficiency (CE) approaching 100%. By contrast, in the case of Ni-UMOFNs, the initial discharge and charge capacities are 1833 and 1226 mAh g−1, respectively, but the reversible capacity rapidly declines and maintains at 546 mAh g−1 after 100 cycles. The relatively small initial CE of Mn-UMOFNs and Ni-UMOFNs (57% and 67%, respectively) could be associated with the inevitable decomposition of electrolyte to form SEI membrane and partial Li+ trapping in active sites, which might be improved via pre-lithiation, nanoengineering, interfacial engineering and so on.35-37 Besides, the gradual rise in capacity of Mn-UMOFNs in the initial 40 cycles should be related with the formation of inorganic SEI film by consuming the electrolyte on the electrode surface, and the reversible growth/dissolution of a polymeric gel-like layer that results from kinetically


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activated electrolyte degradation.38,

39

It should be mentioned here that the large

reversible capacity of Mn-UMOFNs (1187 mAh g-1) is even comparable to the MOF-derived Co3O4 nanosheets (1135 mAh g-1 at 200 mA g-1) reported by Liang et al.,40 demonstrating the promise of MOF ultrathin nanosheets to be used as advanced Li-storage anode materials. Electrochemical impedance spectra (EIS) of the Mn-UMOFNs and Ni-UMOFNs electrodes at the 100th cycle were recorded, as displayed in Figure S5. The much smaller charge-transfer resistance at the electrode-electrolyte interface (Rct) and the steeper Zre–Zim curve at low frequencies of the Mn-UMOFNs electrode illustrate its better Li+ migration kinetics than that of Ni-UMOFNs. Moreover, the ex-situ TEM images of Mn-UMOFNs and Ni-UMOFNs after 100 charge-discharge cycles were taken, as shown in Figure S6a, b. The Mn-UMOFNs were able to maintain the ultrathin naonosheet structure without severe aggregation, which can provide a short ion/electron transfer path to the internal electroactive sites, thus resulting in superior rate capability. By contrast, Ni-UMOFNs present an undiscernable microstructure with inferior SEI layers after cycling, indicating a relatively poor rate capability. Figure 3d presents the rate performances of Mn-UMOFNs and Ni-UMOFNs. The Mn-UMOFNs anode is able to deliver a relatively stable charge capacity of ~1029, ~989, ~935, ~921, ~855, and ~701 mAh g-1 at current densities of 100, 200, 400, 600, 1000, and 2000 mA g-1, respectively. After the high-rate cycling, the reversible capacity of Mn-UMOFNs recovers to ~1332 mAh g−1 as the current rate resumes to 100 mA g−1, demonstrating a good electrochemical stability. By contrast,



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Ni-UMOFNs display much poor capacity retention under the same current density (~717, ~537, ~435, ~355, ~229 mAh g-1 at rates of 200, 400, 600, 1000, and 2000 mA g-1, respectively), accompanying by unstable GDC profiles with high voltage hysteresis (Figure S7b). It should be noted here that the Mn-UMOFNs anode demonstrates a stable GDC profiles with the plateau feature around 0.5 V even at rates higher than 1 A g-1 (Figure S7a). More remarkably, the specific capacities delivered above 0.5 V are almost identical at various current densities, whereas the specific capacities delivered below 0.5 V are significantly different, suggesting that the lithiation/delithiation chemistry of Mn-UMOFNs below 0.5 V is a kinetically limited process. Additional study was then performed at a current density of 1 A g−1 for extended number of charge/discharge cycles, as shown in Figure 3e. The Mn-UMOFNs electrode exhibits a striking durability and reversibility, maintaining an 818 mAh g−1 after 300 cycles while retaining a nearly 100% CE. In comparison, the Ni-UMOFNs electrode also present excellent cyclic stability under such a high rate, but the reversible capacity is significantly lower than that of Mn-UMOFNs (346 mAh g−1 after 300 cycles). To the best of our knowledge, such ultrahigh reversible capacity, outstanding rate capability and long-term cyclic performance of our Mn-UMOFNs outperform all previous reported Mn-based MOFs anode materials.17, 41-43 Besides, considering on the more favorable discharge plateau (0.4 V) compared to that of Ni-UMOFNs, as well as the fact that other previously studied Co/Fe/Cu-based MOFs usually display smooth charge/discharge profiles with a high average operating


