Highly Reversible Conversion Anodes Composed of Ultra-Large

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Highly Reversible Conversion Anodes Composed of Ultra-Large Monolithic Grains with Seamless Intragranular Binder and Wiring Network Chenglong Wu, Jiulin Hu, Zhenguo Yao, Dongguang Yin, and Chilin Li ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b07169 • Publication Date (Web): 06 Jun 2019 Downloaded from http://pubs.acs.org on June 7, 2019

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ACS Applied Materials & Interfaces

Highly Reversible Conversion Anodes Composed of Ultra-Large Monolithic Grains with Seamless Intragranular Binder and Wiring Network

Chenglong Wu†,‡, Jiulin Hu†, Zhenguo Yao†, Dongguang Yin‡* and Chilin Li†*

†State

Key Laboratory of High Performance Ceramics and Superfine Microstructure, Shanghai

Institute of Ceramics, Chinese Academy of Sciences, 585 He Shuo Road, Shanghai 201899, China. ‡School

of Environmental and Chemical Engineering, Shanghai University, Shanghai 200444,

China.

Abstract: Conversion anodes enable a high capacity for lithium-ion batteries due to more than one electron transfer. However the collapse of host structure during cycling would cause huge volume expansion and phase separation, leading to the degradation and disconnection of mixed conductive network of electrode. The initial nanostructuring and loose spatial distribution of active species are often resorted to in order to alleviate the evolution of electrode morphology, but at a cost of the decrease of grain packing density. The utilization of ultra-large micro-sized grains of high density as conversion anode is still highly challenging. Here a proof-of-concept grain architecture characterized by endogenetic binder matrix and wiring network is proposed to guarantee the structural integrity of monolithic grains as large as 50-100 μm during deep conversion reaction. Such big grains were fabricated by self-assembly and pyrolysis of a Keggin-type polyoxometalate-based complex with protonated tris[2-(2-methoxyethoxy)-ethyl]amine (TDA-1-H+). The metal-organic precursor can guarantee a firm adherence of numerous Mo-O clusters and nuclei into highly elastic monolithic structure without evident grain boundaries and intergranular voids. The pyrolyzed TDA-1-H+ not only serves as in-situ binder and conductive wire to glue adjacent Mo-O moieties, but also acts as Li-ion pathway to promote a sufficient lithiation on surrounding Mo-O. Such a monolithic electrode design leads to an unusual high-conversion-capacity performance (1000 mAh/g) with satisfactory reversibility (reaching at least 750 cycles at 1 A/g). These cycled grains are not disassembled even

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after undergoing long-term cycling. The conception of intragranular binder is further confirmed by consolidating MoO2 porous network after in-situ stuffing MoS2 nano-binders.

KEYWORDS: Conversion anode, Intragranular binder, Monolithic grains, Polyoxometalate complex, Li-storage

Introduction Although lithium ion batteries (LIBs) were successfully utilized in the fields of portable, transportation and stationary energy storage devices, they still suffer from the limitation of energy density due to the lacking of high-capacity anodes with micro-sized grains or monolithic structures 1. Insertion-type anodes, e.g. graphite, Li4Ti5O12 and Nb2O5, can show good rate performance and long cycling life even without intentional nano-structuring 2-4. However their reversible capacities are still limited and usually less than 300 mAh/g due to single-electron (or fewer electron) transfer. Conversion or alloying anodes enable a much higher capacity close to or exceeding 1000 mAh/g as long as their low conductivity is compensated and volume expansion is suppressed 5,6. Different from Li4Ti5O12 and Nb2O5, the marginal lithiation in conversion materials does not lead to a remarkable upgrade of bulk electron conductivity, whereas the deep lithiation often triggers a phase segregation with the generation of more grain boundaries, which likely passivate the migration of Li-ions to the next electroactive particle

7,8.

Nano-engineering enables the surface non-stoichiometry or defect

enrichment, and likely delays (or suppresses) the occurrence of phase transformation reaction and the evolution of volume. This factor is favorable for the improvement of bulk conductivity and charge transfer across heterostructure. However the discrete distribution of nanoparticles is often at the cost of low tap density for electrode. Typical alloying anode, e.g. Si, even suffer from a more serious volume expansion/contraction (400% vs. 200% for conversion anodes) 9. The pulverization of conversion/alloying electrodes would cause the loss of electrical connectivity or material exfoliation form the current collector, leading to cycling and capacity degradation. In order to address the issues on electrode cracking and delamination, some methods, such as the interposition of excess conductive carbon nanostructures and the employment of highly viscoelastic binders, have been widely attempted

10,11.

The uniform spread-out of active grains on

graphene-like cushion or the well-aligned stringing of active grains by carbon nanotube enables an


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intimate contact with current collector and effective volume buffering 12,13. Some novel binders with self-healing

function

(e.g.

PAA-P(HEA-co-DMA))

and

multilevel

bonding

(e.g.

poly[3-(potassium-4-butanoate) thiophene]) are recently attempted to preserve the structural integrity of conversion anodes

14,15.

However nano-sizing of active species and their dispersive embedment

are still the prerequisites to achieve the desired effects in most the cases. These extraneous electrode additives are usually electrochemically inert, and compromise the energy density to a certain degree. Therefore exploring highly dense anodes with endogenetic wiring network or binder moiety to retard grain disassembly is highly desired. Punching in the bulk of conductive conversion anode (e.g. nitride) appears to a potential solution to preserve the integrity of cycled micro-sized grains 16. But the optimization of spatial pore distribution and pore size sophisticates the synthesis steps. Extrusion of large active grains from highly elastic conductive skeleton but without grain peel-off is another pathway to compact the conductive network of electrode

17.

However high fraction of inactive

conductive network is sometimes unavoidable. Metal-organic frameworks (MOFs) appear to be a type of good precursor to prepare high-performance conversion anodes 18,19. But their products often suffer from serious internal etching (or hollow structuring) to thin the active species or the formation of loose aggregation of nanoparticles during pyrolysis. In this work, we proposed a proof-of-concept electrode characterized by endogenetic binder moieties and wiring network to guarantee the structural integrity of micro-sized monolithic grains, which are prepared by self-assembly and pyrolysis of a Keggin-type polyoxometalate (POM)-based complex, which also contains the moieties of protonated tris[2-(2-methoxyethoxy)-ethyl]amine (TDA-1-H+) 20. The crystallinity and porosity of resulting MoOx products are tunable depending on the pyrolysis temperature. In contrast to traditional MOF materials, after pyrolysis this novel metal-organic precursor can guarantee an effective adherence of numerous Mo-O clusters or MoO2 nanocrystals into monolithic structure even without evident grain boundaries. Therein the pyrolyzed organic ligand not only serves as in-situ binder and conductive wire to glue adjacent MoOx moieties, but also acts as Li-ion pathway to promote a sufficient lithiation on surrounding MoOx. It therefore leads to an unusual high-conversion-capacity performance (1000 mAh/g) with satisfactory reversibility (reaching at least 750 cycles at 1 A/g) for the highly dense monolithic grains, which are not disassembled even after undergoing long-term cycling. The intragranular binder strategy is further confirmed by sulfurating the loose porous structure of MoO2, which enables the repair of



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grain connectivity and shrinkage of pore volume due to the generation of MoS2 nano-wires as intragranular binder.

Results and discussion Preparation of monolithic grains and microstructural analysis. Figure 1 illustrates the synthetic procedure scheme with the corresponding solution photos. Firstly, commercially available TDA-1 and phosphomolybdic acid H3PMo12O40·nH2O were respectively dissolved in deionized water. Then both the transparent solutions were mixed to trigger the formation of a bright yellow turbid liquid, indicating the occurrence of self-assembly or complexation process. The precipitated material was washed for several times by using deionized water and collected by centrifugation, and then it was moved to a vacuum furnace and dried for 24 h at 50°C. The resulting yellow species was moved to a tube furnace for thermal pyrolysis in nitrogen at 650°C or to 900°C, leading to the products termed as POM-650 or POM-900 respectively. POM-900 was also further sulfurized at 600°C to obtain a product termed as POM-CS. The dissolved POM exits in a form of anionic polynuclear oxide containing transition metal of high valence, e.g. Mo6+ in [PMo12O40]3− 17. The POM anion possesses substantial terminal and bridging oxygen atoms, which enable the interaction with protons and ammonium ions in TDA-1-H+ by electrostatic interaction or covalent bonding. Three ether-containing chains in TDA-1-H+ could modulate their configuration for the effect of 2D alignment manner of POMs, which is indicated from the giant sheet morphology of this POM-based complex with a rectangular geometry with width of 2 μm and length of 5 μm

from the scanning

electron microscopy (SEM) images (Figure S1). The thickness direction of sheet grains is perpendicular to the (001) plane. The POM anionic layers are separated by TDA-1-H+ moieties with a spacing of 14.2 Å, and they are usually located at the outer surface of sheets from the reported negative zeta potential in aqueous solution 21. However when the complex powder is transferred to alcohol solution, the good dispersity disappears and these nanosheets are stacked together to form larger-sized aggregates in view of potential evolution of zeta potential (Figure S2). The X-ray diffraction (XRD) pattern of this complex is consistent with previous report on mono-protonated (TDA-1-H)2H[PMo12O40] (Figure S3) 21. After pyrolysis, the XRD patterns of POM-650 and POM-900 are significantly changed depending on the annealing temperature. POM-650 displays a pattern without remarkable diffraction


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peaks, instead some broad bumps appear in the 2-theta ranges of 25-30o and 30-40o (Figure 2a). These broad peaks are close to the positions of (011) and (-211) peaks of MoO2 phase (JCPDS No. 86-0135), indicating a potential nucleation of MoO2 nanodomains. Higher pyrolysis temperature (900oC) causes a drastic growth of MoO2 grains with better crystallinity as seen from the intensive diffraction peaks typically assigned to monoclinic MoO2 phase. The extra sulfuration step further reinforces these peaks with clearer discernment of shoulder peaks. Meantime some minor peaks appear at 14.5°, 33.0° and 58.3°, and they belong to Plane (002), (100) and (110) in hexagonal MoS2 (JCPDS. No.73-1508). This result implies a successful surface (shallow) sulfuration of POM-900 with the formation of MoS2 wires. The poor crystallinity retards the accurate assignment of POM-650 sample exclusively from XRD result. Therefore we adopted X-ray photoelectron spectroscopy (XPS) measurement to obtain more composition and bonding information in POM-650. Figure 2b shows the Mo 3d peaks, where one can still find a substantial fraction of Mo6+ signal from the pronounced shoulder peak located at 235.2 eV for Mo 3d3/2 22. Another shoulder peak at 229.6 eV belongs to Mo4+ 3d5/2 signal, which also implies the nucleation of MoO2 moieties from the Mo6+-O amorphous matrix 12. The C1s peaks stem from the residual of pyrolytic TDA-1-H+ (Figure 2c). They consist of C-C, C-N (or C-P), C-O and carboxylate carbon (O-C=O) components at 284.6, 285.5, 286.5 and 288.5 eV

23.