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voltage,18, 20-22 our Mn-UMOFNs is more capable to sustain high energy density in a commercial full battery. To more clearly illustrate these results, the comparisons of electrochemical performances and typical charge/discharge profiles between Mn-UMOFNs and other reported MOFs-based anodes are presented in Table S1. It should be underlined here that even when compared with the well-known graphite and Li4Ti5O12 anodes, our Mn-UMOFNs demonstrated here is also more favorable in terms of reversible capacity, energy density, and safety. Lithiation/delithiation mechanism investigation of Mn-UMOFNs and Ni-UMOFNs. In order to probe the changes in the oxidation state of Mn and Ni during cycling, magnetic measurements for both materials were carried out. The discharged and recharged samples were cycled to the desired states-of-charge (SOC) and then retrieved from the cycled cells. The antiferromagnet nature of the obtained Mn-UMOFNs was firstly confirmed via χT versus T plots and M-H curves, as can be seen for the details in Figure S8a, S8b. Magnetic susceptibilities were then considered per gram of electrode and are shown as χelectrode versus T plots in Figure 4. The magnetic susceptibility value of the fully-discharged state (χ100K = 0.207 emu g−1) decreases significantly from that of the pristine electrode consisting of Mn-UMOFNs, acetylene black, and CMC binder (χ100K = 0.580 emu g−1), which suggests the loss of Mn2+ and the occurrence of Mn0. The M-H curves for the fully-discharged Mn-UMOFNs sample imply its weak ferromagnetic ordering at 5 K (Figure S8c), further supporting the occurrence of Mn0 during the discharge process. Upon charging, the crystalline Mn-UMOFNs is not recovered, however, Mn2+ seems to be recovered



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due to the increase in the magnetic susceptibility to χ100K = 0.235 emu g−1. Such phenomenon indicates that Mn2+ ions have different local environment in the pristine and recharged electrode, which has also been observed in the Mn-LCP MOF and MnCO3 systems.17, 44 Similar analyses towards magnetic property can be done for Ni-UMOFNs, as can been seen for the details in Figure S9a, b, c, d. The results also demonstrate the occurrence of Ni0 during the Li+ insertion process, implying a similar lithiation/delithiation chemistry of Ni-UMOFNs and Mn-UMOFNs. To further ascertain the valence changes of Mn during galvanostatic charge-discharge process, electron paramagnetic resonance (EPR) spectra of the cycled Mn-UMOFNs samples were recorded under 2 K. The EPR spectrum of the pristine Mn-UMOFNs (Figure 5a) shows a single Lorentzian line with a g-factor value of 2.003 (the absorption peak is around 321 mT), which is characteristic of antiferromagnetically coupled (A. C.) high-spin Mn2+ ions (S = 3/2).45, 46 During the Li+ insertion process, the EPR signal intensity from Mn2+ (A. C.) rapidly decrease, indicating the loss of antiferromagnetic ordering in Mn-UMOFNs. Besides, the proposed formation of diamagnetic lithium salt upon discharge (most probably lithium carbonate) might also cripple the antiferromagnetic coupling of crystalline Mn-UMOFNs.44 The insert shows an enlarged view of the intermediate samples (0.40d, 0.01d, 0.90c), the intense sharp signal with g-factor value of ~2.00 (~336 mT) stands for the free electrons from acetylene black. It is observed that the EPR lineshape from Mn2+ is broadened with continuous Li+ insertion (0.40d, 0.40 V at discharge), and the characteristic EPR signal from Mn2+ is almost undetected for the