The signals

belonging to hetero-doped carbon are evident, especially those related to the bonding with O atoms. This phenomenon indicates that TDA-1-H+ organic ligand does not undergo a complete pyrolysis into carbon under the moderate annealing temperature (650oC). The functional O groups are still rich with different bonding states, which are beneficial for the multilevel interaction with active Mo-O species and therefore the alleviation of electrode expansion and pulverization as discussed later. The N 1s spectrum stems from the ammonium ions in TDA-1-H+ (Figure 2d), and could be divided into three different C-N bondings, i.e. graphitic-N, pyrrolic-N and pyridinic-N respectively at 403.1, 400.1 and 398.1 eV

17.

A shoulder peak partially overlaps with N 1s spectrum at 396.3 eV, and is

assigned to the signal of Mo 3p3/2

24.

The O 1s (Figure 2e) spectrum can be fitted into four peaks.

The peaks located at 530.9 and 532.1 eV correspond to O2- in MoO2 and MoO3 25,26, and another two peaks located at 533.1 and 534.1 eV are respectively ascribed to the bondings of C-O and O=C-O 27,28.

This result of O 1s spectrum is in good agreement with those of Mo 3d and C 1s spectra. From



Page 6 of 27

XPS, it is concluded that POM-650 has a mixed valence state of Mo6+ and Mo4+ with the presence of pyrolyzed TDA-1-H+ as in-situ O, N-contained binder due to shallow thermal reduction. Thermogravimetric analysis (TGA) was performed under air condition with a heating rate of 10 °C/min to estimate the carbon content of POM-650 (Figure S4). When the temperature is increased to 300 °C, a weight loss of 4 wt% corresponds to the removal of surface adsorbed water in POM-650. From 300 to 500 oC, there is no evident weight loss. Then the weight loss continues from 500 to 700 oC with a value of 3.5 wt%. In this temperature range, the simultaneous oxidation of C and MoO2 to CO2 (gas) and MoO3 occurs 29. Therefore, the carbon content in POM-650 is estimated to be less than 14.4 wt% in view of the already existing MoO3 in POM-650. After 700°C, the TGA curve drops sharply owing to the sublimation of MoO3 30.

Micro-sized monolithic morphology and porosity evolution. POM-650 displays an unusual dense morphology characterized by ultra-large monolithic grains up to 50-100 μm in size from SEM imaging in Figure 3a-c and S5. These grains are difficult to be further grinded to a smaller scale likely due to the existence of intragranular binder of high mechanical strength (i.e. pyrolytic TDA-1-H+). The surface of these grains is roughly smooth without serious grain boundaries. Although the profile of numerous embedded MoOx nanodomains is discernable, they are compactly linked with each other in a seamless manner. Surprisingly, such a thick electrode does not compromise its lithiation performance as shown later. Higher pyrolysis temperature (900oC) causes a porous evolution inside the thick grains of POM-900, however their monolithic morphology is still preserved from the SEM images in Figure 3d-f and S6. The peel-off of some smaller powder grains from the edge of monolithic ones indicates the degradation of grain mechanical strength. The loose nanostructure with rich porosity is evidently observed from the magnified images. Inside the monolithic aggregate, there are numerous small nanoparticles of 20-30 nm in an incompact connecting manner so as to leave plentiful voids. These nanoparticles should be well-crystallized MoO2 nano-crystals, which enable the oxidation and removal of pyrolytic TDA-1-H+ (or resulting C) template during the deoxygenization of MoO3 into MoO2. Although the mesoporous structure is beneficial for electrolyte infiltration, its volume evolution during conversion reaction would degrade the grain connectivity and interrupt the conductive network as shown in the inferior electrochemistry later. We attempted to consolidate the porous structure by additional sulfuration step on POM-900


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(Figure 3g-i and S7). The resulting POM-CS sample presents a denser morphology compared with POM-900. The size of MoO2 primary particles is increased to 40-60 nm after sulfuration. These MoO2 nanoparticles are firmly embedded into the monolithic matrix consisting of the fillers of MoS2 wires. The healing of open pores in monolithic grains of POM-CS would cause a pronounced boosting of capacity retention as discussed later. Note that the sintered samples are thicker than the sheet-like precursor, indicating that the elimination of surface charge in POM-based complex during pyrolysis would cause drastic deformation and mergence of sheet-like grains. In order to get deeper insights into the microstructure of pyrolyzed samples, transmission electron microscopy (TEM) measurement was performed. Since most the monolithic grains are too thick and unsuitable for TEM characterization, we have to choose some relatively thin regions. Figure 4a-d and S8 show the typical regions of POM-650, which disclose the existence of cluster-like species (MoOx) embedded in gel-like matrix. This amorphous matrix is extended to the whole grain with the formation of conformal coating on the grain surface (Figure 4b). This matrix appears to be elastic and could function as in-situ binder for the maintenance of grain integrity, and should belong to the component of pyrolytic TDA-1-H+. In the thinner region, the nucleation process is prone to be more remarkable under strong e-beam irradiation with the appearance of nanoparticle profile (Figure 4c). These particles are as small as 5-10 nm with evident lattice strips corresponding to the (011) and (-211) planes of MoO2 from high-resolution (HR) TEM imaging (Figure 4d). These lattice strips well agree with the diffraction ring assignment in selected area electron diffraction (SAED), typically belonging to MoO2 phase (Figure 4e). When increasing the annealing temperature to 900oC, the MoO2 particles undoubtedly grow, and their sizes are close to 50 nm near the thinnest edge region under the strong e-beam irradiation for high-revolution imaging (Figure 4f, g and S9). No matrix can be found to confine these particles, and also no thick coating to cover the particle surface as observed from the spread of lattice stripes up to the edge of particle. The absence of carbon coating is ascribed to its oxidation combustion by releasing O from MoO2+x. The assignment of lattice stripes in HRTEM is consistent with the single-crystal-like diffraction plots in SAED of individual particle (Figure 4h), corresponding to the (011) and (-102) planes of MoO2. POM-CS enables the growth of MoS2 nanoribbons from the voids between MoO2 particles of ~50 nm (Figure 4i-k and S10). These few-layer MoS2 nanoribbons are in a thickness of about 4-5 nm (i.e. roughly 7 S-Mo-S layers), and function as wires (or fillers) between MoO2 particles to reinforce the



Page 8 of 27

intragranular transport network. The generation of MoS2 moieties is confirmed by the typical wide lattice stripes of (002) planes in HRTEM and diffraction rings (e.g. corresponding to (002), (100) and (110) planes) in SAED, coexistent with those for MoO2 phase (Figure 4l). The compact property of POM-650 and its porosity evolution at higher preparation temperature are also indicted from Brunauer-Emmett-Teller (BET) analysis (Figure S11). The surface area of POM-650 from BET is as low as 2.50 m2/g, and the corresponding pore volume is also small (0.0088 cm3/g). The porous structure of POM-900 enables a higher surface area (8.83 m2/g) and larger pore volume (0.0689 cm3/g) with a range of pore size concentrated around 25 nm. The in-situ formation of ultrathin MoS2 wires in POM-CS further increases the BET surface area to 10.52 m2/g with a comparable pore volume of 0.0569 cm3/g and pore size distribution around 20 nm. The mesoporosity for the latter two samples is reflected from the more pronounced hysteresis in nitrogen sorption isotherms. The higher BET surface area may be also responsible for the lower initial Coulombic efficiency and earlier capacity degradation in view of excessive accumulation of SEI layer, especially for the case of POM-900.

Conversion electrochemistry based on ultra-large monoliths. Metallic MoO2 with negligible bandgap is expected to be a promising conversion anode considering the facile construction of intragranular transport network, as long as sufficient Li-ion flow is provided from the exposed phase boundaries between Li2O and Mo, which are infiltrated by electrolyte during the propagation of conversion reaction

16.

Its theoretical capacity is as high as 838 mAh/g based on four-electron

reaction between Mo0 and Mo4+. This value is more than twice the theoretical capacity (372 mAh/g) of graphite, a commercialized LIB anode material. Moreover, the high density of MoO2 (6.5 g/cm3) is expected to achieve a higher volumetric capacity than that of graphite anode with a much lower density of 2.3 g/cm3. The main concern for MoO2 arises from the large volume change especially during deep conversion reaction, which would give rise to electrode cracking and even pulverization or delamination. These undesired morphology evolutions may interrupt the pathways of electric contact in MoO2 electrode, therefore resulting in a loss of usable active species and capacity decay. MoO2-based nanomaterials have been reported to display a higher electroactivity than that of their bulk counterparts

31.

The shrinkage of particle size is beneficial for the acceleration of charge

transport and alleviation of volume expansion (or stress) across the whole particle. However, the


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loose adherence of nanoparticles by extraneous binder or their discrete distribution in porous network still cannot guarantee a long-term capacity retention in view of the facile degradation of electrode integrity. MoO2-based nanostructures endowed with both the long cycling (e.g. > 50 cycles) and high capacity performances (e.g. > 800 mAh/g) were rarely reported for LIBs 31-36. Our design of ultra-large MoO2 monoliths with compact endogenetic binding network provides a solution to the dilemma between electrode density and electrochemical reversibility. Figure 5a shows the galvanostatic curves of POM-650 from 0.1 to 3.0 V at a current density of 100 mA/g during charge and discharge processes. Interestingly, the lithiation curve during the first discharge displays the similar profile as the following curves, moreover without serious capacity loss after the first cycle. Such a high reversibility is unusual for conversion-type oxide anodes, and is likely associated with the unique microstructure, crystallinity and spatial distribution of MoO2

31.

The curve replicability continues during the following long-term cycling. These curves are sloped in the whole voltage range, and most of the charge capacity is released below 2 V and the total capacity is as high as 900 mAh/g. The curves of cyclic voltammetry (CV) during the early cycles under a scan rate of 0.1 mV/s are consistent with the galvanostatic curves in terms of curve profile (Figure S12). They are highly overlapped between the first cycle and following ones. The anodic process shows a wide characteristic peak from 1 to 2 V. Note that during the first cathodic process the irreversible peak around 0.7 V is weak compared with previous reports on conversion oxides 12. It indicates the suppression of excess formation of solid electrolyte interface (SEI) in view of the protection of pyrolytic TDA-1-H+ coating in POM-650. The absence of plateau-like curves (or sharp CV peaks) is attributed to the preservation of defect-rich MoOx clusters or nuclei under the moderate-temperature annealing. It would alleviate the irreversible Li trapping in lattices or retard the serious phase transformation as indicated from the highly duplicated curves. Figure 5b displays a long cycling performance of POM-650 with reversible capacity preserved at 915 mAh/g after 230 cycles. The Coulombic efficiency (CE) value is close to 80 % for the first cycle and quickly approaches to and is stabilized at 100 % from the second cycle. The excess contribution of capacity by interfacial Li-storage mechanism cannot be ruled out from the observation of practical capacity exceeding the theoretical one 37,38. The reversible capacity at a much higher current density of 1 A/g is still as high as 600 mAh/g or above up to 575 cycles (Figure 5c). The abnormal capacity increase at the 190th cycle should be caused by the reactivation of accumulated “dead mass”. This reactivation