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fully-discharged state (0.01d, 0.01 V at discharge), which indicates the emergence of manganese nanostructures.47, 48 Moreover, two ‘hump’ EPR signals representing hole centers can be observed at ~135 mT and ~218 mT, respectively, which can be only detected at relatively low temperature (Figure 5b).49 Such signals should be derived from the occurrence of electron holes during the reduction process of Mn2+ to Mn0, which might favor the transport of electron and Li-ions. When fully recharged to 3.0 V, the characteristic EPR signal from Mn2+ is recovered but the intensity is much lower than that of pristine Mn-UMOFNs, further implying the change of local environment around Mn2+ ions. EPR studies for the cycled Ni-UMOFNs electrodes were also carried out, as displayed in Figure S10a and S10b. For Ni-UMOFNs, Ni2+ (3d8) EPR signal is not observable, because the Ni2+ of Ni-UMOFNs produces an S =1 ground state of significant zero-field splitting, which is often larger than the available microwave quantum and thus the system is “EPR-silent”.50, 51 However, the Li-ions insertion gives rise to a Dysonian-lineshape EPR absorption signal around ~315 mT with an effective g-factor of ~2.09, corresponding to the delocalized conduct electrons.52 The delocalized conduct electrons released from the Ni 3d-O 2p hybridized orbitals might play an important role in balancing the raised positive charge with continuous Li+ intercalation, which has also been reported for the CoHNta coordination polymer electrode recently.53 However, from EPR point of view, the interconversion between Ni2+ and Ni0 is not possible to be clarified due to the absence of Ni2+ (3d8) and nanosized Ni0 (3d6) EPR signals. Finally, in order to grasp more detailed information about the local environment



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evolution of Mn-UMOFNs and Ni-UMOFNs, we performed the material evolution upon cycling at atomic level by monitoring the ex-situ synchrotron-based soft X-ray spectroscopy (sXAS) evolution for both materials. Figure 6a demonstrates the ex-situ O K-edge sXAS total electron yield (TEY) spectra of the Mn-UMOFNs samples discharged/recharged to the SOC as marked on the galvanostatic electrochemical profile (Figure 6c). The pre-edge region between 530 and 537 eV corresponds to the spectroscopic excitations of O 1s electrons to the hybridized state of O-2p and Mn-3d orbitals, while the two higher main peaks at 539.82 and 543.03 eV stand for the excitations to hybridized states of O-2p and Mn-4sp orbitals.54, 55 The pre-edge region is very sensitive to the local environment evolution around oxygen ions,53, 54 and it is enlarged in Figure 6b for more exhaustive analysis. It is clearly observed that the p1 peak at 534.38 eV is gradually weakened, whereas the p2 peak located at lower energy (533.59 eV) is gradually strengthened with continuous Li+ intercalation; in this manner, the p2 peak becomes the only discernable signal after fully-discharged (0.01d). The pre-edge peak shift to lower photon energy (p1 to p2) can be credited to the rehybridization of O-2p and Mn-3d orbitals and the following decrease of bond covalence.56 During the discharging process of Mn-UMOFNs, the electrons would be readily captured by the high electronegative oxygen atoms, in the meantime, Li-ions would be gradually captured by the oxygen atoms with enhanced electronic density. Furthermore, the p1 peak reappears and becomes the dominant state after charging to 0.29 V, and the spectra are almost identical in the subsequent charge process, suggesting the reversibility of local coordination environment around oxygen ions.


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The O K-edge sXAS TEY spectrum after 10 cycles is similar to the one after initial cycle, further demonstrating the reversibility of the Li+ insertion/extraction process. Mn L-edge sXAS spectra corresponding to the 2p core-hole spin–orbital-coupling split were also recorded, as shown in Figure 6d. The Mn-L sXAS spectra of the Mn-UMOFNs samples after 1 and 10 cycles are similar to that of pristine Mn-UMOFNs, providing further unequivocal evidence that the Mn2+ ions in the octahedral crystal field are reversible upon repeated Li+ insertion/extraction.57, 58 Such finding is also in accordance with the magnetic measurements and EPR analyses results. O K-edge and Ni L-edge sXAS evolution analyses for Ni-UMOFNs were provided in Figure S11. The O K-edge evolution results also prove the rehybridization between O-2p and Ni-3d orbitals and the subsequent decrease of bond covalence upon discharging, and the following recovery of the local environment around oxygen ions upon charging. Likewise, the Ni L-edge TEY spectra confirm the reversibility of octahedral coordinated Ni2+ ions during cycling. Based on the above observations and analyses, the detailed electrochemical redox action mechanism of Mn-UMOFNs was proposed, as depicted in Scheme 1. The Mn-UMOFNs firstly undergo a reduction process between Mn2+ and Mn0 (four Li-ions would be inserted), corresponding to the smooth discharge region above 0.5 V (Figure 3b and Figure S7a), as suggested by EPR analyses. Then, extra Li-ions are gradually inserted to the coordinated oxygen atoms with high electronic density, as well as the benzene rings with delocalized π electrons from Mn-UMOFNs, representing the plateau region below 0.5 V (Figure 3b and Figure S7a). The former



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process is kinetically favored while the later one is kinetically limited, as deduce from Figure S7a. As has been reported before, a benzene ring containing strong electron-withdrawing groups on the edges can theoretically accommodate a maximum of six Li-ions.59 Assuming that the two MnO6 octahedra from Mn-UMOFNs (with two μ3-OH and four O-carboxylate atoms per formula) can additionally accommodate a maximum of six Li-ions, the calculated theoretical capacity of Mn-UMOFNs would be 1392 mAh g-1. The lithiation/delithiation chemistry of Ni-UMOFNs is similar to that of Mn-UMOFNs, corresponding to a theoretical capacity of 1359 mAh g-1. However, the electrochemical performance of Ni-UMOFNs in this work is much worse than that of Mn-UMOFNs. The Li+ diffusion kinetics in the Mn-UMOFNs and Ni-UMOFNs electrodes were further evaluated according to the following formulas:60 D = (R2T2)/(2A2n4F4C2σ2)

(1)

Zre = σω−1/2

(2)

where D, R, T, A, n, F, C, and σ represent the diffusion coefficient (cm2 s-1), gas constant (8.314 J mol-1 K-1), absolute temperature (298 K), electrode area (1.5386 cm2), electron transfer number per formula (16), Faraday constant (96485 C mol-1 ), molar concentration of Li-ions (1×10-3 mol cm-3), and Warburg factor relative to Zre, respectively. The σ-values calculated from the low-frequencies region of EIS (Figure S12) are 14.798 and 38.004, respectively. Hence, the apparent Li+ diffusion coefficients (DLi) for Mn-UMOFNs and Ni-UMOFNs would be 2.48×10-9, 3.77×10-10 cm2 s-1, respectively, further illustrating the much better Li+ diffusion ability of Mn-UMOFNs (6.6-fold larger than that for Ni-UMOFNs). The performance


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inferiority of Ni-UMOFNs over Mn-UMOFNs could be ascribed to the following reasons: (1) Ni2+ has a larger ion radius than Mn2+, resulting in bigger steric hindrance effect in Ni-UMOFNs towards Li+ ions. (2) The obtained Ni-UMOFNs have thicker nanosheets than Mn-UMOFNs, which is supposed to provide longer ion diffusion path and lower percentages of coordinatively unsaturated redox-active Ni2+ sites. Besides, the ultrathin nanosheet microstructure of Ni-UMOFNs cannot be well maintained after cycles (Figure S6), which would lead to the capacity decay. (3) The specific surface area of Ni-UMOFNs is very small (15.04 m2 g-1), which is unfavorable to the infusion and storage of electrolyte ions and the accommodation of volume change during repeated cycling. (4) The b-values of the oxidation and reduction

peaks

for

Ni-UMOFNs

(0.56-0.59) indicate a

relatively small

surface-controlled capacitive contribution to the total capacity when compared with Mn-UMOFNs (Figure S13a, b, c, d).

CONCLUSION In summary, Mn2(OH)2BDC and Ni2(OH)2BDC MOF ultrathin nanosheets (Mn-UMOFNs, Ni-UMOFNs) have been prepared via an expedient and versatile ultrasonic approach and introduced as novel active anode materials for LIBs for the first time. As expected, the obtained Mn-UMOFNs with structure advantages in comparison to Ni-UMOFNs (thinner nanosheets, smaller metal-ion radius, higher specific surface area) exhibit outstanding rate capability and long-term cyclic performance, along with a reasonable charge/discharge profile with low average operating potential. Considering on the low-cost and environmental benignity of Mn



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metals and H2BDC ligands, our Mn-UMOFNs show great promise as a low-cost high-energy anode material for next-generation LIBs. Even when compared with the commercialized graphite and Li4Ti5O12 anodes, our Mn-UMOFNs also shows superiority in terms of reversible capacity, energy density, and safety. Moreover, the detailed electrochemical redox mechanism of Mn-UMOFNs was unraveled by a combination of magnetic measurements, EPR, Mn L-edge and O K-edge sXAS experiments. The results substantiate that both the aromatic BDC2- chelating ligands and the metal centers contribute to lithium storage, providing a high theoretical capacity of 1392 mAh g-1. Further improvement in performance might be possible by finely controlling the thickness of Mn-UMOFNs or grafting carbon matrix into the Mn-UMOFNs network.