Page 10 of 27

phenomenon is also observed at the 376th and 607th cycles. After 600 cycles, the capacity is still preserved at 500 mAh/g up to at least 750 cycles. To the best of our knowledge, such a superior retention of capacity under high current density is unprecedented for these ultra-large (50-100 μm) monolithic grains, which can endure deep conversion reaction even without the assistance of excess amount of binder and conductive additive. The prior cycling at lower rate (e.g. 0.1 A/g for the first 10 cycles) is helpful to stabilize the capacity performance after switching to higher rate. Gradual increase of current density brings about a remarkable improvement on capacity retention even when reaching to much higher current density (Figure 5d). The reversible capacities are still as high as 787, 600 and 380 mA h/g respectively at 1, 2 and 5 A/g. Figure S13 compares our electrochemical result with recently published MoO2-based anodes to emphasize the capacity-rate advantage of our electrode even under the condition of large-sized and high-density grains. The grain growth and crystallinity modification in POM-900 and POM-CS trigger the evolution of electrochemical behavior (Figure 5e and f). Both the electrodes display multi-stage plateaus for both the charge and discharge processes during the early cycles. Most the charge plateaus are below 2 V besides the appearance of an additional plateau at 2.25 V for POM-CS. This phenomenon is caused by electrochemical reaction of MoS2 wires. This MoS2-induced plateau makes the charge capacity of POM-CS 100 mAh/g higher than that of POM-900, and it is 600 mAh/g for the former during the early cycles. The discharge plateaus deviate more or less for the first and following discharge processes, indicating an irreversible evolution of microstructure or Li trapping. These phenomena are also indicated from the CV curves at a scan rate of 0.1 mV/s during the first three cycles (Figure S14). There are roughly two sets of redox peaks respectively around 1.3-1.4 V and 1.6-1.75 V for well-crystallized MoO2. Both the corresponding cathodic peaks have additional shoulder peaks at slightly higher voltage, indicating that lithiation process likely undergoes more phase transformation steps than the corresponding delithiation process. The typical cathodic peaks in the first discharge have a negative displacement compared with the following discharge processes. Therein a broad irreversible peak around 0.72V is also observed, corresponding to the formation of SEI. After sulfuration, the SEI-induced peak shifts to 0.5 V, and meantime an additional irreversible peak appears at 1.1 V. This phenomenon likely indicates a component modulation of SEI due to the catalysis of surface MoS2. In POM-CS, another couple of redox peaks for MoS2 is also highly reversible around 2-2.25 V as those for MoO2 at least during the early cycles. The progress of


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cycling would amorphize MoO2 as indicated by the blurring of plateaus and appearance of sloped curves as in the case of POM-650 (Figure 5e and f). The combination of amorphization and porous architecture can boost the electrolyte wettability and make the reversible capacity climb to 1100-1200 mAh/g (under 100 mA/g). This capacity value has far exceeds the theoretical value of MoO2 (838 mAh/g) in view of a larger contribution from interfacial storage, stemming from the rich porous structure as well as newly generated grain boundaries induced by grain cracking. However the loose porous structure of POM-900 does not keep the electrode integrity under repeated lithiation. The capacity fades quickly after 40 cycles and is kept at 400 mAh/g after 80 cycles (Figure 5g and h). In contrast, POM-CS enables a much better capacity retention with a high reversible capacity of 1100-1200 mAh/g up to 90 cycles. Afterwards, the degradation of capacity is also slow and a capacity of 600 mAh/g is still preservable after 150 cycles. The delayed degradation of capacity benefits from the relatively robust architecture by filling MoS2 wires. These MoS2 nanowires also bring about a progress of rate performance, and the capacities are around 700 and 500 mAh/g respectively at 1 and 2 A/g, which are 200 mAh/g higher compared with POM-900 under the corresponding current densities. The charging under high rate does not compromise the capacity recoverability and retention after 40 cycles. The increase of wired sulfide fraction seems to be promising to further boost the capacity retention, in view of looser nanostructure and more wiring binder amount. However the sulfuration effect is difficult to be deepened from this POM-900 oxide precursor, whereas the facilely sulfurated precursor (e.g. dissoluble Mo-based salts) usually does not lead to a compact morphology consisting of high-density grains after excess sulfuration. On the other hand, excess sulfuration would in turn cause the passivation of electrode network involving substantial less conductive MoS2. It likely degrades the capacity/rate retention.

Discussion on microstructure and kinetics. Resorting to large-sized active grains is crucial to achieve high density electrodes and boost the energy density of LIBs. It is more challenging for conversion-type grains, which often suffer from serious cracking and pulverization during long-term cycling. The disintegration of electrode network can be alleviated by stuffing excess supramolecular binder and conductive additive, especially the highly elastic or self-healing ones

10,11.

However the

introduction of non-active components and redundant porous network would undoubtedly compromise the energy density of LIBs. This extraneous binder strategy has been confirmed to be



effective for insulating Si or less conductive oxides (e.g. Fe2O3), even for micro-sized Si

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14,15.

However the micro-size range is roughly 1-3 μm. Ultra-large grains are never reported to be successful. In this work, an integrated electrode design is proposed with active species, binder and conductive additive all within single grain as large as 50-100 μm. These coarse grains are surprisingly electrochemically reversible even under deep conversion process and high rate. Pyrolytic TDA-1-H+ with rich O-groups is the key component to cement the electroactive MoOx clusters or nuclei, leading to the pulverization difficulty at least during mechanical milling. This endogenetic binder can also serve as fine wiring network to transport electron and ion to surrounding Mo-O moieties inside the individual grain. It only takes a small weight less than 14.4 wt%. The limitation of grain growth and crystallization in MoOx enables a seamless coexistence of active species nanodomains and binder wires. No cracking in POM-650 grains is observed after long 160 cycles at a high current density of 1 A/g (Figure 6a-c), confirming the superior robustness of these ultra-large grains. The coarsing of primary nanoparticles (40~60 nm) within single grain does not influence their compact stacking, because no open pores or voids are left. In order to explore the product distribution after deep discharge (to 0.1 V), we have to focus on the thin region for HRTEM (Figure 6d-g). We find numerous nanoparticles (dark contrast) with a size smaller than 20 nm embedded in pyrolytic TDA-1-H+ matrix (light contrast) after 33 cycles under 100 mA/g. Most the nanoparticles are still compactly stacked with visible phase boundaries, and they correspond to Mo metal and Li2O nanodomains from the lattice strips of (110) plane for Mo and (111) plane for Li2O. These conversion products are also detectable from SAED pattern. Although these products are very fine, they are firmly arrested by the gel-like pyrolytic TDA-1-H+ framework to construct a self-built conductive network and maintain the integrity of monolithic grain. The kinetic advantage of POM-650 over POM-900 and POM-CS is also indicated from the electrochemical impedance results of pristine and cycled electrodes (Figure 7). The corresponding Nyquist plots consist of both the high-frequency and middle-frequency semicircles, respectively representing the SEI film (Rf) and charge transfer (Rct) resistances. The inclined straight line in low frequency region is relevant to a finite-length Warburg impedance (Zflw), which reflects both the diffusion and accumulation of Li (Figure 7a and b)

39.

The intercept value before the beginning of

semicircle corresponds to bulk electrolyte resistance (Re). These spectra are well fitted by suitable equivalent circuit with corresponding elements listed in Figure 7c. POM-650 enables the smallest


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resistance values for both the Rf (6.3 ) and Rct (22.7 ) in pristine Li/POM-650 cell, whereas those are as high as 73.1 and 61.2  for POM-900 as well as 105.6 and 220  for POM-CS. The very small SEI film resistance for POM-650 is in accordance with the suppression of SEI accumulation and initial irreversibility. The attenuation of SEI layer is beneficial for the acceleration of charge transfer across electrode-electrolyte interface and compact conductive network,40 leading to a much smaller Rct for POM-650 than for POM-900 and POM-CS. Note that the interface of POM-CS is more resistant than POM-900 due to the growth of MoS2 wires, which are less conductive than metallic MoO2. However the slight increase of Rf and Rct does not compromise the Li-storage performance. The cycling retention and rate performance are on the contrary better for POM-CS, indicating that the consolidation of electrode network via MoS2 binders and compression of pore volume play more important roles than conductivity of active species. After 20 cycles, the Rf and Rct for Li/POM-650 cell are still as low as 3.4 and 45.4  respectively. The preservation of electrode integrity (compactness) as well as the inhibition of excess grain boundaries and intergranular voids are responsible for the low interface resistance. The Rf values are still correspondingly higher for cycled POM-900 (93.2 ) and POM-CS (88.4 ) than cycled POM-650. The exceeding of Rf of POM-900 over POM-CS after cycling indicates that the loose (porous) electrode network in the former is prone to be degraded with a more accumulation of SEI. However therein the charge transfer across MoO2 particles is remarkably improved from the decrease of Rct value of POM-900 (11.5 ), which is even smaller than that of cycled POM-650. This phenomenon is likely caused by the more exposure of discrete MoO2 particles to the reaction front after the break-down of porous network. Such large-sized monolithic grains were also attempted as Na-storage conversion anodes. Figure S15 shows the performance comparison between POM-650 and POM-CS for Na-storage in diethylene glycol dimethyl ether (DGM) with 1 M NaSO3CF3. Although their Na-storage capacities are lower than the corresponding Li-storage ones because of sluggish mass transport of Na-ions with larger size across multiphase interfaces, Na-driven rate and cycling performances are not bad especially for POM-650. The reversible capacity ranges from 300 to 200 mAh/g under 0.1 A/g, and is preserved at 130, 100 and 70 mAh/g respectively under 0.5, 1 and 2 A/g. After 180 cycles, the reversible capacity is still stabilized at 150 mAh/g with a CE of 96%. POM-CS displays a worse capacity performance with capacity fading from 250 to 70 mAh/g after 120 cycles. These results



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imply that seamless intragranular binder and wiring network also takes positive effect on activating the Na-driven conversion reaction in compact electrode network. We expect that this grain-binding strategy can be extended to more electrode systems also with potentially large volume expansion, e.g. other metal oxides, sulfides and nitrides, provided that their metal-contained precursors can bond with organic ligands to form large-sized or monolithic complex. These complexes should not undergo severe pulverization during the following pyrolysis and the resulting active species should be intrinsically conductive enough. The popular Si anode however appears to be difficult to be tailored by this grain-binding strategy, since the Si-based complex is prone to be oxidized during pyrolysis and the Si-based material is much less conductive. The extra reduction and etching steps cannot guarantee the adhesion of Si particles with each other.

Conclusion In summary, a proof-of-concept conversion electrode architecture characterized by endogenetic binder and wiring network is proposed. It is realized by self-assembly and pyrolysis of a Keggin-type polyoxometalate (POM)-based complex with protonated tris[2-(2-methoxyethoxy)-ethyl]amine (TDA-1-H+), leading to the production of highly dense monolithic grains as large as 50-100 μm. Therein the pyrolyzed TDA-1-H+ function as both the in-situ binder and conductive wire to maintain the structural integrity and morphology stability of monolithic grains during deep conversion reaction. This binder matrix also acts as Li-ion pathway to promote a sufficient lithiation upon surrounding Mo-O clusters and nuclei, which are glued together without evident grain boundaries and inter-granular voids. Such a coarse grain architecture can achieve an unusual high-conversion-capacity performance (1000 mAh/g under 100 mA/g) with excellent reversibility (reaching at least 750 cycles under 1 A/g). This strategy of constructing intra-grain binder can be extended to broader material systems especially for expansible conversion and alloying electrodes.