EXPERIMENTAL DETAILS Synthesis of Mn-UMOFNs and Ni-UMOFNs. All chemicals and solvents were purchased from commercial suppliers and used directly without further purification. The Mn-UMOFNs was fabricated through an ultrasonic method. Firstly, DMF (34 mL), ethanol (10 mL), and water (10 mL) were mixed into a 100 mL polytetrafluoroethylene (PE) tube. Next, 1.50 mmol 1,4-benzenedicarboxylic acid (H2BDC) (0.50 g) was added to the above solution and dissolved under ultrasonication. Then, 1.50 mmol MnCl24H2O (0.297 g) was added, after the Mn2+ salt was fully dissolved, 1.6 mL trimethylamine (TEA) was injected into the solution quickly. Afterwards, the solution was ultrasonicated for another 5 min to obtain a uniform colloidal suspension. Subsequently, the above colloidal suspension was


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untrasonicated under 40 kHz for 8 h under airtight condition (KH5200E, HECHUANG ULTRASONIC Co., China). Finally, the resultant precipitate was collected by centrifugation and rinsed with DMF and ethanol for 3-5 times, followed by vacuum desiccation at 110 C for 12 h. The preparation process of Ni-UMOFNs was same as that of Mn-UMOFNs, except that 1.50 mmol MnCl24H2O was replaced by 1.50 mmol NiCl26H2O, and the DMF, ethanol, and water used were 42 mL, 6 mL, and 6 mL, respectively. Materials Characterization. Scanning electron microscopy (SEM) micrographs were observed using a Quanta 400 FEG field emission scanning electron microscope operated at 20 kV. Transmission electron microscopy (TEM) images and energy dispersive spectroscopy (EDS) mapping images were taken on a Tecnai G2 F20 electron microscope operated at 200 kV. Atomic force microscopy (AFM) measurements were done using a Veeco-Multimode-V scanning probe microscope. Powder X-ray diffraction (PXRD) patterns were carried out on a Holland Panalytical PRO PW3040/60 Diffractometer, with Cu-Kα radiation source (λ = 1.5418 Å). X-ray photoelectron spectroscopy (XPS) analyses were conducted on an ESCALAB 250Xi X-ray photoelectron spectrometer operating at 150 W, with Al−Kα radiation (hν = 1486.6 eV). Fourier transform infrared spectroscopy (FT-IR) spectra were performed on a Nicolet-Nexus 670 infrared spectrometer. N2 sorption isotherms were obtained at 77 K with an ASAP 2020 Accelerated Surface Area and Porosimetry System, the specific surface area and the corresponding pore size distribution were calculated through Brunauer-Emmett-Teller (BET) and Barrett-Joyner-Halenda (BJH) methods.



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Continuous-wave (CW) X-band EPR spectra were performed at 2-290 K with a Bruker EMX plus 10/12 spectrometer at the Steady High Magnetic Field Facilities, High Magnetic Field Laboratory. Microwave power and modulation amplitude were set to 1 mW and 2 G, respectively. Magnetic measurements were conducted using a Physical Property Measurement System (PPMS, B1500A, Quantum Design). The temperature dependence of magnetization was performed under zero field cooled (ZFC) and field cooled (FC) warming conditions between 5 and 300 K with an applied field of 5000 Oe. The M-H measurements were carried out between -20000 Oe and 20000 Oe at 5 K, 35 K, and 300 K, respectively. The discharged/charged samples at given states-of-charge (SOC) were collected by disassembling the CR2032 coin cells inside Ar-filled glovebox, and then rinsed several times with dimethyl carbonate (DMC). The wet electrodes were dried for 12 h at room temperature in Ar-filled glovebox before packing into the paramagnetic quartz tubes for EPR tests, or the capsules for magnetic measurements. sXAS characterization. Soft X-ray absorption spectroscopy (sXAS) was monitored at National Synchrotron Radiation Laboratory (BL12B-α: MCD). Total electron yield (TEY) spectra at Mn/Ni L-edges and O K-edge were recorded via detecting the electric current caused by excited electrons from the Mn-UMOFNs/Ni-UMOFNs samples at various SOC. The cycled samples at given SOC were also collected by disassembling the coin cells inside Ar-filled glovebox, and then washed several times with DMC. The wet Mn-UMOFNs/Ni-UMOFNs electrodes were dried at room temperature for 12 h in Ar-filled glovebox before sXAS tests.


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Electrochemical

Measurements.

The

electrochemical

properties

of

Mn-UMOFNs/Ni-UMOFNs as anode materials were evaluated using CR2032 coin cells with lithium disks as the counter/reference electrode. The working electrodes were composed of the as-fabricated Mn-UMOFNs/Ni-UMOFNs active materials, acetylene black, and carboxymethyl cellulose (CMC) binder in the weight ratio of 70:25:5. The loading amount of active materials was 1.2-1.5 mg cm-2. The electrolyte was made up with 1 M LiPF6 dissolved in dimethyl carbonate (DMC)/ ethylene carbonate (EC)/ diethyl carbonate (DEC) (1:1:1 vol %) containing 5 wt% fluoroethylene carbonate (FEC). Celgard 2325 membrane was employed as the separator. All the cells were assembled in a Ar-filled glovebox and tested at 25 °C. Galvanostatic measurements were conducted between 0.01 and 3.0 V (vs. Li+/Li) on a LAND-CT2001 cycler. All the specific capacities were calculated based on the total mass

of

Mn-UMOFNs

or

Ni-UMOFNs.

Cyclic

voltammetry

(CV)

and

electrochemical impedance spectra (EIS) were recorded using a CHI 660a electrochemical workstation.

ASSOCIATED CONTENT Supporting Information Available. Supplementary XPS spectra, FT-IR spectra, N2 adsorption-desorption isotherms, TEM images, CV curves, EIS spectra, Galvanostatic charge/discharge voltage profiles, Temperature dependence of χT, M-H traces, Temperature dependence of χelectrode, X-band EPR spectra, O K-edge sXAS TEY spectra, Ni L-edge sXAS TEY spectra, the relationships between Zre and ω-1/2, lg i vs.



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log v plots, and Table. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *(B.H.) E-mail: [email protected]. Notes The authors declare no competing financial interest.

ACKNOWLEDGEMENTS This work is supported by National Natural Science Foundation of China for Excellent Young Scholars (Grant No. 21522303), Large Instruments Open Foundation of East China Normal University. We also acknowledge the support from National Synchrotron Radiation Laboratory (NSRL) for the sXAS experiments, and the Steady High Magnetic Field Facilities of High Magnetic Field Laboratory for the EPR experiments. C. L. is also grateful to ECNU Outstanding Doctoral Dissertation Cultivation Plan of Action (No. YB2016031).