Experimental section Preparation of POM-based complex: Phosphomolybdic acid hydrate (H3PMo12O40·nH2O, TCI Reagent Co., denoted as POM) of 821.3 mg and tris[2-(2-methoxyethoxy)-ethyl]amine (C15H33NO6, Alfa Aesar 95%, denoted as TDA-1) of 145.5 mg were dissolved in deionized water of 30 mL respectively. The PH value of TDA-1 solution


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was adjusted to 1.5 by adding 6 M aqueous HCl, which is very important for the formation of polyoxometalate-based complex with sheet-like morphology. Both the solutions underwent ultrasonication step until the clear solutions were obtained. Then, the TDA-1 aqueous solution was mixed with the POM aqueous solution under continuous stirring. The mixed solution evolved into a bright yellow suspension immediately, indicating an occurrence of self-assembly or complexation reaction. After keeping the suspension stir for 2 h, the POM-based complex powder was collected by centrifugation (8000 rpm, 5 min) and was washed several times using pure water. This powder was dried for 24 h at 50 °C in a vacuum oven, followed by hand-milling to get light yellow one. Preparation of POM-650, POM-900 and POM-CS: Typically, for the preparation of POM-650, the POM-based complex was moved into a ceramic holder for calcination for 3h at 650 °C under nitrogen atmosphere based on a heating rate of 3 °C/min. The resulting black powder consisting MoOx clusters (or nuclei) and pyrolyzed TDA-1-H+ was collected and denoted as POM-650. Similarly, POM-900 (made up of MoO2 nanocrystals) was synthesized by annealing POM-based complex at 900 °C for 3h under nitrogen atmosphere based on a heating rate of 3 °C/min. To fabricate POM-CS, POM-900 was milled with sulphur powder, followed by calcination for 10 h at 600 °C under vacuum based on a heating rate of 2 °C/min. At last, the black product containing MoO2 nanocrystals wired by MoS2 was collected and denoted as POM-CS. Physical characterization: The phase assignment of precursor and sintered products was determined by X-ray diffraction (XRD, Philips PW3710, 40 kV/30 mA) with a scanning rate of 2.0°/min in a 2-theta range from 10° to 70° by using Cu Kα radiation. The grain morphology of POM-based complex, annealed products and cycled electrode was observed by scanning electron microscope (SEM, Carl Zeiss Leo 1530VP Gemini). The phase assignment and microstructure situation of POM-based complex, annealed products and cycled electrode were characterized by transmission electron microscopy (TEM, JEOL JSM-6700F) equipped with component of selected area electron diffraction (SAED). The surface elemental and bonding situations of POM-650 were investigated by X-ray photoelectron spectroscopy (XPS, ESCAlab-250). Thermogravimetric analysis (TGA, DSC 800, PerkinElmer) was performed under oxygen atmosphere with a heating rate of 10 °C/min. Gas sorption analyzer



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(Quantachrome Autosorb) was used to obtain nitrogen adsorption/desorption isotherms and the corresponding pore size distribution. Electrochemical characterization: The charge-discharge measurements of POM-650, POM-900 and POM-CS electrodes were run based on coin cells of CR2025 type by using Li (or Na) foils as anodes. During the fabrication of working electrodes, the electroactive grain powder was mixed with conductive additive (Super-P carbon) and electrode binder (polyvinylidene fluoride PVDF dissolved in N-methyl-2-pyrrolidinone, NMP) based on a weight ratio of 7:2:1. The resulting mixture was then ground into uniform slurry manually, which was pasted on the current collector of Cu sheet and was dried in a vacuum oven overnight at 60 °C. The electrode loading is ~1.5 mg/cm2 based on the weight of active species. Celgard 2400 (Whatman) was adopted as the separator. The commercial electrolyte (Sigma Aldrich) for Li cells consists of Li-salt LiPF6 of 1 M dissolved in a mixed solvent of diethyl carbonate (DEC) and ethylene carbonate (EC) based on a volume ratio of 1:1. The electrolyte for Na cells consists of Na-salt NaSO3CF3 of 1.0 M dissolved in diethylene glycol dimethyl ether (DGM). These cells were assembled in a glovebox filled with argon, and therein the contents of water and oxygen are controlled below 0.1 ppm. The galvanostatic measurement was done at room temperature between 0.1 and 3 V in a range of current density from 0.1 to 5 A/g on a battery testing system (LAND-CT2001A). Cyclic voltammetry (CV) was performed at 0.1 mV/s based on a voltage range between 0.1 and 3.0 V on an electrochemical workstation (VersaSTAT3, AMETEK Scientific Instruments). The impedance measurement of POM-650, POM-900 and POM-CS cells at the pristine and cycled stages was carried out in a frequency range between 100 kHz and 0.01 Hz on a Solartron frequency analyzer (1260−1296).

Support information SEM and XRD of pristine POM-based complex precursor, TGA of POM-650, SEM, TEM and BET of POM-650, POM-900 and POM-CS, CV curves of POM-650, POM-900 and POM-CS, capacity performance comparison of POM-650 with previous MoO2-based anodes, Na-storage performance of POM-650. This material is available free of charge via the Internet at http:// pubs.acs.org.


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Corresponding Authors *Email: [email protected] (C. L.). *Email: [email protected] (D. Y.).

Acknowledgments This work was supported by National Key R&D Program of China (2016YFB0901600), National Natural Science Foundation of China (51772313, U1830113), “Hundred Talents” Program of Chinese Academy of Sciences and“Thousand Talents”Program of Shanghai.