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53. Li, C.; Lou, X.; Shen, M.; Hu, X.; Yan, W.; Zou, Y.; Tong, W.; Hu, B. High-Capacity Cobalt-Based Coordination Polymer Nanorods and Their Redox Chemistry Triggered by Delocalization of Electron Spins. Energy Storage Mater. 2017, 7, 195-202. 54. Yoon, W.; Balasubramanian, M.; Chung, K. Y.; Yang, X.; McBreen, J.; Grey, C. P.; Fischer, D. A. Investigation of the Charge Compensation Mechanism on the Electrochemically Li-Ion Deintercalated Li1-xCo1/3Ni1/3Mn1/3O2 Electrode System by Combination of Soft and Hard X-ray Absorption Spectroscopy. J. Am. Chem. Soc. 2005, 127, 17479-17487. 55. Yoon, W.; Balasubramanian, M.; Yang, X.; Fu, Z.; Fischer, D. A.; McBreen, J. Soft X-Ray Absorption Spectroscopic Study of a LiNi0.5Mn0.5O2 Cathode during Charge. J. Electrochem. Soc. 2004, 151, A246. 56. Yoon, W.; Kim, K.; Kim, M.; Lee, M.; Shin, H.; Lee, J.; Lee, J.; Yo, C. Oxygen Contribution on Li-Ion Intercalation−Deintercalation in LiCoO2 Investigated by O K-Edge and Co L-Edge X-ray Absorption Spectroscopy. J. Phy. Chem. B 2002, 106, 2526-2532. 57. Yu, J. H.; Liu, X.; Kweon, K. E.; Joo, J.; Park, J.; Ko, K.; Lee, D. W.; Shen, S.; Tivakornsasithorn, K.; Son, J. S.; Park, J.; Kim, Y.; Hwang, G. S.; Dobrowolska, M.; Furdyna, J. K.; Hyeon, T. Giant Zeeman Splitting in Nucleation-Controlled Doped CdSe:Mn2+ Quantum Nanoribbons. Nat. Mater. 2009, 9, 47-53. 58. Wang, Y.; Liu, J.; Lee, B.; Qiao, R.; Yang, Z.; Xu, S.; Yu, X.; Gu, L.; Hu, Y.; Yang, W.; Kang, K.; Li, H.; Yang, X.; Chen, L.; Huang, X. Ti-Substituted Tunnel-Type


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Na0.44MnO2 Oxide as a Negative Electrode for Aqueous Sodium-Ion Batteries. Nat. Commun. 2015, 6, 6401. 59. Han, X.; Qing, G.; Sun, J.; Sun, T. How Many Lithium Ions Can Be Inserted onto Fused C6 Aromatic Ring Systems? Angew. Chem. Int. Ed. 2012, 51, 5147-5151. 60. Li, H.; Zhou, L.; Zhang, L.; Fan, C.; Fan, H.; Wu, X.; Sun, H.; Zhang, J. Co3O4 Nanospheres Embedded in a Nitrogen-Doped Carbon Framework: An Electrode with Fast Surface-Controlled Redox Kinetics for Lithium Storage. ACS Energy Lett. 2017, 2, 52-59.



a

g

c

b

6.99 nm 2.56 nm 5.18 nm

2.15 nm 1.82 nm

O-K

C-K

f

Mn-UMOFNs simulated

10

(201)

Mn-K

(200)

e

0.5 um

(-201)

d

500 nm

(011)

1 um

Intensity (a.u.)

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

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20 30 2θ (degree)

40

50

Figure 1. (a, b) SEM micrographs of the obtained Mn-UMOFNs under different magnification. (c) AFM image of the obtained Mn-UMOFNs, showing the ultrathin nature of the nanosheets. (d) TEM image of the as-fabricated Mn-UMOFNs. (e) STEM and the corresponding EDS mapping images of the as-fabricated Mn-UMOFNs, showing the uniform distribution of Mn, O, and C elements. (f) PXRD patterns of the obtained Mn-UMOFNs and the simulated one. (g) Layered crystal structure of Mn/Ni-UMOFNs. Color scheme: Mn/Ni, green; O, red; C, gray (H atoms are omitted for clarity).


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a

c

b

8.36 nm

8.33 nm

9.55 nm

9.78 nm

O-K

C-K

f

Ni-UMOFNs simulated (-201)

Ni-K

(201)

e

0.4 um

(011)

d

300 nm

(200)

3 um

Intensity (a.u.)

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


10

20 30 2θ (degree)

40

50

Figure 2. (a, b) SEM micrographs of the obtained Ni-UMOFNs under different magnification. (c) AFM image of the obtained Ni-UMOFNs, showing the ultrathin nature of the nanosheets. (d) TEM image of the as-prepared Ni-UMOFNs. (e) STEM and the corresponding EDS mapping images of the as-prepared Ni-UMOFNs, showing the uniform distribution of Ni, O, and C elements. (f) PXRD patterns of the obtained Ni-UMOFNs and the simulated one.