Competing interests The authors declare no conﬂict of interest.

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Figure 1. Scheme of synthesis steps for POM-650, POM-900 and POM-CS with corresponding photos of solutions and suspension liquid. The synthesis steps consist of chelation, pyrolysis and sulfuration processes. The solutions of TDA-1-H+ (transparent) and H3PMo12O40·nH2O (yellow) were mixed to trigger the chelation process, followed by the precipitation and collection of yellow powder of POM-TDA complex. The powder pyrolyzed at 650°C or 900°C is denoted as POM-650 and POM-900, which was further sulfurized at 600°C to obtain POM-CS. POM-650 displays roughly dense morphology without serious grain boundaries, and therein MoOx clusters or nanodomains are compactly linked with each other in a seamless manner. POM-900 consists of porous monolithic grains composed by numerous well-crystallized MoO2 nanoparticles. POM-CS also consists of MoO2 nanoparticles but wired by MoS2 nanoribbons as nano-binder to stuff the voids between MoO2 particles.


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(a)

(b)

(d)

(c)

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Figure 2. (a) XRD patterns of POM-650, POM-900 and POM-CS. The standard patterns of monoclinic MoO2 and hexagonal MoS2 are also listed as references. POM-650 displays a pattern without remarkable diffraction peaks, instead some broad bumps appear. POM-900 and POM-CS display intensive diffraction peaks typically assigned to MoO2 phase. For POM-CS some minor peaks belong to MoS2 are observable. XPS spectra with their fitting profiles of (b) Mo 3d, (c) C 1s, (d) N 1s and (e) O 1s. POM-650 has a mixed valence state of Mo6+ and Mo4+ with the presence of pyrolyzed TDA-1-H+ as in-situ O, N-contained binder due to incomplete thermal reduction.



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(a)

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Figure 3. (a) Overview SEM image of POM-650, displaying an unusual morphology of monolithic grains with a size up to 50-100 μm. (b,c) SEM images of POM-650 in difference scales, showing roughly smooth and dense surface without serious grain boundaries. (d) Overview image of POM-900, displaying the peel-off of some smaller grains from the edge of monolithic ones. (e,f) SEM images of POM-900 in different scales, showing porous microstructure consisting of numerous nanoparticles as small as 20-30 nm with the preservation of monolithic morphology. (g) Overview image of POM-CS, still displaying the monolithic morphology with the alleviation of grain peel-off from edge compared with POM-900. (h,i) SEM images of POM-CS in difference scales, showing a relatively more compact surface with nanoparticles (40-60 nm in size) firmly embedded into the monolithic matrix compared with POM-900.


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Figure 4. (a) Overview TEM image of POM-650. (b) TEM image of POM-650 with a typical region disclosing the embedment of cluster-like species (MoOx) in gel-like matrix. (c) TEM image of POM-650 with a thinner region showing a nucleation phenomenon with the appearance of nanoparticle profile. (d) HRTEM image and (e) SAED pattern of POM-650 with typical lattice stripes and diffraction rings assigned to MoO2 phase. (f) TEM image of POM-900 showing the well-profiled MoO2 nanoparticles. (g) HRTEM image and (h) SAED pattern of an individual particle in POM-900 with the lattice stripes and diffraction plots corresponding to single-crystal-like MoO2. (i,j) TEM images of POM-CS in different scales disclosing MoS2 nanoribbons growing from the voids between MoO2 particles of ~50 nm. (k) HRTEM image and (l) SAED pattern of POM-CS, indicating the coexistence of MoS2 and MoO2 from the assignment of lattice stripes and diffraction rings.



(a)

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(h)

Figure 5. (a) Galvanostatic curves of POM-650 in a voltage range of 0.1-3.0 V at 100 mA/g depending on different cycling stages. Discharge/charge capacity and CE of POM-650 as a function of cycles at (b) 100 mA/g and (c) 1 A/g. (d) Rate performance of POM-650 from 0.1 to 5 A/g. Galvanostatic curves of (e) POM-900 and (f) POM-CS in a voltage range of 0.1-3.0 V at 100 mA/g depending on different cycling stages. (g) Discharge/charge capacity and CE of POM-900 and POM-CS as a function of cycles at 100 mA/g. (h) Rate performance of POM-900 and POM-CS from 0.1 to 2 A/g.


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Figure 6. (a-c) SEM images of cycled POM-650 grains with different scale and view after long 160 cycles at 1 A/g. No serious grain cracking is observed as a consequence of compact coexistence of product nanodomains and binder wires. (d,e) TEM images of discharged POM-650 in different scales after 33 cycles at 100 mA/g, showing numerous nanoparticles (dark contrast) with a size smaller than 20 nm embedded in pyrolytic TDA-1-H+ matrix (light contrast). (f) HRTEM image and (g) SAED pattern of discharged POM-650, disclosing that nanoparticles are still compactly stacked with visible phase boundaries, and they correspond to Mo metal and Li2O nanodomains from the lattice strips and diffraction rings.



(a)

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(b)

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Figure 7. Electrochemical impedance spectra and their fitting curves of POM-650, POM-900 and POM-CS (a) before cycling and (b) after 20 cycles at 500 mA/g. The high-frequency and middle-frequency semicircles respectively represent SEI film (Rf) and charge transfer (Rct) resistances. A low-frequency inclined straight line is relevant to finite-length Warburg impedance Zflw. (c) Equivalent circuit illustration to fit the impedance plots of (a) and (b). Evolution of (d) Rf and (e) Rct values of POM-650, POM-900 and POM-CS.


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Highly Reversible Conversion Anodes Composed of Ultra-Large

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