2.0 1.5 1.0 0.5

Mn-UMOFNs Ni-UMOFNs

1.5 1.0 0.5 100th

0

2000

c

2000

10th 10th

100

1750

Mn-UMOFNs (discharge) Mn-UMOFNs (charge) Ni-UMOFNs (discharge) Ni-UMOFNs (charge)

1500 1250

60 40

750 500

20 250

100th

0

200 400 600 800 1000 1200 Specific capacity (mAh g-1)

0

20

40 60 Cycle number

e

1500 1250

100

1000

200

100 80

60 400

600

1000 2000

750

40

500 20 250 unit: mA g-1

0 0

20

40 Cycle number

60

0 80

Specific capacity (mAh g )


-1

100

1750

80

1000

2000

2000

Coulombic efficiency (%)

d

100th

2.0

0.0 500 1000 1500 Specific capacity (mAh g-1)

10th 10th

2.5

0.0 0

100th


2.5

3.0

0 100

80

100

1750


1500 1250 1000

80 60 1 A g-1

40

750 500

20 250 0 0

50

100

150 Cycle number

200

250


b

1st

Mn-UMOFNs Ni-UMOFNs

Voltage (V vs. Li+/Li)


3.0

Specific capacity (mAh g-1)

a

Specific capacity (mAh g-1)

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

Page 36 of 40

0 300

Figure 3. (a, b) Galvanostatic discharge-charge potential profiles between 0.01 and 3 V (vs. Li+/Li) of the obtained Mn-UMOFNs and Ni-UMOFNs electrodes in the initial cycle and the 10th/100th cycle, respectively. (c) Cycling performances of Mn-UMOFNs and Ni-UMOFNs at a current density of 100 mA g−1. (d) Rate performances of Mn-UMOFNs and Ni-UMOFNs at various current densities from 100 to 2000 mA g−1. (e) Cycling performances of Mn-UMOFNs and Ni-UMOFNs at a high current rate of 1 A g−1. To activate the electrodes, 100 mA g-1 was applied in the first 10 cycles.


Page 37 of 40

4.0 pristine fully-discharged fully-charged

3.5

χelectrode (emu g-1)

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


3.0 2.5 2.0 1.5 1.0 0.5 0.0 0

50

100

150 T (K)

200

250

300

Figure 4. Temperature dependence of χ for selected samples: pristine electrode fabricated with Mn-UMOFNs, fully-discharged electrode, and fully-charged electrode.



a

Mn2+ (A. C.)

Intensity (a.u.)

Intensity (a.u.)

Mn2+ (A. C.)

acetylene black

0.40d 0.01d 0.90c

0

100

200 300 400 500 Magnetic field (mT)

600

700

pristine 0.40d 0.01d 0.90c 3.00c

0

200 400 Magnetic field (mT)

600

b Intensity (a.u.)

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

Page 38 of 40

acetylene black

2K 30 K 45 K

0

100

200 300 400 500 Magnetic field (mT)

600

700

Figure 5. (a) X-band EPR spectra recorded on the Mn-UMOFNs series at 2 K. Here ‘d’ stands for ‘discharge’, ‘c’ stands for ‘charge’, and the number before stands for the corresponding SOC. The inset shows an enlarged view of the intermediate samples. (b) X-band EPR spectra recorded on the fully-discharged Mn-UMOFNs sample under variable temperature.


Page 39 of 40

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


Figure 6. (a) O K-edge sXAS TEY spectra of the Mn-UMOFNs samples discharged/recharged to the SOC as marked on the galvanostatic electrochemical profile. (b) Enlarged O K-edge sXAS TEY spectra of the pre-edge region. (c) The galvanostatic electrochemical profile of Mn-UMOFNs /Li cycled at a current density of 100 mA g-1. (d) Mn L-edge sXAS TEY spectra of pristine Mn-UMOFNs, and the samples after 1 and 10 galvanostatic charge/discharge cycles.



3.0 Mn-UMOFNs

2.5


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

Page 40 of 40

2.0 Mn0

1.5

Li+ Mn2+

Li+

1.0 0.5

Li+

Li+

0.0 Capacity (mAh g-1)

Scheme 1. The proposed electrochemical redox action mechanism of Mn-UMOFNs. The Mn-UMOFNs firstly undergo a reduction process between Mn2+ and Mn0, corresponding to the smooth discharge region above 0.5 V. Then, extra Li-ions are gradually inserted to the coordinated oxygen atoms with high electronic density, as well as the benzene rings with delocalized π electrons, representing the plateau region below 0.5 V.


Ultrathin Manganese-Based MetalâOrganic Framework Nanosheets

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